Diffusion in GaAs and other III-V Semiconductors - Stephen J. Pearton

Diffusion in GaAs 

and other III-V Semiconductors 

10 Years of Research 

Editor: 

D.J. Fisher 

SCITEC PUBLICATIONS

Notes: 

Each item in this section of the volume begins with a graphical compilation of relevant diffusion data which 

have been reported during the past decade. The plotted data are also tabulated as indicated on the graph. In some 

cases, the tabulated data have been obtained by digitizing published graphs and the values may not correspond 

exactly with the author's unpublished raw data. 

Refers to table N 

3N Bulk Diffusion - Quantitative Data 

The migration of Ag from epitaxial layers and into (111) samples of Si, 

during annealing at temperatures of between 450 and 500C, was studied 

by means of secondary ion mass spectrometric depth profiling. It was 

found that the diffusivities lay between 8 x 10 -16 and 1.6 x 10 -15 cm 2 /s 

(table N). These values were lower than were expected on the basis of 

previous data. 

T.C.Nason, G.R.Yang, K.H.Park, T.M.Lu: Journal of Applied Physics, 

1991, 70[3], 1392-6 

[446-91/92-027] 

Indicates volume and page number in 

DDF where abstract first appeared

AlAs 

Ag 

AlAs/GaAs: Ag Diffusion 

Various elements were diffused into a superlattice structure at temperatures of between 

700 and 1000C. Their disordering effect upon the superlattice was assessed by using a 

small-angle polishing method. The diffusion of Ag had no disordering effect upon the 

superlattice. The results were explained in terms of the interstitial-substitutional 

mechanism, and of the solubility of the given dopant in GaAs. 

H.P.Ho, I.Harrison, N.Baba-Ali, B.Tuck, M.Henini: Journal of Electronic Materials, 

1991, 20[9], 649-52 

[446-84/85-002] 

Al 

31 AlAs/GaAs: Al Diffusion 

The intermixing of superlattices was investigated as a function of the Si concentration 

following annealing at temperatures ranging from 500 to 900C. The superlattice samples 

were prepared by means of molecular beam epitaxy, and the near-surface layers were 

doped with Si to concentrations of between 2 x 10 17 and 5 x 10 18 /cm 3 . The Si and Al 

depth profiles were measured by means of secondary ion mass spectrometry. The 

diffusion length and activation energy of Al, as a function of Si dopant concentration, 

were deduced from the secondary ion mass spectrometry data. Within the above 

temperature range a single activation energy, for Al diffusion, of about 4eV was observed 

(table 1). The Al diffusion coefficient increased rapidly with Si concentration. 

P.Mei, H.W.Yoon, T.Venkatesan, S.A.Schwarz, J.P.Harbison: Applied Physics Letters, 

1987, 50[25], 1823-5 

[446-157/159-227] 

227

Al AlAs Au 

Table 1 

Diffusivity of Al in AlAs/GaAs 

Si (/cm 3 ) Temperature (C) D (cm 2 /s) 

5 x 10 17 900 6.1 x 10 -17 

5 x 10 17 850 3.0 x 10 -17 

1 x 10 18 795 5.0 x 10 -17 

2 x 10 18 750 6.6 x 10 -17 

2 x 10 18 750 5.4 x 10 -17 

5 x 10 18 700 4.2 x 10 -17 

5 x 10 17 795 2.3 x 10 -18 

2 x 10 18 695 1.1 x 10 -17 

2 x 10 18 700 5.6 x 10 -18 

5 x 10 18 650 5.6 x 10 -18 

1 x 10 18 745 2.6 x 10 -18 

2 x 10 18 655 1.0 x 10 -18 

2 x 10 18 655 7.3 x 10 -19 

2 x 10 17 900 1.2 x 10 -17 

5 x 10 17 745 3.6 x 10 -19 

- 850 4.2 x 10 -20 

- 800 2.6 x 10 -20 

AlAs/GaAs: Al Diffusion 

Enhanced layer interdiffusion. in Te-doped (2 x 10 17 to 3 x 10 18 /cm 3 ) organometallic 

chemical vapor deposited superlattices was studied by using secondary ion mass 

spectrometry. It was found that, at temperatures ranging from 800 to 1000C, the Al 

diffusion coefficient had an activation energy of 3eV and was approximately proportional 

to the Te content. In the case of Si-induced mixing, the activation energy for Al diffusion 

was 4.1eV and exhibited a power-law dependence upon the Si content. 

P.Mei, S.A.Schwarz, T.Venkatesan, C.L.Schwartz, E.Colas: Journal of Applied Physics, 

1989, 65[5], 2165-7 

[446-72/73-002] 

Au 

AlAs/GaAs: Au Diffusion 



small-angle polishing method. The diffusion of Au had no disordering effect upon the 

228

Au AlAs Si 




1991, 20[9], 649-52 

[446-84/85-002] 

Cu 

AlAs/GaAs: Cu Diffusion 



small-angle polishing method. The diffusion of Cu had no disordering effect upon the 




1991, 20[9], 649-52 

[446-84/85-002] 

Mn 

AlAs/GaAs: Mn Diffusion 



small-angle polishing method. It was found that Mn induced disordering of the 

superlattice. However, the disordering effect which arose from Mn diffusion could be 

entirely inhibited if the fraction of As in the diffusion source were considerably higher 

than that of Mn. This inhibition effect was related to the formation of MnAs or MnAs 2 . 

This left very little Mn, in the vapor phase, which was available for diffusion. The results 

were explained in terms of the interstitial-substitutional mechanism, and of the solubility 

of the given dopant in GaAs. 


1991, 20[9], 649-52 

[446-84/85-002] 

Si 

AlAs/AlGaAsP/GaAs: Si Diffusion 

The diffusion of Si III -Si V neutral pairs versus the diffusion of Si III -V III complexes in III-V 

crystals was considered with regard to experimental data which revealed the effect of Si 

diffusion upon the self-diffusion of column-III and column-V lattice atoms. Secondary 

ion mass spectroscopy was used to compare the enhanced diffusion of column-III or 

column-V atoms in various Si-diffused heterostructures which were closely lattice- 

229

Si AlAs Surface 

matched to GaAs. An enhancement of lattice atom self-diffusion, due to impurity 

diffusion, was found to occur predominantly on the column-III lattice. The data supported 

the Si III -V III diffusion model and indicated that the main native defects which 

accompanied Si diffusion were column-III vacancies. These diffused directly on the 

column-III sub-lattice. 

D.G.Deppe, W.E.Plano, J.E.Baker, N.Holonyak, M.J.Ludowise, C.P.Kuo, R.M.Fletcher, 

T.D.Osentowski, M.G.Craford: Applied Physics Letters, 1988, 53[22], 2211-3 

[446-64/65-157] 

Zn 

AlAs/GaAs: Zn Diffusion 

The diffusion of Zn into superlattices was studied by using transmission electron 

microscopy and secondary ion mass spectroscopy. It was found that micro-defects existed 

near to the Zn diffusion front. These defects were interstitial dislocation loops. It was 

suggested that the diffusion of Zn into the present materials was similar to Zn diffusion 

into GaAs. This was considered to be evidence for an interstitial mechanism for the 

enhancement of interdiffusion. 

I.Harrison, H.P.Ho, B.Tuck, M.Henini, O.H.Hughes: Semiconductor Science and 

Technology, 1989, 4[10], 841-6 

[446-72/73-002] 

AlAs/GaAs: Zn Diffusion 

The effect of an As pressure upon the disordering effect of Zn diffusion into superlattices 

was studied. It was found that the degree of disordering increased when no excess As was 

added to the ampoule. It had previously been found that dislocation loops formed near to 

the Zn diffusion front. The same effect was observed here, except when Zn diffusion was 

carried out in the absence of excess As. The Zn penetration was found to be greatest 

when no excess As was added to the diffusion ampoule. 

I.Harrison, H.P.Ho, B.Tuck, M.Henini, O.H.Hughes: Semiconductor Science and 

Technology, 1990, 5[6], 561-5 

[446-74-001] 

Surface Diffusion 

Al 

AlAs: Al Surface Diffusion 

During the molecular beam epitaxial growth of AlAs on the vicinal (100) surface of 

GaAs, reflection high-energy electron diffraction was used to measure the transition 

temperature between 2-dimensional nucleation and pure step propagation which occurred 

when sub-monolayer amounts of Sn were present on the surface. In the case of samples 

230

Surface AlAs Surface 

which were misoriented by 0.5º with respect to the [011] or [01¯1] direction, the transition 

temperature decreased by approximately 100C after the deposition of 0.6 of a monolayer 

of Sn. The presence of Sn increased the surface mobility of Al adatoms on (100) AlAs 

surfaces; as indicated by the annealing behavior of the AlAs surface at 600C. 

G.S.Petrich, A.M.Dabiran, P.I.Cohen: Applied Physics Letters, 1992, 61[2], 162-4 

[446-93/94-001] 

AlAs/GaAs: Al Surface Diffusion 

A study was made of reflection high-energy electron diffraction specular-beam intensity 

oscillations on vicinal (001)AlAs which had been grown onto GaAs(001) substrates that 

were misoriented by 2 or 3° towards [110], [010], or [¯110]. It was found that the 

temperature dependence of the oscillation behavior on vicinal surfaces was similar to that 

on GaAs(001) and InAs(001). Contrary to the case of GaAs(001), however, the surface 

reconstruction could not be kept constant during the growth-mode transition and it was 

therefore difficult to analyze AlAs(001) data in as much detail as that for GaAs(001). 

Nevertheless, from the similarity between them it was estimated that the effective surface 

migration barrier for Al adatoms on AlAs(001) was about 1.74eV. 

T.Shitara, J.H.Neave, B.A.Joyce: Applied Physics Letters, 1993, 62[14], 1658-60 

[446-106/107-007] 

Ga 

AlAs: Ga Surface Diffusion 

During the molecular beam epitaxial growth of AlAs on the vicinal (100) surface of 






of Sn. This indicated that the Ga mobility had increased. 


[446-93/94-001] 

-miscellaneous 

AlAs: Surface Diffusion 

An investigation was made of surface kinetics, during metalorganic vapor-phase epitaxial 

growth, by means of high-vacuum scanning tunnelling microscopic observations of 2- 

dimensional nuclei and denuded zones. Monte Carlo simulations were carried out which 

were based upon the solid-on-solid model. Two-dimensional nucleus densities were used to 

deduce that the surface diffusion coefficient of AlAs was equal to 1.5 x 10 -7 cm 2 /s at 530C. 

The activation energy for migration was estimated to be 0.80eV. The 2-dimensional nucleus 

size in the [110] direction was about twice that in the [¯110] direction. 

231

Surface AlAs General 

This anisotropy was attributed mainly to a difference in the lateral sticking probabilities 

between steps along [¯110] and those along [110]. The ratio of the sticking probabilities 

was estimated to be greater than 3:1. The denuded zone widths on the upper terraces were 

some 2 times wider than those on the lower terraces. This suggested that the sticking 

probability at descending steps was 10 to 300 times larger than the probability at 

ascending steps. 

M.Kasu, N.Kobayashi: Journal of Crystal Growth, 1997, 170, 246-50 

[446-141/142-093] 

AlAs: Surface Diffusion 

The mechanisms of molecular beam epitaxy were investigated by growing and analyzing 

the shapes of facet structures which consisted of an (001) top surface and two (111)B side 

surfaces. The diffusion of Al was found to be almost negligible; regardless of the As flux. 

By analyzing the shape of the facet, the diffusion length of Al on a (001) surface was 

estimated to be about 0.02µ at 580C. 

S.Koshiba, Y.Nakamura, M.Tsuchiya, H.Noge, H.Kano, Y.Nagamune, T.Noda, 

H.Sakaki: Journal of Applied Physics, 1994, 76[7], 4138-44 

[446-117/118-159] 

AlAs/GaAs: Surface Diffusion 

After depositing 1/6 of a monolayer of AlAs onto a very flat GaAs (001) surface by 

means of metalorganic vapor-phase epitaxy, a study was made of AlAs 2-dimensional 

nuclei by means of high-vacuum scanning tunnelling microscopy. The AlAs 2- 

dimensional nuclei elongated in the [110] direction, like GaAs. The density of AlAs 2- 

dimensional nuclei in the saturation region was 5 x 10 10 /cm 2 at 580C. The saturated AlAs 

2-dimensional nucleus density decreased with increasing temperature. On the basis of the 

saturated AlAs 2-dimensional nucleus densities, the surface diffusion coefficient of AlAs 

on GaAs was estimated to be 1.5 x 10 -7 cm 2 /s at 530C. This was an order of magnitude 

lower than that of GaAs on GaAs. 

M.Kasu, N.Kobayashi: Applied Physics Letters, 1995, 67[19], 2842-4 

[446-125/126-111] 

General 

AlAs/GaAs: Self-Diffusion 

Cation self-diffusion in superlattices was examined in terms of the activation enthalpy. It 

was suggested that cation diffusion should be mediated by As-antisite point defects, via 

the use of (As)antisite-rich materials or As-rich diffusion sources. It was also proposed 

that (As)antisite-mediated cation diffusion should exhibit a characteristic activation 

enthalpy of about 2.5eV under extrinsic conditions. Published data on interdiffusion in 

superlattices revealed a Fermi level dependence of the activation enthalpy. On this basis, 

232

General AlAs Interdiffusion 

it was concluded that the As-antisite defect was responsible for p-type impurity-enhanced 

cation self-diffusion. 

H.Iguchi: Japanese Journal of Applied Physics, 1989, 28[12], L2115-8 

[446-74-001] 

Interdiffusion 

AlAs/GaAs: Interdiffusion 

It was noted that undoped superlattices which had been grown at low temperatures 

underwent marked interface intermixing upon increasing the annealing temperature up to 

900C. Quantum confinement shifts which were caused by the intermixing of lowtemperature 

re-grown and normal superlattices were studied by using electro-modulation 

spectroscopy. The effective activation energy for intermixing in the low-temperature 

superlattices during isochronal post-growth annealing (30s) was found to be 0.32eV. This 

value was anomalously lower than that for superlattices that were grown at normal 

temperatures. Roughening of the interfaces, due to As precipitates, was associated with 

the intermixing. 

I.Lahiri, D.D.Nolte, J.C.P.Chang, J.M.Woodall, M.R.Melloch: Applied Physics Letters, 

1995, 67[9], 1244-6 

[446-125/126-111] 

AlAs/GaAs: Interdiffusion 

The dopant-induced intermixing of Al and Ga in as-grown short-period superlattices was 

studied by means of atomic resolution cross-sectional scanning tunnelling microscopy. In 

the case of Si-doped n-type AlAs/GaAs short-period superlattices, the intermixing 

increased with increasing Si concentration (0 to 5 x 10 18 /cm 3 ). In the case of Be-doped p- 

type AlAs/GaAs short-period superlattices, no intermixing of Al and Ga was observed; 

regardless of the Be concentration (0 to 5 x 10 18 /cm 3 ). 

J.F.Zheng, M.Salmeron, E.R.Weber: Solid State Communications, 1995, 93[5], 419-23 

[446-119/120-187] 

233

(Al,Ga)As 

Al 

AlGaAs/AlAs: Al Diffusion 

Undoped superlattice structures were grown, with or without the presence of 120 Sn 

implants, by using molecular beam epitaxy. They were then annealed under Si 3 N 4 , SiO 2 

or encapsulant films. It was found that an enhancement of the Al-Ga interdiffusion 

coefficient occurred under the Si 3 N 4 and SiO 2 films, due to the in-diffusion of Si from the 

films. The enhancement was greater during diffusion of the Sn implant. Intermixing 

enhancement was attributed to the operation of the Fermi effect. Beneath the WN x film, 

interdiffusion was suppressed even in the presence of the Sn dopant. 

E.L.Allen, C.J.Pass, M.D.Deal, J.D.Plummer, V.F.K.Chia: Applied Physics Letters, 1991, 

59[25], 3252-4 

[446-84/85-006] 

AlGaAs/GaAs: Al Diffusion 

Data were presented which showed that the Al-Ga interdiffusion coefficient for an 

Al x Ga 1-x As-GaAs quantum-well heterostructure or a superlattice was highly dependent 

upon the crystal encapsulation conditions. The activation energy for Al-Ga interdiffusion, 

and thus layer-disordering, was smaller (about 3.5eV) for dielectric encapsulated samples 

than after capless annealing (about 4.7eV). The interdiffusion coefficient for Si 3 N 4 - 

capped samples was almost an order of magnitude smaller than for the cases of capless or 

SiO 2 -capped samples at temperatures of between 800 and 875C. As well as the type of 

encapsulant, the encapsulation geometry (stripes or capped stripes) was important 

because of strain effects. These were a major source of anisotropic Al-Ga interdiffusion. 

L.J.Guido, N.Holonyak, K.C.Hsieh, R.W.Kaliski, W.E.Plano, R.D.Burnham, 

R.L.Thornton, J.E.Epler, T.L.Paoli: Journal of Applied Physics, 1987, 61[4], 1372-9 

[446-60-002] 

234

Al (Al,Ga)As Be 

AlGaAs/GaAs: Al Diffusion 

Photoluminescence spectroscopy was used to determine the temperature and 

compositional dependence of the interdiffusion of Al and Ga in (Al,Ga)As/GaAs 

superlattices. The position of the band-to-band luminescence in the superlattices was 

measured before and after thermal annealing. The diffusion equation was solved for a 

fixed value of the diffusion coefficient in order to establish the potential profile of the 

superlattice structure after annealing. A solution of the Schrödinger equation, where the 

electron or hole wave function was expanded as a Fourier series, was used to determine 

the position of the superlattice band edges before and after annealing and thus deduce the 

expected luminescence peak positions. The value of the coefficient which yielded a 

calculated shift which was in agreement with the measured shift in the luminescence was 

taken to be the actual value of the interdiffusion coefficient. For structures consisting of 

GaAs wells and Al x Ga 1-x As barriers, where x was 1 or 0.3, the interdiffusion process was 

characterized by an activation energy of 6.0eV and a value of 4 x 10 -19 cm 2 /s at 850C. 

When x was equal to 0.7, the interdiffusion was characterized by an activation energy of 

4.0eV and a value of 7 x 10 -18 cm 2 /s at 850C. 

J.C.Lee, T.E.Schlesinger, T.F.Kuech: Journal of Vacuum Science and Technology B, 

1987, 5[4], 1187-90 

[446-55/56-002] 

Be 

AlGaAs: Be Diffusion 

A close relationship between Be surface segregation and diffusion, in molecular beam 

epitaxial AlGaAs layers which were heavily doped with Be, was analyzed within the 

framework of a thermodynamic approach to segregation effects.. The effect of growth 

parameters (excess As pressure, substrate temperature, growth rate) and dopant level 

upon the likelihood of Be segregation layer formation was considered. 

S.V.Ivanov, P.S.Kopev, N.N.Ledentsov: Journal of Crystal Growth, 1991, 108[3-4], 661- 

9 

[446-81/82-002] 

GaAlAs: Be Diffusion 

A study was made of the contributions of segregation, diffusion and aggregation to the 

broadening of d-doped planes of Be in Ga 0.67 Al 0.33 As. It was found that sharp spikes of 

Be could be obtained for sheet densities which were below 10 13 /cm 2 and for growth 

temperatures of 500C or less. At higher temperatures or densities, segregation or 

concentration-dependent rapid diffusion could occur; thus causing significant spreading 

even during growth. The co-deposition of Si and Be markedly reduced this broadening. 

235

Be (Al,Ga)As Be 

J.J.Harris, J.B.Clegg, R.B.Beall, J.Castagné, K.Woodbridge, C.Roberts: Journal of 

Crystal Growth, 1991, 111[1-4], 239-45 

[446-91/92-001] 


A molecular beam epitaxial technique was developed in order to suppress Be diffusion by 

incorporating In into a GaAlAs epilayer. The diffusion coefficients of Be-doped 

In x (Ga 0.9 Al 0.1 ) 1-x As, grown at 600C, decreased from 10 -14 to 2 x 10 -15 cm 2 /s when the InAs 

mole fraction, x, was increased from 0 to 0.07. This indicated that compressive stresses in 

the epilayer, caused by the incorporation of In, played an important role in suppressing 

Be diffusion. 

T.Tomioka, T.Fujii, H.Ishikawa, S.Sasa, A.Endoh, Y.Bamba, K.Ishii, Y.Kataoka: 

Japanese Journal of Applied Physics, 1990, 29[5], L716-9 

[446-76/77-001] 


The suppression of Be diffusion in molecular beam epitaxial Ga 0.9 Al 0.1 As was reported 

here for the first time. It was achieved by incorporating In into the epilayer. The minimum 

Be diffusion coefficient in In-doped layers with a carrier concentration of 7 x 10 19 /cm 3 

and an InAs mole fraction of 0.07, which had been grown at 600C, was equal to about 2 x 

10 -15 cm 2 /s. This value was 5 times smaller than that which was observed in the absence 

of In. The photoluminescence intensity of the layers decreased markedly in In x (Al,Ga) 1-x 

when x was greater than 0.05. This behavior was attributed to a crystal degradation which 

resulted from the presence of misfit dislocations. 

T.Fujii, T.Tomioka, H.Ishikawa, S.Sasa, A.Endoh, Y.Bamba, K.Ishii, Y.Kataoka: Journal 

of Vacuum Science and Technology B, 1990, 8[2], 154-6 

[446-74-003] 


The migration of ion-implanted Be was studied as a function of Al concentration and 

annealing temperature and was compared with its diffusivity in GaAs. The behavior of Be 

in AlGaAs was similar to that in GaAs, and it even exhibited the anomalous characteristic 

of increased redistribution with decreasing temperature. The results could be described 

by: 

Ga 0.8 Al 0.2 As: D(cm 2 /s) = 1.8 x 10 -9 exp[-0.90(eV)/kT] 


The diffusivity of Be appeared to increase with Al content. This was suggested to be due 

to an increase in the bond strength of matrix atoms upon adding Al. This prevented the 

easy transfer of Be from interstitial to substitutional sites. An over-saturation of Be 

interstitials could also explain the persistence of anomalous diffusion in AlGaAs with 

respect to the annealing temperature. The results were explained in terms of a 

substitutional-interstitial diffusion mechanism, the relative amounts of interstitial and 

236

Be (Al,Ga)As Be 

substitutional Be, and the relative difficulty of moving from an interstitial to a 

substitutional site. 

C.C.Lee, M.D.Deal, J.C.Bravman: Applied Physics Letters, 1995, 66[3], 355-7 

[446-123/124-160] 

AlGaAs/AlAs: Be Diffusion 

Dopant diffusion was investigated via depth-profiling using secondary ion mass 

spectrometric analysis of heterostructures which contained Be-doped short-period Al x Ga 1- 

xAs superlattices, where x ranged from 0.3 to 0.8, that had been grown by means of 

molecular beam epitaxy. Out-diffusion of Be into the undoped GaAs layers was observed 

only at a substrate temperature of 660C, when the Be concentration was 2 x 10 18 /cm 3 . At 

a dopant concentration of 2 x 10 19 /cm 3 , a marked increase in diffusion occurred at all 

growth temperatures. The solubility limits of Be were 10 19 /cm 3 at x = 0.6, and 2 x 

10 18 /cm 3 at x = 0.8. Secondary ion mass spectrometry profiles revealed that the amount of 

diffused Be in the active region was twice as high in samples with a thin (450 to 600nm) 

p-type cladding layer. 

A.Gaymann, M.Maier, W.Bronner, N.Gruen, K.Koehler: Materials Science and 

Engineering B, 1997, 44[1-3], 12-5 

[446-157/159-237] 

AlGaAs/GaAs: Be Diffusion 

It was pointed out that the characteristic features of Be diffusion in GaAs substrates and 

GaAs/AlGaAs superlattices could be explained in terms of a kick-out mechanism in 

which the doubly positively charged Ga self-interstitial governed Ga self-diffusion. Such 

characteristics included much lower diffusivities of Be under out-diffusion conditions 

than under in-diffusion conditions. It was found that the Longini mechanism was able to 

explain most of the features. 

S.Yu, T.Y.Tan, U.Gösele: Journal of Applied Physics, 1991, 69[6], 3547-65 

[446-86/87-009] 


The effect of substrate orientation upon Be transport during GaAs molecular beam 

epitaxial growth was evaluated by means of secondary ion mass spectrometry and 

measurements of the current-voltage characteristics of AlGaAs/GaAs heterojunction 

bipolar transistors. The Be doping level was between 2 x 10 19 and 9 x 10 19 /cm 3 . The Be 

transport which was observed for the conventional (100) orientation increased rapidly 

upon increasing the growth temperature from 530 to 630C. However, with a substrate 

misorientation away from (100) and towards (111)A, Be transport decreased at 630C and 

reached a minimum value for the (311)A orientation. The maximum current gain, of 

AlGaAs/GaAs heterojunction bipolar transistors which had been grown at 560C, was 

equal to 264 for the (411)A orientation and 3 for the (100) orientation. It was concluded 

that this confirmed the applicability of substrate orientations other than the conventional 

(100) one for obtaining a sharp Be profile. 

237

Be (Al,Ga)As D 

K.Mochizuki, S.Goto, T.Mishima, C.Kusano: Japanese Journal of Applied Physics, 1992, 

31[1-11], 3495-9 

[446-99/100-059] 


A secondary ion mass spectrometry investigation was made of Be diffusion, during the 

molecular beam epitaxial growth of graded-index separate confinement heterostructure 

laser structures. In the case of growth at 700C, it was found that Be from the p-type 

AlGaAs cladding layer diffused into the quantum well and beyond. As a result, the p-n 

junction was displaced from the heterojunction. The extent of Be diffusion was found to 

depend upon the dopants in the graded-index regions which adjoined the GaAs active 

layer. When the graded-index segments were left undoped, Be diffused through the entire 

p-side graded-index region, the quantum well active region, and a significant portion of 

the n-side graded-index region. However, when the graded-index regions were doped 

with Be and Si on the p-side and n-side, respectively, the displacement of the p-n junction 

which was caused by Be diffusion was significantly reduced. Upon assuming that Be 

diffused from a constant surface source and into an n-type layer, as a singly charged 

interstitial donor, the present analysis predicted that increasing the doping of the n-type 

layer would retard Be diffusion; whereas increasing the doping of the p-type layer would 

enhance it. Upon including the electric field of the p-n junction in the model, peaks and 

inflections were predicted which resembled those that were observed in experimental 

secondary ion mass spectroscopy profiles. It was concluded that, because of Be-related O 

contamination and Be diffusion in the p-side graded-index region, the presence of Be 

should be avoided on the p side. However, Si additions to the n side were expected to be 

beneficial as they minimized Be diffusion and p-n junction displacement. 

V.Swaminathan, N.Chand, M.Geva, P.J.Anthony, A.S.Jordan: Journal of Applied 

Physics, 1992, 72[10], 4648-54 

[446-106/107-015] 

D 

GaAlAs: D Diffusion 

The diffusion of D in Si-doped Al x Ga 1-x As was studied for x-values of up to 0.30. It was 

found that, for x = 0, the diffusion profile could be closely fitted by using an erfc 

function. When the x-value was greater than 0.055, the profiles exhibited a plateau that 

was followed by a sharp decrease. It was suggested that, in Si-doped samples, the D 

behaved like a deep acceptor with a level, H -/0 , which was slightly resonant in the 

conduction band of GaAs. It appeared as a localized state, for x-values above 0.07, as the 

band-gap energy increased. In this region, the H - species became dominant and were 

trapped on the positively charged donors during diffusion. 

J.Chevallier, B.Machayekhi, C.M.Grattepain, R.Rahbi, B.Theys: Physical Review B, 

1992, 45[15], 8803-6 

[446-86/87-001] 

238

Ga (Al,Ga)As Ga 

Ga 

32 AlGaAs: Ga Diffusion 

The intermixing of AlGaAs-based interfaces was enhanced by capping wafers with a 

layer of SiO 2 . By assuming that this enhancement resulted from the introduction of 

additional Ga vacancies into the sample, it was possible to estimate the temperaturedependent 

equilibrium Ga vacancy diffusivity. Experiments were performed in which 

SiO 2 -capped quantum-well samples were annealed at temperatures ranging from 800 to 

1025C. The calculated photoluminescence shifts were compared with the measured 

spectra and a relationship, for the Ga vacancy diffusivity, of the form: 

D (cm 2 /s) = 0.962 exp[-2.72(eV)/kT] 

(table 2) was obtained. By using this relationship, the equilibrium Ga vacancy 

concentration could be estimated via Monte Carlo simulation. The resultant expression 

was: C (/cm 3 ) = 1.25 x 10 31 exp[-3.28(eV)/kT]. 

K.B.Kahen, D.L.Peterson, G.Rajeswaran, D.J.Lawrence: Applied Physics Letters, 1989, 

55[7], 651-3 

[446-70/71-103] 

Table 2 

Diffusivity of Ga Vacancies in AlGaAs 

Temperature (C) D (cm 2 /s) 

1030 2.8 x 10 -11 

1005 1.7 x 10 -11 

950 6.3 x 10 -12 

900 2.3 x 10 -12 

AlGaAs/AlAs: Ga Diffusion 





films. The enhancement was greater during diffusion of the Sn implant. In both cases, 

intermixing enhancement was attributed to the operation of the Fermi effect. Beneath the 

WN x film, interdiffusion was suppressed even in the presence of the Sn dopant. 


59[25], 3252-4 

[446-84/85-006] 

239

Ga (Al,Ga)As Ga 

33 AlGaAs/GaAs: Ga Diffusion 

Photoluminescence spectroscopy was used to determine the temperature and 

compositional dependence of the interdiffusion of Al and Ga in (Al,Ga)As/GaAs 

superlattices. The position of the band-to-band luminescence in the superlattices was 

measured before and after thermal annealing. The diffusion equation was solved for a 

fixed value of the diffusion coefficient in order to establish the potential profile of the 

superlattice structure after annealing. A solution of the Schrödinger equation, where the 

electron or hole wave function was expanded as a Fourier series, was used to determine 

the position of the superlattice band edges before and after annealing and thus deduce the 

expected luminescence peak positions. The value of the coefficient which yielded a 

calculated shift which was in agreement with the measured shift in the luminescence was 

taken to be the actual value of the interdiffusion coefficient. For structures consisting of 

GaAs wells and Al x Ga 1-x As barriers, where x was 1 or 0.3, the interdiffusion process was 

characterized by an activation energy of 6.0eV and a value of 4 x 10 -19 cm 2 /s at 850C. 

When x was equal to 0.7, the interdiffusion was characterized by an activation energy of 

4.0eV and a value of 7 x 10 -18 cm 2 /s at 850C (table 3). 

J.C.Lee, T.E.Schlesinger, T.F.Kuech: Journal of Vacuum Science and Technology B, 

1987, 5[4], 1187-90 

[446-55/56-002] 

Table 3 

Interdiffusivity (Al-Ga) in Al x Ga 1-x As/GaAs 

x Temperature (C) D (cm 2 /s) 

0.3 905 1.1 x 10 -17 

0.3 875 3.4 x 10 -18 

0.3 865 8.2 x 10 -19 

0.3 850 4.0 x 10 -19 

0.3 820 1.1 x 10 -19 

0.7 820 2.4 x 10 -18 

0.7 790 1.1 x 10 -18 

0.7 775 4.6 x 10 -19 

0.7 750 1.3 x 10 -19 

0.7 750 1.1 x 10 -19 

AlGaAs/GaAs: Ga Diffusion 

It was pointed out that the characteristic features of Be and Zn diffusion in GaAs 

substrates and GaAs/AlGaAs superlattices could be explained in terms of a kick-out 

mechanism in which the doubly positively charged Ga self-interstitial governed Ga selfdiffusion. 

It was found that the Longini mechanism was able to explain most of these 

features. However, the predictions of the Longini mechanism with regard to Ga self- 

240

Ga (Al,Ga)As In 

diffusion disagreed with experimental observations of the effect of superlattice 

disordering. 


[446-86/87-009] 

H 

AlGaAs: H Diffusion 

Diffusion experiments were performed on samples of Si-doped Al x Ga 1-x As epitaxial 

layers, with x-values which ranged from 0 to 0.30, as a function of the Si doping level 

and the diffusion temperature. For each composition, calculated H diffusion profiles 

which had been deduced by using Mathiot's model were fitted to the experimental 

profiles. It was assumed that H behaved as a deep acceptor, and that H o and H - were the 

diffusing species. The trapping of H - by Si + donors, and their acceleration by an electric 

field, were incorporated into the model. As well as the diffusion coefficient of H, and the 

dissociation constant of the SiH complexes, the model provided for a compositional 

dependence of the H acceptor level in AlGaAs alloys. It was concluded that the H 

acceptor level was localized in the band-gap of the present AlGaAs alloys, and deepened 

below the Γ conduction band as x increased. 

B.Machayekhi, R.Rahbi, B.Theys, M.Miloche, J.Chevallier: Materials Science Forum, 

1994, 143-147, 951-6 

[446-113/114-001] 

GaAlAs: H Diffusion 

Layers of material, which was doped with various group-VI donors (S, Se, Te), were 

exposed to H plasma. By using secondary ion mass spectroscopy it was shown that, as in 

the case of Si-doped materials, the diffusivity of H depended strongly upon the AlAs 

content. Electronic measurements indicated that, after H diffusion, the electron 

concentration systematically decreased while their mobility increased; thus demonstrating 

the passivation of the group-VI donors by H. 

B.Theys, B.Machayekhi, J.Chevallier, K.Somogyi, K.Zahraman, P.Gibart, M.Miloche: 

Journal of Applied Physics, 1995, 77[7], 3186-93 

[446-121/122-052] 

In 

AlGaAs/GaAs: In Diffusion 

Data were presented on the disordering of an AlGaAs/GaAs laser structure using In solid 

sources. By using the independent or combined diffusion of Si and In from thin-film 

sources, it was deduced that In had a higher diffusion coefficient than Si and led to a 

similar degree of impurity-induced disordering. The degree of index guiding was tested 

by making excess-loss measurements in single-mode raised-cosine s-bends. It was found 

241

In (Al,Ga)As Mn 

that structures which were patterned by SiO 2 /In disordering suffered excess losses which 

were similar to those in structures that were patterned with SiO 2 . An 0.26mm transition 

length for 3dB loss was measured for 1000nm-wide guides with 0.1mm guide offsets. 

This corresponded to a lateral index of refraction difference of between 0.8 and 1.0%. 

There was no evidence for an increased linear loss due to the presence of a dilute InGaAs 

alloy at a measurement wavelength of 870nm. 

T.K.Tang, J.J.Alwan, C.M.Herzinger, T.M.Cockerill, A.Crook, T.A.DeTemple, 

J.J.Coleman, J.E.Baker: Applied Physics Letters, 1991, 59[22], 2880-2 

[446-91/92-005] 

Mg 

AlGaAs: Mg Diffusion 

Layer samples were diffused, at 785C, with Mg from an As-saturated Ga solution that 

contained 0.1wt%Mg. Secondary ion mass spectrometry and differential Hall effect 

measurements revealed that the depth profile consisted of a high-Mg concentration region 

close to the surface, and a lower-concentration plateau within the sample. The diffusion 

of Mg into GaAs, Al 0.5 Ga 0.5 As and Al 0.7 Ga 0.3 As for 0.33h resulted in diffusion fronts at 

about 0.002, 0.004 and 0.006mm from the surface, respectively. The depth, for a fixed 

hole concentration, was proportional to the square root of the diffusion time in both GaAs 

and Ga 0.65 Al 0.35 As. 

S.Mukai, Y.Kaneko, T.Nukui, M.Mori, M.Watanabe, H.Itoh, H.Yajima: Japanese Journal 

of Applied Physics, 1989, 28[1], L1-3 

[446-64/65-158] 

AlGaAs: Mg Diffusion 

An investigation was made of dual Mg/F or Mg/Ar implantation. It was found that the 

dual implantation suppressed Mg diffusion, but degraded the electrical properties. This 

was more apparent in material with lower Al contents. The Ar dual implantation 

suppressed Mg diffusion and improved the electrical properties of material with a high Al 

content. It was suggested that Mg-F bonds formed as a result of F dual implantation and 

that successive annealing suppressed Mg diffusion and disturbed Mg activation. The 

extensive radiation damage which was caused by Ar dual implantation caused Mg atoms 

to occupy lattice sites in AlGaAs with a high Al content. 

N.Hara, H.Suehiro, S.Kuroda: Materials Science Forum, 1995, 196-201, 1943-8 

[446-127/128-107] 

Mn 

AlGaAs/GaAs: Mn Diffusion 

The diffusion of Mn was carried out in sealed quartz ampoules, using 4 types of Mn 

source. These were: solid crystalline Mn grains, Mn 3 As, MnAs, and Mn thin films on 

242

Mn (Al,Ga)As Si 

GaAs substrates. It was found that only MnAs led to the formation of a smooth GaAs 

surface and a uniform dopant distribution. In the case of the other sources, interactions 

between the source materials and the substrate gave rise to poor surface morphologies and 

inhomogeneous distributions. In the case of diffusion at 800C, surface p-type carrier 

concentrations of the order of 10 20 /cm 3 were obtained. The diffusion profiles which were 

determined by using capacitance-voltage techniques resembled those which were 

obtained for Zn diffusion. It was suggested that a substitutional-interstitial mechanism 

was the predominant one. It was also noted that layer disordering could be produced in 

AlGaAs-GaAs superlattices by Mn impurities. 

C.H.Wu, K.C.Hsieh, G.E.Höfler, N.El-Zein, N.Holonyak: Applied Physics Letters, 1991, 

59[10], 1224-6 

[446-84/85-007] 

O 

AlGaAs: O Diffusion 

A layer of SiO 2 , deposited by sputtering, was used as a diffusion source for O impurities, 

as well as a source of Ga vacancies which enhanced impurity diffusion and permitted 

reductions to be made in the required annealing temperatures and times. A self-aligned 

native oxide of an AlGaAs cladding layer was used to form a Zn diffusion mask and 

dielectric layer. 

R.S.Burton, T.E.Schlesinger, D.J.Holmgren, S.C.Smith, R.D.Burnham: Journal of 

Applied Physics, 1993, 73[4], 2015-8 

[446-106/107-008] 

Si 

AlGaAs: Si Diffusion 

Data were presented which showed that the native oxide that could form on Al x Ga 1-x As 

confining layers (where x was greater than 0.7) on Al y Ga 1-y As/Al z Ga 1-z As superlattices or 

quantum-well heterostructures (where y was greater than z) served as an effective barrier 

to Si impurity diffusion. It thus impeded impurity-induced layer disordering. High-quality 

native oxide was produced by the conversion of high x-value Al x Ga 1-x As confining layers 

(which could be grown on a variety of heterostructures) via H 2 O vapor oxidation (at 

temperatures of more than 400C) in N 2 carrier gas. 

J.M.Dallesasse, N.Holonyak, N.El-Zein, T.A.Richard, F.A.Kish, A.R.Sugg, 

R.D.Burnham, S.C.Smith: Applied Physics Letters, 1991, 58[9], 974-6 

[446-81/82-002] 


The diffusion and drift of Si were studied by means of capacitance-voltage measurements. 

These revealed that low substrate temperatures, during growth via molecular beam 

243

Si (Al,Ga)As Si 

epitaxy, were required in order to achieve Dirac d-like dopant profiles. It was further 

shown theoretically that the random Poisson distribution, which was usually assumed for 

dopant distributions in semiconductors, should be modified at high dopant concentrations. 

This was because of repulsive interactions between the impurities. 

E.F.Schubert, C.W.Tu, R.F.Kopf, J.M.Kuo, L.M.Lunardi: Applied Physics Letters, 1989, 

54[25], 2592-4 

[446-70/71-103] 


The self-aligned diffusion of Si was studied. It was found that the use of a Si film for the 

diffusion led to major problems of morphology degradation and dopant contamination 

during Si diffusion. A method which involved both a SiO 2 encapsulant and a sputtered Si 

film source (Si diffusion) or mask (Zn diffusion) was investigated. The optimum 

thicknesses of the Si and SiO 2 films were 18 and 55nm, respectively. 

W.X.Zou, S.Corzine, G.A.Vawter, J.L.Merz, L.A.Coldren, E.L.Hu: Journal of Applied 

Physics, 1988, 64[4], 1855-8 

[446-72/73-003] 


Data were presented which demonstrated that the surface encapsulant and As 4 overpressure 

strongly affected Si diffusion in Al x Ga 1-x As, and were important parameters in 

impurity-induced layer disordering. An increase in the As 4 over-pressure resulted in a 

decrease in the diffusion depth for Al x Ga 1-x As. In addition, the band-edge exciton was 

observed in absorption on an Al x Ga 1-x As-GaAs superlattice that was diffused with Si and 

was converted to bulk crystal Al y Ga 1-y As via impurity-induced layer disordering. The data 

indicated that the Si diffusion process and the properties of the diffused material were 

different for GaAs and Al x Ga 1-x As-GaAs superlattices which were converted into uniform 

Al y Ga 1-y As (where y was between 0 and 1) via impurity-induced layer disordering with 

amphoteric dopant Si. 

L.J.Guido, W.E.Plano, D.W.Nam, N.Holonyak, J.E.Baker, R.D.Burnham, P.Gavrilovic: 

Journal of Electronic Materials, 1988, 17[1], 53-6 

[446-60-004] 


The migration of Si into Al x Ga 1-x As from a sputtered Si film was described, where x 

ranged from 0 to 0.4. It was shown that both the diffusion rate and the surface Si 

concentration decreased with increasing Al mole fraction. The diffusion behavior of Si 

was explained in terms of the binding energy of the Al-As bond and of the disorder of the 

mixed crystal. 

E.Omura, X.S.Wu, G.A.Vawter, E.L.Hu, L.A.Coldren, J.L.Merz: Applied Physics 

Letters, 1987, 50[5], 265-6 

[446-55/56-001] 

244



A new method for self-aligned Si-Zn diffusion was described. In this method, closed-tube 

Si diffusion was carried out by using a sputtered SiN x film. Then, Zn diffusion which was 

self-aligned to the Si diffusion window was carried out by re-using the SiN x film as a 

mask. The key factor was that the SiN x film should have the correct refractive index 

profile. 

W.X.Zou, R.Boudreau, H.T.Han, T.Bowen, S.S.Shi, D.S.L.Mui, J.L.Merz: Journal of 


[446-121/122-045] 


A layer of SiO 2 , deposited by sputtering, was used as a diffusion source for Si impurities, 

as well as a source of Ga vacancies which enhanced impurity diffusion and permitted 

reductions to be made in the required annealing temperatures and times. A self-aligned 

native oxide of an AlGaAs cladding layer was used to form a Zn diffusion mask and 

dielectric layer. 

R.S.Burton, T.E.Schlesinger, D.J.Holmgren, S.C.Smith, R.D.Burnham: Journal of 


[446-106/107-008] 

GaAlAs: Si Diffusion 

The mechanism of Si diffusion in Ga 0.7 Al 0.3 As was studied by using photoluminescence 

and secondary ion mass spectrometry, and transmission electron microscopy across the 

corner of a wedge-shaped sample. The diffusion source was a grown-in highly Si-doped 

layer. It was deduced that Frenkel defects (column-III vacancies and interstitials) were 

generated within the highly Si-doped region. The column-III interstitials rapidly diffused 

towards the surface, where they reacted with the column-III vacancies which were 

generated at the surface during annealing in a gaseous As ambient. This caused a 

supersaturation, of column-III vacancies in the Si-doped region, which drove Si diffusion. 

Annealing in vacuum reduced the supersaturation of column-III vacancies, and thus 

decreased Si diffusion. A predominant Si-donor plus column-III vacancy complex 

emission band was found in spectra from the Si-diffused region. The results supported the 

concept of a vacancy-assisted mechanism for Si diffusion and impurity-induced 

disordering. 

L.Pavesi, N.H.Ky, J.D.Ganière, F.K.Reinhart, N.Baba-Ali, I.Harrison, B.Tuck, M.Henini: 


[446-86/87-002] 


The effect of the substrate temperature, during molecular beam epitaxial growth, upon the 

migration of Si atoms in d-doped or planar-doped Ga 0.75 Al 0.25 As was investigated by 

using secondary ion mass spectrometry. For substrate temperatures of 580 to 640C, the Si 

spread over about 35nm in d-doped Ga 0.75 Al 0.25 As. For substrate temperatures below 

245


580C, the measured width of the Si profile was limited by the resolution of the secondary 

ion mass spectrometer. Magneto-transport measurements were also performed in order to 

determine dopant spreading. The Si migration which was measured by means of 

secondary ion mass spectrometry was in qualitative agreement with the transport results. 

However, the secondary ion mass spectrometry data indicated larger Si areal densities. 

Two mechanisms, auto-compensation and electron localization by a DX center, were 

believed to be responsible for the latter observations. 

A.M.Lanzillotto, M.Santos, M.Shayegan: Applied Physics Letters, 1989, 55[14], 1445-7 

[446-72/73-002] 



broadening of d-doped planes of Si in Ga 0.67 Al 0.33 As. It was found that sharp spikes of Si 

could be obtained for sheet densities which were below 10 13 /cm 2 and for growth 

temperatures of 500C or less. At higher temperatures or densities, segregation or 


even during growth. The co-deposition of Si and Be markedly reduced this broadening. 



[446-91/92-001] 

AlGaAs/GaAs: Si Diffusion 

Under growth conditions which were optimized so as to give the best transport with 

normal-side doping, the migration of the Si dopant towards the inverted interface during 

growth was the main reason for a reduced inverted well mobility. This discovery 

permitted the preparation of modulation-doped inverted quantum wells of unprecedented 

quality. 

L.Pfeiffer, E.F.Schubert, K.W.West, C.W.Magee: Applied Physics Letters, 1991, 58[20], 

2258-60 

[446-84/85-007] 


The migration of Si during the metal-organic vapor-phase epitaxial growth of laser 

structures was studied by means of secondary ion mass spectroscopy. The migration 

process was found to depend mainly upon the Si concentration in the AlGaAs layer; for 

both silane and disilane doping gases. Above a critical concentration of about 3 x 

10 18 /cm 3 , Si migrated into the nominally undoped GaAs layer. This shift in the Si front 

became even more pronounced when the GaAs layer was grown at a lower rate than that 

of the AlGaAs layer. The Si depth profile had the same gradient as the Al depth profile; 

even in layers with a large shift of the Si front. Migration appeared to occur preferentially 

towards the growth front. It was concluded that the process was governed not only by 

diffusion, but also by surface kinetics. The effect of Si migration upon the threshold 

246


current density of broad-area lasers was significant only for a large shift of the Si front 

into the active GaAs layer. 

E.Veuhoff, E.Baumeister, R.Treichler: Journal of Crystal Growth, 1988, 93, 650-5 

[446-64/65-159] 



strongly affected Si diffusion in GaAs and Al x Ga 1-x As, and were important 

parameters in impurity-induced layer disordering. An increase in the As 4 over-pressure 

resulted in an increase in the diffusion depth in the case of GaAs, and a decrease in the 

diffusion depth for Al x Ga 1-x As. In addition, the band-edge exciton was observed in 

absorption on an Al x Ga 1-x As-GaAs superlattice that was diffused with Si and was 

converted to bulk crystal Al y Ga 1-y As via impurity-induced layer disordering. In contrast, 

the exciton was not observed during absorption on GaAs that was diffused with Si, in 

spite of the high degree of compensation. The data indicated that the Si diffusion process 

and the properties of the diffused material were different for GaAs and Al x Ga 1-x As-GaAs 

superlattices which were converted into uniform Al y Ga 1-y As (where y was between 0 and 

1) via impurity-induced layer disordering with amphoteric dopant Si. 



[446-60-004] 


Data were presented which showed that dislocations and Si diffusion promoted 

accelerated layer disordering of quantum well heterostructures which were grown on 

GaAs/Si substrates by metalorganic chemical vapor deposition. The accelerated impurityinduced 

layer disordering was more extreme at temperatures greater than 900C, and was 

virtually non-existent at temperatures below 775C. 

W.E.Plano, D.W.Nam, K.C.Hsieh, L.J.Guido, F.A.Kish, A.R.Sugg, N.Holonyak, 

R.J.Matyi, H.Shichijo: Applied Physics Letters, 1989, 55[19], 1993-5 

[446-72/73-005] 


Data were presented on the disordering of an AlGaAs/GaAs laser structure using In solid 

sources. By using the independent or combined diffusion of Si and In from thin-film 

sources, it was deduced that In had a higher diffusion coefficient than Si and led to a 

similar degree of impurity-induced disordering. The degree of index guiding was tested 

by making excess-loss measurements in single-mode raised-cosine s-bends. It was found 

that structures which were patterned by SiO 2 /In disordering suffered excess losses which 

were similar to those in structures that were patterned with SiO 2 . An 0.26mm transition 

length for 3dB loss was measured for 1000nm-wide guides with 0.1mm guide offsets. 

This corresponded to a lateral index of refraction difference of between 0.8 and 1.0%. 

247

Si (Al,Ga)As Sn 

There was no evidence for an increased linear loss due to the presence of a dilute InGaAs 

alloy at a measurement wavelength of 870nm. 

T.K.Tang, J.J.Alwan, C.M.Herzinger, T.M.Cockerill, A.Crook, T.A.DeTemple, 

J.J.Coleman, J.E.Baker: Applied Physics Letters, 1991, 59[22], 2880-2 

[446-91/92-005] 

GaAlAs/GaAs: Si Diffusion 

The Si migration and impurity-induced layer intermixing from a buried impurity source 

were studied by means of transmission electron microscopic and secondary ion mass 

spectroscopic studies of isolated Si-doped GaAs layers in an undoped Ga 0.6 Al 0.4 As/GaAs 

superlattice, and by photoluminescence measurements of Si-doped quantum wells with 

undoped Ga 0.6 Al 0.4 As barriers. In annealed samples, the Si profiles suggested the 

occurrence of a Si diffusion process which involved multiply ionized column-III 

vacancies. The width of the resultant Si profile, and the spatial extent and completeness 

of intermixing, depended strongly upon the initial Si concentration in the doped layer. 

K.J.Beernink, R.L.Thornton, G.B.Anderson, M.A.Emanuel: Applied Physics Letters, 

1995, 66[19], 2522-4 

[446-121/122-053] 

Sn 

GaAlAs: Sn Diffusion 


broadening of d-doped planes of Sn in Ga 0.67 Al 0.33 As. It was found that the Sn planes 

were severely broadened by all 3 processes. At higher temperatures or densities, 

segregation or concentration-dependent rapid diffusion could occur; thus causing 

significant spreading even during growth. 



[446-91/92-001] 

AlGaAs/AlAs: Sn Diffusion 





films. The enhancement was greater during diffusion of the Sn implant. In both cases, 

intermixing enhancement was attributed to the operation of the Fermi effect. Beneath the 

WN x film, interdiffusion was suppressed even in the presence of the Sn dopant. 


59[25], 3252-4 

[446-84/85-006] 

248

Te (Al,Ga)As Zn 

Te 

AlGaAs: Te Diffusion 

Monte Carlo simulation was used to model the enhanced disordering of AlGaAs-based 

interfaces in the presence of high concentrations of Te atoms. The model was based upon 

the experimental finding that the thermal interdiffusion process was similar to the selfdiffusion 

of Ga in GaAs. The model agreed well with experimental data for both Ga selfdiffusion 

and for intermixing. The intermixing was found to be caused by the enhanced 

solubility of Ga vacancy acceptors in the presence of donor Te atoms, and not by 

diffusion of the Te atoms. The activation energy for the process was found to be about 

2.7eV. 

K.B.Kahen: Applied Physics Letters, 1988, 53[21], 2071-3 

Zn 

[446-64/65-158] 

AlGaAs: Zn Diffusion 





profile. 



[446-121/122-045] 


Data were presented which showed that the native oxide that could form on Al x Ga 1-x As 

confining layers (where x was greater than 0.7) on Al y Ga 1-y As/Al z Ga 1-z As superlattices or 

quantum-well heterostructures (where y was greater than z) served as an effective barrier 

to Zn impurity diffusion. It thus impeded impurity-induced layer disordering. Highquality 

native oxide was produced by the conversion of high x-value Al x Ga 1-x As 

confining layers (which could be grown on a variety of heterostructures) via H 2 O vapor 

oxidation (at temperatures of more than 400C) in N 2 carrier gas. 

J.M.Dallesasse, N.Holonyak, N.El-Zein, T.A.Richard, F.A.Kish, A.R.Sugg, 

R.D.Burnham, S.C.Smith: Applied Physics Letters, 1991, 58[9], 974-6 

[446-81/82-002] 


The self-aligned diffusion of Zn was studied. It was found that the use of a Si film for the 

diffusion led to major problems of morphology degradation and dopant contamination 

249

Zn (Al,Ga)As Zn 

during Si diffusion. A method which involved both a SiO 2 encapsulant and a sputtered Si 

film source (Si diffusion) or mask (Zn diffusion) was investigated. The optimum 

thicknesses of the Si and SiO 2 films were 18 and 55nm, respectively. 

W.X.Zou, S.Corzine, G.A.Vawter, J.L.Merz, L.A.Coldren, E.L.Hu: Journal of Applied 

Physics, 1988, 64[4], 1855-8 

[446-72/73-003] 


The diffusion of Zn was studied by using liquid-phase epitaxy methods, and Si-doped n- 

type substrate material. The measurements were carried out at 850C, and dopant 

concentrations which ranged from 10 18 to 10 19 /cm 3 were introduced. It was found that the 

Zn concentration in the solid depended upon the square root of the atomic fraction of Zn 

in the liquid. The diffusivity was dominated by the interstitial-substitutional process, and 

exhibited a cubic dependence upon the Zn content. The Zn interstitial was mainly doublyionized 

Zn i 2+ . It was noted that Al played the role of a catalyst in the diffusion process. 

The Zn diffusion coefficient in Al 0.85 Ga 0.15 As was some 4 times greater than that in 

GaAs. 

C.Algora, G.L.Araujo, A.Marti: Journal of Applied Physics, 1990, 68[6], 2723-30 

[446-86/87-002] 


The use of thin Si films for the selective-area diffusion of Si and Zn was described. It was 

found that Si films behaved as ideal masks for Zn diffusion at temperatures below 750C. 

Ideal lateral Zn diffusion profiles were also observed when using these films; regardless 

of the stress at the interface. 

G.A.Vawter, E.Omura, X.S.Wu, J.L.Merz, L.Coldren, E.Hu: Journal of Applied Physics, 

1988, 63[11], 5541-7 

[446-72/73-003] 


The state of Zn diffusion at the hetero-interface of 660nm double-hetero light-emitting 

diodes was investigated by using the electron beam-induced current method. The p-n 

junction penetrated towards the n-cladding layer, as a result of Zn diffusion, when the 

carrier concentration of the p-active layer was greater than 10 18 /cm 3 . The distance 

between the hetero-interface and the p-n junction was related to the optical output and 

modulation band-width of a light-emitting diode. The dependence of the Zn effective 

diffusion coefficient upon the carrier concentration of the p-active layer and n-cladding 

layer was investigated. It was concluded that a suitable growth temperature for the active 

layer was about 850C. 

M.Yoneda, Y.Nakamura, A.Tsushi, K.Ichimura: Japanese Journal of Applied Physics, 

1993, 32[1-9A], 3770-4 

[446-109/110-025] 

250


GaAlAs: Zn Diffusion 

The entry of Zn into Ga 0.7 Al 0.3 As and single heterostructures was studied. It was found 

that the depth of the diffusion front was proportional to the square root of the diffusion 

time. In the case of heterostructures, the Ga 0.7 Al 0.3 As layer thickness modified the 

relationship by decreasing the junction depth to an extent which was some multiple of the 

layer thickness. The relationship could be used to predict diffusion fronts in double 

heterostructures. 

H.J.Yoo, Y.S.Kwon: Journal of Electronic Materials, 1988, 17[4], 337-9 

[446-62/63-203] 


Samples of Ga 0.62 Al 0.38 As were diffused with Zn, via a 200 to 300nm protective ZrO 2 

layer. The diffusion depth exhibited a square-root time dependence. The absolute 

diffusivity values depended slightly upon the diffusion conditions. The layer had 

essentially no effect upon the carrier concentration profile or the activation energy. 

J.E.Bisberg, A.K.Chin, F.P.Dabkowski: Journal of Applied Physics, 1990, 67[3], 1347-51 

[446-74-003] 


Samples of Ga 0.7 Al 0.3 As were grown onto a Si substrate and were diffused with Zn from a 

ZnAs 2 source. It was found that the Zn diffusivity was greater in these layers than in 

layers which were grown on a GaAs substrate. The effective diffusion coefficient was 

related to the defect density in the GaAlAs, and to the diffusion depth. A simple model 

showed that the diffusivity along dislocations was some 5 times higher than that in 

dislocation-free bulk material. 

S.Sakai, Y.Terauchi, N.Wada, Y.Shintani: Japanese Journal of Applied Physics, 1991, 

30[9A], 1942-3 

[446-84/85-011] 


A simple method for the open-tube diffusion of Zn from (ZnO) x (SiO 2 ) 1-x film sources, 

and into Ga 0.8 Al 0.2 As was described. The oxide films were deposited by using metalorganic 

chemical vapor deposition. A capping layer of SiO 2 was deposited on top of the 

source films, and diffusion was carried out in flowing N at 650C. Diffusion depths of 

between 200nm and several microns could be easily obtained. The diffusion front in n- 

type substrates was sharp. The dependence of the diffusion depth upon the source film 

composition (for x-values of 0.04 to 1) was determined by using sectioning methods. 

D.J.Lawrence, F.T.Smith, S.T.Lee: Journal of Applied Physics, 1991, 69[5], 3011-5 

[446-78/79-002] 

251


AlGaAs/GaAs: Zn Diffusion 

The diffusivity of ion-implanted Zn was deduced from secondary ion mass spectrometry 

profiles. Diffusion annealing was carried for various times at 750C. It was found that the 

diffusivity of Zn was proportional to the square of the Zn concentration. This implied the 

existence of local thermal equilibrium. The absolute values were 200 times smaller than 

those which had been reported for gaseous-source Zn diffusion at 650C in GaAs. The 

superlattice disordering rate increased with increasing Zn concentration and was 

attributed to the diffusion of positively charged interstitials such as Ga n+ or Al n+ , where n 

was between 2 and 3. 

E.P.Zucker, A.Hashimoto, T.Fukunaga, N.Watanabe: Applied Physics Letters, 1989, 

54[6], 564-6 

[446-64/65-159] 


It was recalled that previous work had indicated that (Si 2 ) x (GaAs) 1-x could be formed 

within the GaAs quantum well of an Al x Ga 1-x As-GaAs quantum well heterostructure. It 

was shown here that the Si concentration in the (Si 2 ) x (GaAs) 1-x layer greatly exceeded 

typical doping levels. The stability of the quantum well heterostructures was investigated 

with respect to thermal annealing and Zn impurity-induced layer disordering. The results 

showed that the (Si 2 ) x (GaAs) 1-x alloy was stable with respect to thermal annealing unless 

a rich source of Ga vacancies was provided. Also, relatively low-temperature Zn diffusion 

greatly enhanced the disordering of the alloy layer. 

L.J.Guido, N.Holonyak, K.C.Hsieh, R.W.Kaliski, J.E.Baker, D.G.Deppe, R.D.Burnham, 

R.L.Thornton, T.L.Paoli: Journal of Electronic Materials, 1987, 16[1], 87-91 

[446-51/52-111] 


It was pointed out that the characteristic features of Zn diffusion in GaAs substrates and 

GaAs/AlGaAs superlattices could be explained in terms of a kick-out mechanism in 

which the doubly positively charged Ga self-interstitial governed Ga self-diffusion. Such 

characteristics included a square-law dependence of the Zn diffusivity upon its own 

background concentration under Zn iso-concentration diffusion conditions, various Zn indiffusion 

profiles, much lower diffusivities of Zn under out-diffusion conditions than 

under in-diffusion conditions, and a huge enhancement of Zn in-diffusion during 

GaAs/AlGaAs superlattice disordering. It was found that the Longini mechanism was able 

to explain most of these features. 


[446-86/87-009] 


Data were presented on the reduction of layer intermixing in quantum-well 

heterostructures, during high-temperature annealing, by using an initial low-temperature 

252

Zn (Al,Ga)As General 

blocking diffusion of Zn. Room-temperature photoluminescence measurements of the 

increase in the lowest electron to heavy-hole transition energy in the quantum-wells were 

used to characterize the extent of layer intermixing. Doped (C and Si) samples which had 

been annealed (850C, 12h) after a low-temperature blocking Zn diffusion (480C) 

exhibited reductions in energy shift from about 0.177eV, to as little as 0.018eV. Similar 

effects were observed, but to a lesser extent, in the case of undoped samples. The 

improved thermal stability was attributed to a Zn diffusion-induced reduction in the 

number of column-III vacancies in the active layers. This was confirmed by secondaryion 

mass spectroscopy measurements. 

M.R.Krames, A.D.Minervini, E.I.Chen, N.Holonyak, J.E.Baker: Applied Physics Letters, 

1995, 67[13], 1859-61 

[446-125/126-112] 

GaAlAs/GaAs: Zn Diffusion 

The migration of thin highly p-doped layers in single and double heterostructures, grown 

using metalorganic vapor-phase epitaxy, was studied using capacitance-voltage etch 

profiling and secondary ion mass spectrometry. It was deduced that the diffusivity of Zn 

in Ga 0.7 Al 0.3 As could be described by: 

D (cm 2 /s) = 1.5 x 10 -3 exp[-2.2(eV)/kT] 

for rapid thermal annealing. A model which was based upon an interstitial cum 

substitutional diffusion mechanism, with certain kinetic limitations, was successfully used 

to simulate the observed dopant concentration profiles. Markedly anomalous diffusion of 

Zn, from GaAs and into highly n-doped GaAlAs, was found. 

N.Nordell, P.Ojala, W.H.Van, G.Landgren, M.K.Linnarsson: Journal of Applied Physics, 

1990, 67[2], 778-86 

[446-74-004] 

- miscellaneous 

AlGaAs/GaAs: Diffusion 

A model was presented which accounted for the anomalous diffusion of p-type dopants 

during the growth of bipolar transistors. The model was based upon Fermi-level pinning 

at the crystal surface during epitaxial growth. This led to an increased concentration of 

column-III interstitial defects in heavily n-type AlGaAs or GaAs. The excess column-III 

interstitials which were generated in the n-type crystal then flowed into the p + base region 

and led to a transfer of p-type impurity atoms from column-III lattice sites to interstitial 

positions, via a kick-out mechanism. Once located in interstitial positions, the impurity 

atoms diffused rapidly. The model was consistent with previously proposed mechanisms 

for both impurity diffusion and column-III self-diffusion. 

D.G.Deppe: Applied Physics Letters, 1990, 56[4], 370-2 

[446-74-004] 

253

Surface (Al,Ga)As Surface 


AlGaAs: Surface Diffusion 

It was found that surface migration was effectively enhanced by evaporating Ga or Al 

atoms onto a clean surface under an As-free atmosphere or low As pressure. This 

characteristic was exploited by alternately supplying Ga and/or Al and As to the substrate 

surface in order to grow atomically-flat GaAs-AlGaAs hetero-interfaces, and also to grow 

high-quality AlGaAs layers at very low temperatures. The migration characteristics of 

surface adatoms were investigated by using reflection high-energy electron diffraction 

measurements. It was found that differing growth mechanisms operated at high and low 

temperatures. Both mechanisms were expected to yield flat heterojunction interfaces. By 

applying this method, GaAs-AlGaAs single quantum-well structures could be grown at 

substrate temperatures of 200 and 300C, respectively. 

Y.Horikoshi, M.Kawashima, H.Yamaguchi: Japanese Journal of Applied. Physics, 1988, 

27[2], 169-79 

[446-60-001] 

AlGaAs/GaAs: Surface Diffusion 

Ribbed crystals could be grown in a single processing step because the ribs were defined 

by the two non-growing (111)B surfaces which developed at each edge of (011) mesas on 

a patterned GaAs substrate during the organometallic chemical vapor deposition of 

GaAs/AlGaAs structures. The study revealed the importance of surface diffusionenhanced 

crystal growth when a growth surface was adjacent to a non-growing surface 

such as a (111)B facet. The magnitude of this effect suggested that the present deposition 

technique was well-suited to the growth of structures which were tapered in 3 

dimensions. 

E.Colas, A.Shahar, W.J.Tomlinson: Applied Physics Letters, 1990, 56[10], 955-7 

[446-74-005] 

AlGaAs/GaAs: Surface Diffusion 

Monte Carlo methods were used to model the adatom migration of cations in Al x Ga 1-x As, 

where x was equal to 0, 0.5 or 1, and the resultant atomic arrangements on a reconstructed 

As-stabilized GaAs(001) surface with an adatom coverage of up to 0.5. It was found that 

the cation adatom migration depended strongly upon the adatom coverage of the surface. 

Randomly impinging cations occupied lattice sites on the As dimers at coverages of less 

than 0.1. As the coverage was increased from 0.1 to 0.3, the impinging cations migrated 

mainly along the missing dimer rows. At a coverage of more than 0.3, adatoms tended to 

favor lattice sites on the As dimers; including those with non-tetrahedral coordination. In 

the case of Al 0.5 Ga 0.5 As, lattice sites along the missing dimer were occupied mainly by Al 

adatoms, while those on As dimers were favored by Ga adatoms. This was because Al 

adatom migration was several times slower than Ga adatom migration. The resultant 

254

Surface (Al,Ga)As Interdiffusion 

atomic arrangements were explained in terms of a coverage dependence of the migration 

potential. 

T.Ito, K.Shiraishi, T.Ohno: Applied Surface Science, 1994, 82-83, 208-13 

[446-121/122-047] 


1.0E-14 

1.0E-15 

1.0E-16 

table 3 

table 4 

table 5 

table 6 

table 32 

D (cm 2 /s) 

1.0E-17 

1.0E-18 

1.0E-19 

1.0E-20 

7 8 9 10 

10 4 /T(K) 

Figure 1: Interdiffusivity in AlGaAs/GaAs 

AlGaAs/GaAs: Interdiffusion 

A systematic study was made of impurity-free Al-Ga interdiffusion in superlattices which 

were sealed into ampoules. Four structures were used, with superlattice periods that 

ranged from 9 to 52nm. Three ambients were explored: along the Ga-rich solidus, with no 

excess Ga or As in the evacuated ampoule, or with an excess As content which was less 

than that required to reach the As-rich solidus limit. In each of the ambients, the 

Arrhenius dependence of the Al-Ga interdiffusion coefficient could be described by a 

single activation energy at temperatures ranging from 700 to 1050C. Excellent agreement 

was obtained for the Al-Ga interdiffusion coefficients which were measured by using 

superlattices on Si-doped and undoped GaAs substrates. By normalization to a constant 

255

Interdiffusion (Al,Ga)As Interdiffusion 

As over-pressure of 1atm, the Ga- and As-rich activation energies were deduced to be 

3.26 and 4.91eV, respectively. These activation energies were in the range which was 

predicted for Al-Ga interdiffusion, mediated by group-III vacancy second-nearest 

neighbor hopping. An increase in energy which occurred upon going from Ga- to As-rich 

conditions was attributed to a shift in the Fermi-level position, towards the valence band; 

with an increase in the ionized group-III vacancy concentration. 

B.L.Olmsted, S.N.Houde-Walter: Applied Physics Letters, 1993, 63[4], 530-2 

[446-106/107-016] 


Impurity-induced layer disordering experiments were performed on quantum-well 

heterostructures which were heavily doped with C. The results showed that the presence 

of C retarded Al and Ga interdiffusion, as compared with un-doped material. 

Interdiffusion in C-doped quantum-well heterostructures was not enhanced by the use of 

a Ga-rich rather than an As-rich annealing ambient. The data were inconsistent with most 

Fermi-level effect models for layer disordering which did not include a chemical impurity 

dependence or sub-lattice dependence, and which did not take account of the possibility 

of inhibited Al and Ga interdiffusion in extrinsic crystals. 

L.J.Guido, B.T.Cunningham, D.W.Nam, K.C.Hsieh, W.E.Plano, J.S.Major, E.J.Vesely, 

A.R.Sugg, N.Holonyak, G.E.Stillman: Journal of Applied Physics, 1990, 67[4], 2179-82 

[446-74-004] 


Impurity-induced layer disordering or intermixing in quantum well heterostructures and 

superlattices were reviewed. On the basis of the behavior of column-III vacancies and 

interstitials, suitable models for layer disordering were developed. 

D.G.Deppe, N.Holonyak: Journal of Applied Physics, 1988, 64[12], R93-113 

[446-72/73-005] 

34 AlGaAs/GaAs: Interdiffusion 

Data were presented which showed that the Al-Ga interdiffusion coefficient for an 

Al x Ga 1-x As-GaAs quantum-well heterostructure or a superlattice was highly dependent 

upon the crystal encapsulation conditions (table 4). The activation energy for Al-Ga 

interdiffusion, and thus layer-disordering, was smaller (about 3.5eV) for dielectric 

encapsulated samples than after capless annealing (about 4.7eV). The interdiffusion 

coefficient for Si 3 N 4 -capped samples was almost an order of magnitude smaller than for 

the cases of capless or SiO 2 -capped samples at temperatures of between 800 and 875C. 

As well as the type of encapsulant, the encapsulation geometry (stripes or capped stripes) 

was important because of strain effects. These were a major source of anisotropic Al-Ga 

interdiffusion. 

L.J.Guido, N.Holonyak, K.C.Hsieh, R.W.Kaliski, W.E.Plano, R.D.Burnham, 

R.L.Thornton, J.E.Epler, T.L.Paoli: Journal of Applied Physics, 1987, 61[4], 1372-9 

[446-60-002] 

256


Table 4 

Interdiffusivity (Al-Ga) in AlGaAs/GaAs 

Conditions Temperature (C) D (cm 2 /s) 

capless 875 1.1 x 10 -17 

SiO 2 cap 875 9.6 x 10 -18 

capless 850 2.9 x 10 -18 

SiO 2 cap 850 2.9 x 10 -18 

SiO 2 cap 825 2.8 x 10 -18 

capless 825 1.0 x 10 -18 

Si 3 N 4 cap 875 2.2 x 10 -18 

Si 3 N 4 cap 850 7.8 x 10 -19 

SiO 2 cap 800 7.0 x 10 -19 

capless 800 4.5 x 10 -19 

Si 3 N 4 cap 825 3.8 x 10 -19 

Si 3 N 4 cap 800 2.2 x 10 -19 


A model was presented which described the diffusion of Si into GaAs from grown-in 

dopant sources. The effects of background impurities upon Si diffusion and layer 

interdiffusion in the present superlattices were also described. Epitaxial GaAs samples 

with alternating doped and undoped layers, and superlattices with Mg- or Si-doped layers, 

were studied. Various annealing conditions were used to study interactions between 

grown-in impurities and native defects. A model which described impurity diffusion and 

Al-Ga layer interdiffusion was based upon the behavior of column-III vacancies (V III ) and 

interstitials (I III ), and the control of their contents. The results indicated that n-type 

superlattices underwent enhanced layer interdiffusion because of an increased solubility 

of the V III defect. Enhanced layer interdiffusion in p-type superlattices was attributed to 

an enhanced solubility of I III . 

D.G.Deppe, N.Holonyak, W.E.Plano, V.M.Robbins, J.M.Dallesasse, K.C.Hsieh, 

J.E.Baker: Journal of Applied Physics, 1988, 64[4], 1838-44 

[446-72/73-006] 


It was shown that donor diffusion and layer intermixing were greatly enhanced in the 

presence of defects which were created by crystal overgrowth on locally laser-melted 

substrates. Accelerated defect and impurity-induced layer disordering, and donor 

diffusion from a solid SiO 2 source, a Ge vapor source or a grown-in Se source were 

257


observed in regions of high defect density. Enhanced donor diffusion and crystal selfdiffusion 

were attributed to an increased density of column-III defects and dislocations. 

F.A.Kish, W.E.Plano, K.C.Hsieh, A.R.Sugg, N.Holonyak, J.E.Baker: Journal of Applied 

Physics, 1989, 66[12], 5821-5 

[446-74-005] 


A study was made of the in-diffusion of various group-IV and group-VI n-type impurities. 

In all cases, the n-type dopants enhanced the Al-Ga interdiffusion coefficient above that 

which could be attributed to an As over-pressure alone. The Si-induced enhancement had 

previously been attributed to a change in Fermi-level position with doping and could 

therefore account for disordering by other n-type impurities. However, important 

differences were observed in the interdiffusion characteristics that were induced by Si or 

Ge, and by S or Se. The disordering was attributed to an enhancement of the group-III 

vacancy concentration for each of these n-type impurities. This was also true of undoped 

crystals which were disordered by an As ambient alone at 855C. 


[446-106/107-016] 


The As vapor pressure dependence of interdiffusion in a hetero-interface at high 

temperatures was studied by measuring the wavelength shift of the photoluminescence in 

a multi quantum well. It was found that interdiffusion at a temperature of 850C was 

minimized by an As pressure of 100torr and was enhanced at both lower and higher As 

pressures. A degradation of the photoluminescence intensity was observed only at higher 

As pressures. These effects were attributed to the presence of excess Al and Ga 

vacancies, and their associated defects. 

A.Furuya, O.Wada, A.Takamori, H.Hashimoto: Japanese Journal of Applied Physics, 

1987, 26[6], L926-8 

[446-55/56-002] 


Transmission electron microscopy and carrier concentration measurements were used to 

characterize the layer interdiffusion mechanism of a Se-doped Al x Ga 1-x As-GaAs 

superlattice during high-temperature annealing. By varying the annealing environment 

and comparing the results with similarly annealed un-doped superlattices and Mg-doped 

superlattices, it was found that layer interdiffusion occurred via the interaction of the Se 

impurity with native defects which were associated with As-rich conditions. The most 

likely candidate was suggested to be the column-III vacancy. 

D.G.Deppe, N.Holonyak, K.C.Hsieh, P.Gavrilovic, W.Stutius, J.Williams. Applied 

Physics Letters, 1987, 51[8], 581-3 

[446-55/56-002] 

258



A study was made of the role of defect diffusion, from crystal surfaces, in the disordering 

of a multiple quantum well structure that was Si-doped during molecular beam epitaxial 

growth. The distribution of the native defects was deduced from photoluminescence 

spectroscopic, secondary ion mass spectrometric, and electrochemical C-V profiling data. 

No significant difference was observed between the Al-Ga interdiffusion coefficients of 

Si-doped and undoped superlattices when they were annealed with excess Ga. This was 

attributed to the lack of a source of group-III vacancies. Only a small fraction of the 

enhancement which was predicted to result from Si doping was observed when excess As 

was used instead. The largest Fermi-level enhancement was observed when no excess Ga 

or As was present in the evacuated ampoule. The results indicated that the crystal surface 

was both source and sink for the native defects which were known to mediate Al-Ga 

interdiffusion. Significant electrical compensation of the donors was observed after 

annealing in both As-rich or Ga-rich ambients. This was attributed to ionized group-III 

vacancy generation in the former case, and to Si atoms which moved from group-III to 

group-V sites in the latter case. 


[446-106/107-017] 


Spatially resolved values of the Al/Ga interdiffusion coefficient for p-i-n and n-i-p 

AlGaAs-GaAs device structures were found to be almost identical in magnitude, but 

varied with position (by a factor of 2) across a 1µ-thick multiple quantum well active 

region. These observations contrasted with theoretical predictions, given that the Fermi 

level to valence-band energy separation changed by 0.7eV across the intrinsic region, and 

suggested that impurity-free layer disordering did not provide the necessary uniformity in 

energy shift for photonic integrated circuit fabrication. 

S.Seshadri, L.J.Guido, P.Mitev: Applied Physics Letters, 1995, 67[4], 497-9 

[446-123/124-157] 


Direct optical observations were made of diffusion-related deep levels that were 

associated with interdiffusion in superlattice structures. Low-energy cathodoluminescence 

spectroscopy was used to investigate the formation and evolution of deep levels under 

various conditions. It was found that the spatial distribution of the deep levels was 

strongly related to the extent of superlattice intermixing, as measured using secondary ion 

mass spectrometry and photoluminescence spectroscopy. The results strongly suggested 

that a larger interdiffusion rate of the Si-induced layer intermixing was related to the 

formation of a deep level which was associated with an optical emission at 1.3eV. 

R.E.Viturro, B.L.Olmsted, S.N.Houde-Walter, G.W.Wicks: Journal of Vacuum Science 

and Technology B, 1991, 9[4], 2244-50 

[446-88/89-008] 

259



The effect of pressure and stoichiometry upon the Al-Ga interdiffusion of undoped 

multiple quantum wells was investigated over the entire composition range of tile GaAs 

solidus. The occurrence of a two orders of magnitude increase in the interdiffusion 

coefficient suggested that interdiffusion in an intrinsic crystal was mediated mainly by 

column-III vacancies over the whole solidus range. It was noted that the 

photoluminescence intensity of the Ga-rich crystal was more than 3 orders of magnitude 

greater than that of the As-rich crystal. 


[446-88/89-008] 

35 AlGaAs/GaAs: Interdiffusion 

The interdiffusion of Ga and Al in AlGaAs alloys which were subjected to various 

annealing temperatures, times and environments was considered. The interdiffusion 

coefficients (table 5) and activation energies were determined by relating shifts in the 

photoluminescence peaks to calculated transition energies which were based upon an erf 

composition profile. It was noted that a Ga over-pressure reduced interdiffusion whereas 

an As over-pressure increased interdiffusion. This was thought to be the first study of the 

effect of a Ga over-pressure upon the interdiffusion of Al and Ga in AlGaAs. 

K.Y.Hsieh, Y.C.Lo, J.H.Lee, R.M.Kolbas: Institute of Physics Conference Series, 1989, 

96, 393-6 

[446-125/126-120] 

Table 5 

Interdiffusivity (Al-Ga) in AlGaAs/GaAs 


940 9.0 x 10 -18 

940 5.9 x 10 -18 

920 2.5 x 10 -18 

920 1.9 x 10 -18 

890 1.2 x 10 -18 

845 1.9 x 10 -19 


Quantum-well heterostructures were annealed in an AsH 3 /H 2 atmosphere, and the use of 

photoluminescence spectroscopy revealed a uniform and reproducible increase in the 

effective quantum-well band-gap. The energy shift data indicated that Al/Ga 

interdiffusion occurred under non-equilibrium conditions. The activation energies varied 

from about 5.2eV, in the equilibrium case, to about 3.4eV in the non-equilibrium case. 

260


S.Seshadri, L.J.Guido, T.S.Moise, J.C.Beggy, T.J.Cunningham, R.C.Barker, R.N.Sacks: 


[446-93/94-004] 

GaAlAs/GaAs: Interdiffusion 

A study was made of Al-Ga interdiffusion and C acceptor diffusion in C-doped 

Ga 0.6 Al 0.4 As/GaAs superlattices which had been annealed, under various ambient As 4 

pressure conditions, at temperatures ranging from 825 to 960C. The superlattices were 

doped with C to an initial acceptor concentration of about 2.9 x 10 19 /cm 3 . The Al-Ga 

interdiffusion was found to be most predominant in Ga-rich annealing ambients. The 

interdiffusivity values were about 2 orders of magnitude lower than those predicted by the 

Fermi-level effect model. In As-rich ambients, the interdiffusion values were in 

approximate agreement with those which were predicted by the Fermi-level effect model. 

By analyzing measured hole concentration profiles, it was deduced that both C acceptor 

diffusion and reduction occurred during annealing. Both the C acceptor diffusivity and 

the C acceptor reduction coefficient data could be characterized approximately by a ¼- 

power dependence upon the As 4 pressure. These pressure dependences indicated that C 

diffused via the interstitialcy or interstitial-substitutional mechanism, while hole reduction 

was governed by a C acceptor precipitation mechanism. 

H.M.You, T.Y.Tan, U.M.Gösele, S.T.Lee, G.E.Höfler, K.C.Hsieh, N.Holonyak: Journal 

of Applied Physics, 1993, 74[4], 2450-60 

[446-109/110-028] 

GaAlAs/GaAs: Interdiffusion 

Superlattices of Ga 0.7 Al 0.3 As/GaAs which had been grown by means of metalorganic 

chemical vapor deposition, and which were heavily doped with C using CCl 4 , were 

annealed (825C, 24h) under various ambients and encapsulants. Photoluminescence 

monitoring at 1.7K was used to determine approximate interdiffusion coefficients, D Al-Ga , 

for various annealing conditions. For all of the encapsulants which were studied, D Al-Ga 

increased with increasing As 4 pressure in the annealing ampoule. This result disagreed 

with trends which had been reported for Mg-doped crystals, and with the predictions of 

the charged point-defect (Fermi-level) model. It was noted that a Si 3 N 4 cap provided the 

most effective surface protection against ambient-stimulated layer interdiffusion (D Al-Ga = 

1.5 x 10 -19 to 3.9 x 10 -19 cm 2 /s). The most extensive layer intermixing occurred for 

uncapped superlattices which were annealed in an As-rich ambient (D Al-Ga equal to about 

3.3 x 10 -18 cm 2 /s). These values were up to 40 times greater than those which had 

previously been reported for nominally undoped AlGaAs/GaAs superlattices. This 

implied that doping slightly enhanced layer intermixing; but this was significantly less 

than that predicted by the Fermi-level effect. The discrepancies between the experimental 

data and the model were considered. Marked changes in the optical properties of the 

annealed superlattices, as a function of storage time at room temperature, were also 

reported. It was suggested that these changes might reflect a degraded thermal stability of 

the annealed crystals, due to lattice defects which were generated at high temperatures. It 

261


was proposed that this was related to the failure to prepare buried heterostructure 

quantum-well lasers, via impurity-induced layer disordering, in similar doped crystals. 

I.Szafranek, M.Szafranek, J.S.Major, B.T.Cunningham, L.J.Guido, N.Holonyak, 

G.E.Stillman: Journal of Electronic Materials, 1991, 20[6], 409-18 

[446-91/92-005] 

36 GaAlAs/GaAs: Interdiffusion 

The migration of Al and Ga in Ga 0.6 Al 0.4 As/GaAs quantum wells was investigated by 

measuring the photoluminescence of samples which had been annealed at temperatures 

ranging from 850 to 1065C; with and without a SiO 2 cap. At 1000C, under a SiO 2 cap, 

the Al-Ga interdiffusion coefficient was at least 2 orders of magnitude higher for a 

GaAlAs/GaAs quantum well than for an InAlGaP/GaInP quantum well, within the same 

sample. By comparing the calculated photoluminescence shifts with measured values, an 

activation energy of 4.5eV was estimated (table 6) for Al-Ga interdiffusion in a 

GaAlAs/GaAs quantum well under a SiO 2 cap. 

K.J.Beernink, D.Sun, D.W.Treat, B.P.Bour: Applied Physics Letters, 1995, 66[26], 3597- 

9 

[446-125/126-120] 

Table 6 

Interdiffusivity (Al-Ga) in GaAlAs/GaAs 

Conditions Temperature (C) D (cm 2 /s) 

RTA/SiO 2 1065 2.5 x 10 -15 

RTA/SiO 2 1025 1.8 x 10 -15 

RTA/SiO 2 1000 4.3 x 10 -16 

no SiO 2 1000 4.1 x 10 -17 

RTA/SiO 2 975 1.8 x 10 -16 

FA/SiO 2 925 3.0 x 10 -17 

FA/SiO 2 900 2.3 x 10 -17 

FA/SiO 2 870 4.7 x 10 -18 

FA/SiO 2 845 1.6 x 10 -18 

262

(Al,Ga,In)P 

Si 

InAlGaP/GaAs: Si Diffusion 





column-V atoms in various Si-diffused heterostructures which were closely latticematched 

to GaAs. An enhancement of lattice atom self-diffusion, due to impurity 







[446-64/65-157] 

Zn 

AlGaInP: Zn Diffusion 

The behavior of Zn acceptors in p-type layers during thermal annealing in gaseous AsH 3 - 

H 2 mixtures was investigated. It was found that the electrical activity of Zn acceptors was 

greatly affected by H atoms which originated from the arsine during annealing. In 

addition, observations of atomic disordering suggested that H passivation of Zn acceptors 

suppressed atomic diffusion. The results could be consistently explained by a simple H 

passivation model which involved the termination of dangling bonds by H atoms. 

A.Ishibashi, M.Mannoh, I.Kidoguchi, Y.Ban, K.Ohnaka: Applied Physics Letters, 1994, 

65[10], 1275-7 

[446-119/120-188] 

263

Interdiffusion (Al,Ga,In)P|(Al,Ga)Sb Be 


InAlGaP/GaInP: Interdiffusion 

The migration of Al and Ga in In 0.5 Al 0.3 Ga 0.2 P/Ga 0.6 In 0.4 P quantum wells was 

investigated by measuring the photoluminescence of samples which had been annealed at 

temperatures ranging from 850 to 1065C; with and without a SiO 2 cap. At 1000C, under a 

SiO 2 cap, the Al-Ga interdiffusion coefficient was at least 2 orders of magnitude higher 

for a GaAlAs/GaAs quantum well, than for an InAlGaP/GaInP quantum well, within the 

same sample. 

K.J.Beernink, D.Sun, D.W.Treat, B.P.Bour: Applied Physics Letters, 1995, 66[26], 3597- 

9 

[446-125/126-120] 

(Al,Ga)Sb 

Be 

AlGaSb: Be Diffusion 

The ion-implantation p-type doping of Al 0.75 Ga 0.25 Sb was studied. The surface 

morphology and electrical properties were shown, by using atomic force microscopy and 

Hall measurements, to be degraded after rapid thermal annealing at 650C. The 

implantation of Be resulted in sheet hole concentrations which were twice those of the 

implanted acceptor dose (10 13 /cm 2 ) after 600C annealing. This was attributed to double 

acceptor or antisite defect formation. Implanted C acted as an acceptor, but also 

demonstrated an excess hole conduction which was attributed to implantation-induced 

defects. Implanted Zn required a higher annealing temperature than did Be, in order to 

achieve 100% effective activation for a dose of 10 13 /cm 2 . It was suggested that this was 

probably the result of the greater implantation-induced damage that was created by the 

264

Be (Al,Ga)Sb Mg 

heavier Zn ion. The secondary ion mass spectroscopy of as-implanted and annealed Be, 

Mg and C samples was studied. The diffusion of implanted Be (5 x 10 13 cm 2 , 45keV) was 

shown to exhibit an inverse dependence upon temperature. This was attributed to a 

substitutional-interstitial diffusion mechanism. Implanted C (2.5 x 10 14 /cm 2 , 70keV) 

exhibited no redistribution, even after annealing at 650C. 

J.C.Zolper, J.F.Klem, A.J.Howard, M.J.Hafich: Journal of Applied Physics, 1996, 79[3], 

1365-70 

[446-131/132-168] 

Mg 

AlGaSb: Mg Diffusion 

The ion-implantation p-type doping of Al 0.75 Ga 0.25 Sb was studied. The surface 

morphology and electrical properties were shown, by using atomic force microscopy and 

Hall measurements, to be degraded after rapid thermal annealing at 650C. The 

implantation of Mg resulted in sheet hole concentrations which were twice those of the 

implanted acceptor dose (10 13 /cm 2 ) after 600C annealing. This was attributed to double 

acceptor or antisite defect formation. Implanted C acted as an acceptor, but also 

demonstrated an excess hole conduction which was attributed to implantation-induced 

defects. Implanted Zn required a higher annealing temperature than did Mg, in order to 

achieve 100% effective activation for a dose of 10 13 /cm 2 . It was suggested that this was 

probably the result of the greater implantation-induced damage that was created by the 

heavier Zn ion. The secondary ion mass spectroscopy of as-implanted and annealed Mg 

and C samples was studied. Implanted Mg (10 14 /cm 2 , 110keV) exhibited marked 

redistribution, and losses of up to 56% at the interface, after 600C annealing. Implanted C 

(2.5 x 10 14 /cm 2 , 70keV) exhibited no redistribution, even after annealing at 650C. 

J.C.Zolper, J.F.Klem, A.J.Howard, M.J.Hafich: Journal of Applied Physics, 1996, 79[3], 

1365-70 

[446-131/132-168] 

265

(Al,In)As 

Be 

InAlAs: Be Diffusion 

The behavior of implanted Be + ions was investigated during rapid thermal annealing at 

temperatures of between 600 and 900C. It was found that the apparent activation energy 

for Be was equal to 0.43eV. Lower activation efficiencies of the dopant were found in 

InAlAs, as compared with InGaAs. Anomalously low activation was detected for lowdose 

Be implants. The latter effect was attributed to a lack of vacant sites for the Be 

atoms to occupy. Extensive redistribution of the Be was observed after annealing (750C, 

10s). 

E.Hailemariam, S.J.Pearton, W.S.Hobson, H.S.Luftman, A.P.Perley: Journal of Applied 

Physics, 1992, 71[1], 215-20 

[446-86/87-038] 

Fe 

InAlAs: Fe Diffusion 

Single (200keV) and multiple-energy Fe implantation of n-type material was carried out 

on In 0.52 Al 0.48 As at room temperature or 200C. Secondary ion mass spectrometry profiles 

revealed marked out-diffusion during all of the rapid thermal annealing treatments which 

were used; regardless of the implantation temperature. The Fe implantation peaks which 

were observed after the annealing of other In-based compounds were not observed here. 

J.M.Martin, R.K.Nadella, M.V.Rao, D.S.Simons, P.H.Chi, C.Caneau: Journal of 

Electronic Materials, 1993, 22[9], 1153-8 

[446-109/110-039] 

Si 

InAlAs: Si Diffusion 

The behavior of implanted Si + ions was investigated during rapid thermal annealing at 

temperatures of between 600 and 900C. The apparent activation energy for Si was equal 

266

Si (Al,In)As General 

to 0.58eV. Lower activation efficiencies of the dopant were found in InAlAs, as 

compared with InGaAs. The Si underwent no migration, even after annealing at 850C. 


Physics, 1992, 71[1], 215-20 

[446-86/87-038] 

Ti 

InAlAs: Ti Diffusion 

The Ti implantation of p-type material was carried out on In 0.52 Al 0.48 As at room 

temperature or 200C. The implanted Ti exhibited only slight in-diffusion and outdiffusion 

after both room temperature and 200C implantation. Rutherford back-scattering 

measurements of annealed samples which had been implanted at 200C revealed a crystal 

quality which was similar to that of virgin material. The resistivity of all of the samples 

after annealing was greater than 10 6 Ωcm. 

J.M.Martin, R.K.Nadella, M.V.Rao, D.S.Simons, P.H.Chi, C.Caneau: Journal of 


[446-109/110-039] 

General 

InAlAs: Diffusion 

It was shown that systematic variations in the experimental parameters could turn multilayers 

into so-called microscopic laboratories for the study of point defects. In this way, 

the effects of composition, doping and strain could be separated. It was also possible to 

determine the nature and charge state of the mediating defect. The present results showed 

that interdiffusion in the present system was mediated by a double-acceptor vacancy-like 

defect. The activation energy, in the case of In 0.52 Al 0.48 As, was equal to 4eV. Its value 

varied to the extent of 0.051eV for every percent of strain. 

F.H.Baumann, J.H.Huang, J.A.Rentschler, T.Y.Chang, A.Ourmazd: Physical Review 

Letters, 1994, 73[3], 448-51 

[446-115/116-141] 

267

AlN 

1.0E-07 

1.0E-08 

1.0E-09 

D (cm 2 /s) 

1.0E-10 

1.0E-11 

1.0E-12 

1.0E-13 

table 7 

table 8 

table 9 

1.0E-14 

1.0E-15 

1.0E-16 

4 5 6 

10 4 /T(K) 

Figure 2: Diffusivity of O in AlN 

D 

AlN: D Diffusion 

The out-diffusion of H was studied, using 2 H plasma-treated (250 or 400C, 0.5h) or 2 H + - 

implanted samples, during annealing at temperatures ranging from 300 to 900C. 

268

D AlN O 

Secondary ion mass spectrometry was used to measure the resultant distributions. At 

concentrations that were greater than 10 20 /cm 3 , there was a near-surface (less than 0.3µ) 

region that was probably due to the formation of platelet defects. At concentrations of 

about 10 18 /cm 3 , a plateau region was present which extended throughout the film 

thickness of about 1µ. This was attributed to the pairing of 2 H with point defects. In 

implanted samples, 2 H redistribution occurred in the same manner as the bulk population 

in plasma-treated material. The thermal stability of the D profile in the nitride was much 

higher than that in GaAs and similar compounds. 

R.G.Wilson, S.J.Pearton, C.R.Abernathy, J.M.Zavada: Journal of Vacuum Science and 

Technology A, 1995, 13[3], 719-23 

[446-140-029] 

Table 7 

Diffusivity of O in AlN 


1500 8.0 x 10 -16 

1600 3.5 x 10 -15 

1700 6.7 x 10 -15 

1800 1.3 x 10 -14 

1900 1.9 x 10 -14 

N 

AlN: N Permeation 

A method was described for the determination of ion migration numbers on the basis of 

permeability data. The latter data were here obtained by using an emf technique. The 

permeability was equal to 8.66cm 2 /MPa-s. It was pointed out that the present method 

could establish the presence of an ionic component of the conductivity for nitrides, and its 

importance could be estimated for cases where a dense ceramic could not be prepared. 

R.P.Lesunova, L.S.Karenina, V.K.Gilderman, S.F.Palguev: Izvestiya Akademii Nauk 

SSSR - Neorganicheskie Materialy, 1989, 25[11], 1926-8. (Inorganic Materials, 1989, 

25[11], 1633-5) 

[446-76/77-131] 

O 

37 AlN: O Diffusion 

Interdiffusion of N and O was investigated by means of electron energy-loss spectroscopy 

and transmission electron microscopy. Diffusion couples, Al 2 O 3 /AlN, were prepared by 

oxidizing AlN ceramics, and were annealed at temperatures ranging from 1500 to 1900C 

269

O AlN O 

(table 7). The couples were encapsulated in a Ta ampoule in order to ensure an inert 

atmosphere, and the O concentration profiles across the oxide/nitride interface were 

measured by means of electron energy-loss spectroscopy. It was found that O/N 

interdiffusion in AlN could be described by: 

D(cm 2 /s) = 1.5 x 10 -8 exp[-240(kJ/mol)/RT] 

The magnitude and temperature-dependence of the interdiffusion were comparable to 

those which had been reported for other non-oxide ceramic materials. Under normal AlN 

sintering conditions, the O/N interdiffusion was too slow to provide an effective means 

for O removal from AlN grains. 

M.Sternitzke, G.Müller: Journal of the American Ceramic Society, 1994, 77[3], 737-42 

[446-123/124-267] 

Table 8 

Bulk Diffusivity of O in AlN 


1600 3.43 x 10 -14 

1700 1.00 x 10 -13 

1800 4.27 x 10 -13 

1900 1.43 x 10 -12 

38,9 AlN: O Grain Boundary Diffusion 

The diffusion of O in commercial nitride substrates was studied by carrying out 

interdiffusion-type experiments. The diffusion of O within AlN grains and along grain 

boundaries was investigated. By using as-received AlN substrates and an electron 

microprobe as an analytical tool, it was found that the rate of O diffusion along grain 

boundaries was strongly affected by the presence of impurities and/or other phases at 

these boundaries. The diffusion of O within AlN grains was studied at temperatures of 

between 1600 and 1900C (table 8) in a flowing N atmosphere by using secondary ion 

mass spectrometry to determine O concentration profiles. The chemical diffusion 

coefficient of O in the AlN lattice was described by: 

log[D(cm 2 /s)] = -1.68 - 427(kJ/mol)/2.303RT 

The O concentration profiles which were determined by means of secondary ion mass 

spectrometry also revealed a contribution that arose from diffusion along grain boundaries 

(table 9). It was therefore possible to determine values of the product of grain boundary 

width and grain boundary O diffusivity. 

H.Solmon, D.Robinson, R.Dieckmann: Journal of the American Ceramic Society, 1994, 

77[11], 2841-8 

[446-123/124-268] 

270

Interdiffusion AlN Interdiffusion 


AlN/Al 2 OC: Interdiffusion 

It was recalled that AlN and Al 2 OC were isostructural. The AlN-Al 2 OC system formed 

homogeneous solid solutions above 1950C. Interdiffusion was studied in the solidsolution 

regime in order to clarify differences in the kinetics of phase separation when 

samples were annealed at lower temperatures. The diffusion couples, (AlN) 0.7 (Al 2 OC) 0.3 

/(AlN) 0.3 (Al 2 OC) 0.7 , were prepared by hot pressing and were annealed at 2273K. It was 

found that the interdiffusion coefficients in this system were much larger than those in the 

SiC-AlN system. 

Q.Tian, A.V.Virkar: Journal of the American Ceramic Society, 1996, 79[8], 2168-74 

[446-150/151-214] 

Table 9 

Grain Boundary Diffusivity of O in AlN 

Temperature (C) 

D (cm 2 /s) 

1700 9.50 x 10 -10 

1800 1.59 x 10 -9 

1900 1.78 x 10 -8 

AlN/SiC: Interdiffusion 

It was recalled that AlN and 2H-type SiC were isostructural. The SiC-AlN system formed 

homogeneous solid solutions above 2000C. Interdiffusion was studied in the solidsolution 

regime in order to clarify differences in the kinetics of phase separation during 

annealing at lower temperatures. Diffusion couples, (SiC) 0.3 (AlN) 0.7 /(SiC) 0.7 (AlN) 0.3 , 

were prepared by hot pressing. Interdiffusion coefficients were measured at 2373, 2473 

and 2573K, and the corresponding activation energy was estimated to be 632kJ/mol. 

Q.Tian, A.V.Virkar: Journal of the American Ceramic Society, 1996, 79[8], 2168-74 

[446-150/151-214] 

271

BN 

D 

BN: D Diffusion 

Re-emission curves and thermodesorption spectra were measured for D ions which had 

been implanted into this material. The thermodesorption spectra consisted of several 

peaks at temperatures ranging from 400 to 1100K. The re-emission curves could be 

described by a simple mathematical model which included the effects of diffusion, 

second-order thermodesorption and defect trapping. The recycling factor and defecttrapping 

factor were found to depend exponentially upon the temperature, at temperatures 

above 600K. They deviated from this behavior at room temperature. It was supposed that 

radiation-enhanced and thermally-activated processes predominated at room and high 

temperatures, respectively. 

A.A.Pisarev, V.M.Smirnov, S.K.Zhdanov, A.V.Varava, V.V.Bandurko: Journal of 

Nuclear Materials, 1992, 187[3], 254-9 

[446-91/92-075] 

272

GaAs 

1.0E-07 

1.0E-08 

1.0E-09 

1.0E-10 

D (cm 2 /s) 

1.0E-11 

1.0E-12 

1.0E-13 

1.0E-14 

1.0E-15 

1.0E-16 

1.0E-17 

1.0E-18 

1.0E-19 

1.0E-20 

Be (table 10) 

Cr (table 11) 

Fe (table 12) 

Ga (tables 13 to 15) 

Ge (table 16) 

H (tables 17 and 18) 

S (table 19) 

Si (tables 21 to 25) 

Sn (table 26) 

Zn (tables 27 to 31) 

6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 

10 4 /T(K) 

Figure 3: Summary of Diffusivities of Various Elements in GaAs 

Al 

GaAs: Al Diffusion 

An infra-red absorption spectroscopic investigation was made of the thermal diffusion of 

Al from monatomic Al layers which were embedded in a GaAs epitaxial film. After 

273

Al GaAs Al 

annealing, the absorption peak of the 2-dimensionally localized vibrational modes at 

358/cm (due to Al layers) decreased while the peak at 362/cm (due to isolated Al atoms) 

increased. The 362/cm peak height was compared with the fraction of isolated Al atoms, 

as calculated by assuming second-nearest neighbor hopping diffusion from a monatomic 

Al layer and into the GaAs matrix. It was thus deduced that the diffusion coefficient of Al 

in GaAs was equal to 2 x 10 -19 cm 2 /s at 700C. It was concluded that this was a simple and 

reliable method for the investigation of impurity diffusion in crystals. 

H.Ono. N.Ikarashi, T.Baba: Applied Physics Letters, 1995, 66[5], 601-3 

[446-121/122-053] 

GaAs/Al/GaAs: Al Diffusion 

Theoretical and experimental aspects of the growth of heterostructures were investigated. 

In these heterostructures, GaAs was grown on top of the buried metal layer via migrationenhanced 

epitaxy at low temperatures (200 or 400C) in order to provide a kinetic barrier 

to the out-diffusion of Al during super-layer growth. The crystallinity and orientation of 

the Al film which was deposited on (100)GaAs at about 0C was studied by using 

transmission electron diffraction, dark-field imaging, and X-ray diffraction methods. It 

was found that the Al was polycrystalline, with a grain size of about 6nm, and that the 

preferred growth orientation was (111). This could be textured in the plane, but oriented 

out of the plane. The quality of the GaAs super-layer, which was grown on top of the Al 

by means of migration-enhanced epitaxy, was very sensitive to the growth temperature. A 

layer which was grown at 400C had a good structural and optical quality, but was 

associated with considerable Al out-diffusion at the Al/GaAs hetero-interface. At 200C, 

where the interface had good structural integrity, the super-layer exhibited twinning and 

no luminescence was observed. 

P.Bhattacharya, J.E.Oh, J.Singh, D.Biswas, R.Clarke, W.Dos Passos, R.Merlin, 

N.Mestres, K.H.Chang, R.Gibala: Journal of Applied Physics, 1990, 67[8], 3700-5 

[446-78/79-030] 

GaAs/AlAs: Al Diffusion 

The implantation of Be ions into heterostructures at room temperature or liquid N 

temperatures was investigated. It was found that room-temperature implantation created 

dislocation loops at the first interface; a distance which was far short of the maximum 

projected range. Implantation at low temperatures caused twinning. The latter could be 

removed by annealing (900C, 1200s), without leading to the interdiffusion of Al. The 

presence of dislocation networks tended to enhance intermixing. The Be concentrations 

were sufficiently low to prevent Be-induced intermixing. 

S.Mitra: Semiconductor Science and Technology, 1990, 5[11], 1138-40 

[446-76/77-017] 

GaAs/AlGaAs: Al Diffusion 

The effect of room-temperature electron irradiation upon interdiffusion at quantum-well 

interfaces was investigated by using low-temperature cathodoluminescence spectroscopy. 

It was found that interdiffusion was enhanced by the presence of defects which were 

274

Al GaAs As 

generated by irradiation with a 400keV electron beam. After irradiation at room 

temperature to doses of between about 1.5 x 10 17 and 2.5 x 10 17 /cm 2 , followed by rapid 

thermal annealing (900C, 60s), an interdiffusion length of 0.3 to 0.5nm was found. The 

resultant damage tended to saturate with increasing irradiation dose. The formation of 

defect clusters at high doses limited the degree of defect introduction, and therefore the 

extent of interdiffusion at the interface. 

Y.J.Li, M.Tsuchiya, P.M.Petroff: Applied Physics Letters, 1990, 57[5], 472-4 

[446-76/77-018] 

GaAs/AlGaAs: Al Diffusion 

An investigation was made of the proton-implantation enhanced intermixing of quantum 

wells, for H + doses which ranged from 5 x 10 13 to 10 16 /cm 2 . The implantation of 20keV 

H + , followed by high-temperature rapid thermal annealing, led to the enhanced diffusion 

of Al into the GaAs quantum well. Shifts in the electron heavy-hole recombination 

energies, due to compositional changes, were observed by using room-temperature 

cathodoluminescence methods. Diffusion lengths of more than 2nm were deduced from 

the energy shifts in a 5nm well, and were found to depend upon the implanted dose and 

the annealing time. It was suggested that this was to be expected if the enhanced 

interdiffusion was caused by defects which were introduced by implantation. 

G.F.Redinbo, H.G.Craighead, J.M.Hong: Journal of Applied Physics, 1993, 74[5], 3099- 

102 

[446-106/107-079] 

As 

GaAs: As Diffusion 

Atomistic thermodynamic calculations were made of the energetics of self-diffusion. An 

assessment of the activation enthalpy of the saddle-point configuration of various modes 

of vacancy self-diffusion indicated that second-nearest neighbor hopping was the most 

energetically favorable mechanism, if vacancies were available in equilibrium 

concentrations. An assessment of the activation entropy indicated that normal diffusion 

pre-exponential factors, of the order of 10 -5 to 0.1cm 2 /s, were consistent with vacancy 

self-diffusion via second-nearest neighbor hopping. It was proposed that self-diffusion 

which was characterized by pre-exponential factors of the order of 10 7 to 10 8 cm 2 /s, and 

activation energies of the order of 6eV, involved processes in which surface vacancy 

generation was inhibited and self-diffusion was mediated by Frenkel pair generation. 

J.F.Wager: Journal of Applied Physics, 1991, 69[5], 3022-31 

[446-78/79-011] 


The growth behavior and mechanisms of epitaxial lateral overgrowth of GaAs on (001) 

GaAs substrates were investigated. It was found that the lateral growth exhibited a strong 

dependence upon the orientation of the seed. Lateral growth was slowest when the aligned 

275

As GaAs As 

seed was oriented in the [100], [110], [010], [¯110], or equivalent, directions. However, it 

increased sharply when the seed was tilted away from these orientations. Monomolecular 

growth steps on the surface, with a strong contrast and a high lateral spatial resolution, 

were successfully observed by using Nomarski differential interference contrast 

microscopy plus image processing. Their average propagation velocity was estimated to 

be between 0.003 and 0.03mm/s. It was found that the surface of the epitaxial lateral 

overgrowth layer was extremely flat at the atomic scale. The experimental results 

indicated that vertical growth was governed mainly by the propagation of steps that 

originated from their source. On the other hand, lateral growth was limited by the 

diffusion of As in Ga solution until (111)- or (001)-type facets appeared at the lateral 

growth front. The growth then became limited by kinetic processes. It was concluded that 

substrates of (001)GaAs, as well as (111)GaAs, were suitable for obtaining a very smooth 

epitaxial lateral overgrowth layer with a large ratio of the lateral growth width to the 

vertical growth thickness. 

S.Zhang, T.Nishinaga: Journal of Crystal Growth, 1990, 99, 292-6 

[446-76/77-007] 


The chemical reactions and Schottky-barrier characteristics of W(200nm)/Si(0 to 

2.5nm)/GaAs contacts when annealed at 800C were investigated. The Si interfacial layers 

and W films were sputter-deposited onto chemically etched GaAs substrates. The 

W/Si/GaAs diodes clearly exhibited the same Schottky-barrier characteristics as those of 

(WSi 0.6 )/GaAs diodes. By using secondary ion mass spectrometry, the Si layer was found 

to suppress As atom diffusion from GaAs substrates and into W films during annealing 

(800C, 1h). A reduction in natively oxidized GaAs surfaces was also observed in the 

initial stages of Si layer deposition by X-ray photo-emission spectroscopy. These results 

suggested that the Si layer eliminated native oxides from GaAs surfaces, resulting in 

tungsten-silicide/GaAs intimate contact formation at the interface. The Si obstructed the 

diffusion paths of As atoms at W grain boundaries with W-Si-O ternary compounds. 

Y.Kuriyama, S.Ohfuji, J.Nagano: Journal of Applied Physics, 1987, 62[4], 1318-23 

[446-55/56-005] 


The in-diffusion of As vacancies, and their interaction with the mid-gap electron trap, 

EL2, during unprotected and proximity high-temperature annealing was modelled. By 

fitting existing data, it was found that the diffusive capture of V As by EL2 was inhibited 

by a large (greater than 1eV) repulsive barrier of unknown origin. When taken together 

with other published results, the model indicated that the diffusivity of V As was described 

by: 

D(cm 2 /s) = 0.004 exp[-1.8(eV)/kT] 

However, this value was thought to be uncertain by at least an order of magnitude. A new 

discovery was the existence of a strong repulsive barrier between V As and EL2. This 

inhibited their interaction, and was of the order of 1.2eV. When interpreted as being a 

Coulomb barrier, it required that EL2 (or a portion of EL2) should carry a positive 

charge. 

276

As GaAs As 

In order to keep EL2 neutral, there would then have to be a compensating acceptor in 

EL2. 

K.M.Luken, R.A.Morrow: Journal of Applied Physics, 1996, 79[3], 1388-90 

[446-134/135-124] 


Quantitative determinations were made of the contribution which As self-interstitials 

made to the As self-diffusion coefficient. Values of the As self-interstitial contribution 

were deduced from S in-diffusion profiles, which were simulated on the basis of a kickout 

mechanism. 

M.Uematsu, P.Werner, M.Schultz, T.Y.Tan, U.M.Gösele: Applied Physics Letters, 1995, 

67[19], 2863-5 

[446-125/126-120] 


Samples with a 100nm Co over-layer, which had been subjected to rapid thermal 

annealing (400 to 650C, 60s), were analyzed by using mass and energy dispersive recoil 

spectrometry. Separate characterizations of the C, O, Co, Ga, and As depth distributions 

were carried out. It was found that As migrated to the surface at annealing temperatures 

which were higher than 450C. The composition at various depths was determined at a 

number of temperatures. On the basis of Arrhenius plots, the apparent activation energies 

were estimated to be equal to about 0.6eV for phase formation and equal to 1.3eV for 

diffusion. The X-ray diffraction data indicated that CoGa and CoAs were present in all of 

the annealed samples. Scanning electron microscopy showed that the surface was 

reticulated after heat treatment, and that grain growth occurred at higher temperatures. 

M.Hult, H.J.Whitlow, M.Ostling, M.Andersson, Y.Andersson, I.Lindeberg, K.Stähl: 


[446-117/118-165] 


An investigation was made of near-bandedge photoluminescence from semi-insulating 

crystals after they had been annealed in wafer or bulk form. The results, with respect to 

uniformity after annealing, were in agreement with previous data. The 1.360eV emission 

band which was seen in annealed crystals and which had been assumed to imply that a 

V As -related rapid-diffusion process was the mechanism which was responsible for the 

annealing-induced uniformity, was shown to be unconnected with it. The involvement of 

V As in the band was questioned. From literature data, it was estimated that the diffusion 

coefficient of V As (1.2 x 10 -12 at 1050C, 5.8 x 10 -14 at 850C and 6.7 x 10 -16 cm 2 /s at 650C) 

was too low to permit bulk equilibrium and uniformity via vacancy diffusion from the 

surface at the annealing temperatures which were used. It was concluded that local 

rearrangement of defects was a viable mechanism for producing uniformity during postgrowth 

annealing. 

V.Swaminathan, R.Caruso, S.J.Pearton: Journal of Applied Physics, 1988, 63[6], 2164-7 

[446-157/58-273] 

277

As GaAs B 

GaAs[l]: As Diffusion 

The As concentration profiles ahead of a crystal interface, when advancing into a Ga-rich 

solution, were determined (during the electro-epitaxial growth of layers) by using 

computer simulation techniques. The effects of Peltier heating or cooling, and of 

electromigration, during growth were incorporated. The growth velocity in the absence or 

presence of convection, due to the Peltier effect and to electromigration, was calculated 

under various conditions. It was observed that there was a transition, in the movement of 

As atoms towards the crystal interface, from smooth and orderly to turbulent and wavy as 

the intensity of the electric field increased during electro-epitaxial growth. 

R.S.Q.Fareed, R.Dhanasekaran, P.Ramasamy: Journal of Applied Physics, 1994, 75[8], 

3953-8 

[446-117/118-165] 

Au 

GaAs: Au Diffusion 

Rutherford back-scattering spectrometry and X-ray photoelectron spectroscopy were used 

to investigate compositional changes, in thin metal-semiconductor systems, which were 

caused by Ar + and N + ion bombardment or by annealing. The investigation was carried 

out on contacts of Ni-Au-Ge on GaAs, as well as on irradiated and non-irradiated Au- 

GaAs. The structures were bombarded with Ar + and N + ions to doses of between 10 14 and 

3 x 10 16 /cm 2 . The experimental results indicated that the interdiffusion of Au and Ga 

atoms depended upon the bombardment dose. Upon annealing the samples, blocking of 

interdiffusion was observed in Au-GaAs structures (which had been deposited on 50keV 

Ar + ion-irradiated and pre-treated GaAs) at certain radiation-defect concentrations. This 

behavior was attributed to Au-Ga bond formation, and appeared to depend upon Ga 

interstitial atoms. 

L.B.Guoba, A.A.Vitkauskas, J.V.Kameneckas, V.R.Sargünas, A.P.Sakalas: Physica 

Status Solidi A, 1989, 111[2], 507-13 

[446-64/65-162] 

B 

GaAs: B Diffusion 

The introduction of B into As sites, as deduced from the strength of vibrational modes at 

601.7 and 628.3/cm, was studied as a function of the fluence of 2MeV electrons. 

Simultaneous monitoring of the strength of the Is-2p electronic transitions of neutral 

shallow acceptors, and of the neutral 0.078eV acceptor and its singly-ionized 0.203eV 

level, provided accurate data on the position of the Fermi level during irradiation. The 

results were inconsistent with previous models for the location of B on As sites. A model 

278

B GaAs Be 

was proposed which was based upon the onset of enhanced B diffusion when the Fermi 

level lay above 0.078eV. 

W.J.Moore, R.L.Hawkins: Journal of Applied Physics, 1988, 63[12], 5699-702 

[446-72/73-010] 

Be 

310 GaAs: Be Diffusion 

Spatial localisation of Be in d-doped material, within a few lattice constants (less than 

2nm), was achieved at low growth temperatures for Be concentrations of less than 

10 14 /cm 2 ; as revealed by capacitance-voltage profiles and secondary ion mass 

spectroscopy. At high growth temperatures, and at higher Be concentrations, significant 

spreading of the dopants occurred and was explained in terms of Fermi-level pinninginduced 

segregation, repulsive Coulomb interactions of dopants, and diffusion. The 

highest Be concentration which was obtained at low growth temperatures exceeded 2 x 

10 20 /cm 3 , and was limited by repulsive dopant interactions. It was shown that the 

repulsive Coulomb interaction resulted in a correlated non-random dopant distribution. It 

was deduced that the diffusivity of Be (table 10) could be described by: 

D (cm 2 /s) = 0.00002 exp[-1.95(eV)/kT] 

The present values were much lower than those previously reported. 

E.F.Schubert, J.M.Kuo, R.F.Kopf, H.S.Luftman, L.C.Hopkins, N.J.Sauer: Journal of 


[446-157/159-279] 

Table 10 

Diffusivity of Be in GaAs 


1005 4.7 x 10 -13 

945 1.5 x 10 -13 

895 6.0 x 10 -14 

800 1.4 x 10 -14 

700 1.7 x 10 -15 

600 8.6 x 10 -16 

GaAs: Be Diffusion 

It was pointed out that Be was one of the main p-type dopants which were used for the 

fabrication of devices that were based upon GaAs or related III-V materials. The element 

dissolved substitutionally on the group-III sub-lattice, and diffused via a kick-out 

mechanism which involved group-III self-interstitials. Non-equilibrium concentrations of 

these self-interstitials had a marked effect upon the diffusivity of Be. Various situations 

were considered in which non-equilibrium point defects played a role in Be diffusion. 

279

Be GaAs Be 

These included the in-diffusion of such dopants from an external source, the diffusion of 

grown-in dopants, self-interstitial generation by Fermi-level surface pinning, and 

recombination-enhanced Be diffusion during device operation. It was noted that the 

diffusion behavior of C, which was found on the group-V sub-lattice of GaAs, was much 

less sensitive to non-equilibrium point defects. It was therefore used to replace Be as a p- 

type dopant. 

M.Uematsu, K.Wada, U.Gösele: Applied Physics A, 1992, 55[4], 301-12 

[446-93/94-008] 



broadening of d-doped planes of Be. It was found that sharp spikes of Be could be 

obtained for sheet densities which were below 10 13 /cm 2 and for growth temperatures of 

500C or less. At higher temperatures or densities, segregation or concentration-dependent 

rapid diffusion could occur; thus causing significant spreading even during growth. The 

co-deposition of Si and Be markedly reduced this broadening. 



[446-91/92-001] 


The slow positron technique was used to study undoped and Be-doped samples, and 

thereby determine the effect of Be upon the creation and migration of Ga vacancies, V Ga , 

during annealing. It was deduced that a V Ga mono-vacancy which was created in Bedoped 

material resulted in an enhanced Coulombic interaction between an As vacancy, 

V As , and a Be acceptor, Be Ga . In the case of undoped material, the formation of divacancies, 

V Ga -V As , predominated. The migration length of the vacancies was shorter in 

Be-doped material than in undoped material. It was therefore suggested that Ga 

interstitials, I Ga , existed in the Be-diffused layer and interacted with V Ga which were 

introduced from the surface. It was suggested that a kick-out mechanism governed Be 

diffusion in this material. 

J.L.Lee, L.Wei, S.Tanigawa, M.Kawabe: Journal of Applied Physics, 1991, 69[9], 6364-3 

[446-86/87-009] 


Recombination-enhanced impurity diffusion was observed for the first time in Be-doped 

GaAs. It was found that Be diffusion under forward bias was enhanced by a factor of 

about 10 15 at room temperature, and that the activation energy for diffusion decreased 

from 1.8eV for thermal diffusion: 

D(cm 2 /s) = 8.3 x 10 -7 exp[-1.8(eV)/kT] 

to 0.6eV under recombination-enhanced conditions: 

D(cm 2 /s) = 8.7 x 10 -11 exp[-0.59(eV)/kT] 

280


M.Uematsu, K.Wada: Applied Physics Letters, 1991, 58[18], 2015-7 

[446-84/85-013] 


A close relationship between Be surface segregation and diffusion, in molecular beam 

epitaxial GaAs layers which were heavily doped with Be, was analyzed within the 

framework of a thermodynamic approach to segregation effects. Good agreement between 

the theoretical and experimental results suggested that the main cause of extremely fast 

Be in-diffusion in Be-doped GaAs, and the deterioration of its surface morphology and 

luminescence properties, was Be surface segregation. This resulted in near-surface solidphase 

layer enrichment with Be, as compared with bulk Be-doped GaAs. The effect of 

growth parameters (excess As pressure, substrate temperature, growth rate) and dopant 

level upon the likelihood of Be segregation layer formation was considered. 

S.V.Ivanov, P.S.Kopev, N.N.Ledentsov: Journal of Crystal Growth, 1991, 108[3-4], 661- 

9 

[446-81/82-002] 


The effect of the substrate orientation upon Be transport during GaAs molecular beam 

epitaxy was studied by means of secondary ion mass spectrometry. The substrates were 

misoriented from (100) towards (111)A, and epitaxial growth was performed at 630C for 

Be dopant contents of between 5 x 10 19 and 7 x 10 19 /cm 3 . Surface segregation and 

anomalous diffusion similarly depended upon the substrate orientation. In the case of the 

(311)A orientation, Be transport was sharply reduced from its value for the conventional 

(100) orientation. The results were explained qualitatively by considering the effect of 

atomic steps upon the growing surface. 

K.Mochizuki, S.Goto, C.Kusano: Applied Physics Letters, 1991, 58[25], 2939-41 

[446-81/82-009] 


The out-diffusion of implanted Be was found to be identical after capless or capped 

(nitride or oxide) rapid thermal annealing at temperatures of 900 to 1000C. It depended 

upon the Be dose and its proximity to the surface. Out-diffusion was more pronounced 

when the Be implant was shallow (less than 100nm) and/or the Be + dose was high 

(greater than 10 15 /cm 2 ). It was demonstrated that Be out-diffusion was driven by the 

presence of a highly damaged surface layer. Auger results indicated the formation of a 

BeO x compound at the surface of a high-dose (10 16 /cm 2 ) Be-implanted sample that was 

subjected to capless rapid thermal annealing (1000C, 1s). It appeared that BeO x formation 

occurred when the out-diffused Be interacted with native Ga/As oxides during annealing. 

All of the Be which remained in the GaAs, after rapid thermal annealing at temperatures 

above 900C for 2s, was electrically active. 

H.Baratte, D.K.Sadana, J.P.De Souza, P.E.Hallali, R.G.Schad, M.Norcott, F.Cardone: 


[446-78/79-012] 

281



High depth-resolution secondary ion mass spectrometry profiling was used to investigate 

the broadening of d-doped planes of Be in material which had been prepared by using 

molecular beam epitaxial methods. It was confirmed that concentration-dependent 

diffusion was the predominant broadening process for Be at growth temperatures of less 

than 600C. By incorporating Si atoms into the same plane, it was shown that the 

broadening could be completely inhibited. This suggested that the rapid diffusion process 

resulted from mutual repulsion between the Be Ga - ions, and was prevented by the reverse 

field which arose from Si Ga + ions or by the formation of low-mobility Si Ga + -Be Ga 

- 

complexes. The rapid diffusion of Si as Si Ga -Si As pairs was also reduced. The latter was 

attributed to a Fermi-level effect, with compensation by Be tending to reduce the 

probability of Si As formation. The surface segregation of Si was unaffected, whereas that 

of Be was reduced. This indicated that the surface fields which existed during growth 

contributed to the behavior of Be, but not to that of Si. 

J.J.Harris, J.B.Clegg, R.B.Beall, J.Castagné: Semiconductor Science and Technology, 

1990, 5[7], 785-8 

[446-76/77-007] 


The redistribution of Be implants during post-implantation annealing was studied in order 

to evaluate the effect of implantation damage upon the diffusion process. The Be implants 

exhibited only uniform concentration-dependent diffusion, unlike the rapid up-hill 

diffusion which was observed in the peak of Mg implants. This difference was explained 

by invoking a substitutional-interstitial diffusion mechanism and by performing computer 

simulations of damage-generated point defects. In the up-hill diffusion region, the dopants 

diffused from areas of excess interstitial concentration towards areas of excess vacancy 

concentration. A critical point defect concentration was necessary in order to initiate uphill 

diffusion. This behavior could be induced, in the case of Be implants, by coimplanting 

with a heavier element such as Ar. 

H.G.Robinson, M.D.Deal, D.A.Stevenson: Applied Physics Letters, 1990, 56[6], 554-6 

[446-74-009] 


The behavior of implanted Be was studied by annealing samples which had been 

implanted with low or high Be doses. The high-dose (10 14 /cm 2 ) samples exhibited an 

increase in diffusion upon increasing the annealing temperature from 700 to 900C. 

However, the low-dose (2 x 10 13 /cm 2 ) samples exhibited a decrease in diffusivity as the 

temperature increased. The temperature dependence in the low-dose case could be 

reversed by the co-implantation of B (10 14 /cm 2 ). This behavior was explained in terms of 

282


the substitutional-interstitial diffusion mechanism and the relative concentrations of 

interstitial and substitutional Be atoms in the various cases. 

M.D.Deal, H.G.Robinson: Applied Physics Letters, 1989, 55[10], 996-9 

[446-72/73-010] 


Layers of Be-doped material were grown at 300C by using the migration enhanced 

epitaxy technique. The layers exhibited essentially no electrical activation. Rapid thermal 

annealing of the layers resulted in a mobility and hole concentration which were 

comparable to those of conventional molecular beam epitaxial layers which were grown 

at 600C. Secondary ion mass spectroscopy showed that Be diffusion in annealed 

migration enhanced epitaxial layers was much smaller than that in conventional molecular 

beam epitaxial layers; especially highly-doped ones. Raman spectroscopy and 4K 

photoluminescence experiments were also performed. It was concluded that the migration 

enhanced epitaxy method could replace the conventional molecular beam epitaxy method 

in applications which required a high hole concentration and little diffusion. 

B.Tadayon, S.Tadayon, W.J.Schaff, M.G.Spencer, G.L.Harris, P.J.Tasker, C.E.C.Wood, 

L.F.Eastman: Applied Physics Letters, 1989, 55[1], 59-61 

[446-70/71-106] 


Abrupt Be doping profiles were obtained by means of organometallic vapor phase 

epitaxy. Secondary ion mass spectroscopy was used to study the annealing behavior of 

profiles for Be concentrations of 2 x 10 18 /cm 3 . The diffusion fronts were non-Gaussian 

and abrupt. Estimates of the diffusion coefficient of Be were obtained by assuming a 

quadratic concentration dependence. The Be diffusion coefficient was equal to about 10 - 

15 cm 2 /s at 825C. This was at least an order of magnitude lower than that reported for Zn 

profiles which were grown by means of organometallic vapor phase epitaxy. In addition, 

anomalous surface tailing growth was observed. This was very similar to that which was 

reported to occur during Be doping via molecular beam epitaxy. 

M.J.Tejwani, H.Kanber, B.M.Paine, J.M.Whelan: Applied Physics Letters, 1988, 53[24], 

2411-3 

[446-64/65-162] 


A review was presented of self-diffusion mechanisms and doping-enhanced superlattice 

disordering. With regard to the influence of Be p-type dopants, the Fermi level effect had 

to be considered; together with dopant diffusion-induced Ga self-interstitial 

supersaturation or undersaturation. In accord with its effect upon superlattice disordering, 

Be diffusion appeared to be governed by the kick-out mechanism. It was concluded that 

dislocations in this material and in other III-V compounds were only moderately efficient 

sinks or sources for point defects. 

T.Y.Tan, U.Gösele: Materials Science and Engineering, 1988, B1, 47-65 

[446-62/63-208] 

283



The diffusion profiles of buried Be dopant which had been implanted by using a focussed 

ion beam were determined after annealing. The diffusion coefficient of the Be was 

determined by fitting the results of computer calculations. It was found that the diffusion 

coefficient of the Be was enhanced by excess interstitial Be. The Be diffusion profiles 

expanded upon annealing at 850C. The diffusion coefficient of Si which had been 

introduced by using a focussed ion beam was undetectably small when compared with 

that of Be at 850C. 

T.Morita, J.Kobayashi, T.Takamori, A.Takamori, E.Miyauchi, H.Hashimoto: Japanese 


[446-55/56-005] 


Models were presented for the distribution profiles of Be in ion-doped layers after 

implantation and annealing. The possibility of predicting the mean free path of Be in III- 

V compounds was considered. The effect of defect-impurity interactions upon Be 

diffusion was also examined. It was found that a flux of impurities towards the surface 

occurred which was not diffusive in nature. 

G.I.Koltsov, V.V.Makarov, S.J.Yurchuk: Fizika i Tekhnika Poluprovodnikov, 1996, 

30[10], 1907-16 (Semiconductors, 1996, 30[10], 996-1000) 

[446-148/149-171] 


Atomic-resolution images of Be delta-doped layers were obtained by means of crosssectional 

scanning tunnelling microscopy. In the case of samples which had been grown 

at 480C, it was observed that the doping layer width for concentrations of up to 10 13 /cm 2 

was less than 1nm. At higher dopant concentrations, it was found that the dopant layer 

thickness increased markedly with dopant concentration. It was suggested that this 

broadening of the dopant layer at high dopant concentrations was due to Coulombic 

repulsion between individual Be ions. The effect of Coulombic repulsion could also be 

observed in the spatial distribution of the dopant atoms in the plane of the dopant layer. 

P.M.Koenraad, M.B.Johnson, H.W.M.Salemink, W.C.Van der Vleuten, J.H.Wolter: 

Materials Science and Engineering B, 1995, 35[1-3], 485-8 

[446-136/137-109] 


A study was made of the role that was played by the wafer surface in the transient 

diffusion of Be. Samples were doped during molecular beam epitaxial growth, and were 

annealed (900C, 0.25 or 2h) under 2 different caps. In some of the annealed samples, the 

dopant was initially located near to the surface. In other samples, the dopant was initially 

located in a buried layer. Both types of sample were analyzed by means of secondary ion 

mass spectrometry. It was found that variations in the diffusion behavior under the 

various experimental conditions could all be qualitatively explained in terms of a model 

284


which took account of 3 important effects. These were the transient evolution of point 

defect populations, the injection of Ga vacancies by an oxide cap, and the efficiency of 

the surface in restoring point-defect equilibrium. 

Y.M.Haddara, M.D.Deal, J.C.Bravman: Applied Physics Letters, 1996, 68[14], 1939-41 

[446-134/135-124] 


The migration of ion-implanted Be was studied as a function of Al concentration. The 

behavior of Be in AlGaAs was similar to that in GaAs, and it even exhibited the 

anomalous characteristic of increased redistribution with decreasing temperature. The 

results could be described by: 



GaAs: D(cm 2 /s) = 8.9 x 10 -10 exp[-1.00(eV)/kT] 

The diffusivity of Be appeared to increase with Al content. This was suggested to be due 

to an increase in the bond strength of matrix atoms upon adding Al. This prevented the 

easy transfer of Be from interstitial to substitutional sites. An over-saturation of Be 

interstitials could also explain the persistence of anomalous diffusion in AlGaAs with 

respect to the annealing temperature. The results were explained in terms of a 

substitutional-interstitial diffusion mechanism, the relative amounts of interstitial and 

substitutional Be, and the relative difficulty of moving from an interstitial to a 

substitutional site. 


[446-123/124-160] 


The diffusion of implanted Be in liquid-encapsulated Czochralski material was modelled 

by using a computer simulation. The so-called plus-one approach to defect generation 

after implantation, as well as the assumed existence of local Ga interstitial sinks, were 

successfully used to simulate a high Be diffusivity, up-hill diffusion and a time-dependent 

Bi diffusivity. The fast diffusion of implanted Be could be simulated by using the same 

intrinsic Bi diffusivity as that used in simulating the slow diffusion of molecular beam 

epitaxially grown-in Be. Account was taken of the roles which were played by extended 

defects and non-equilibrium Ga point defects in affecting the anomalous diffusion 

behavior of implanted Be. 

J.C.Hu, M.D.Deal, J.D.Plummer: Journal of Applied Physics, 1995, 78[3], 1606-13 

[446-123/124-161] 


Diffusion was studied in samples of molecular beam epitaxial material with grown-in Be. 

The diffusion profiles of samples which had been annealed under various conditions were 

determined by using secondary ion mass spectrometry, and a computer simulation was 

used to analyze the experimental results and extract diffusion parameters. The Be 

285


diffusion profiles exhibited kinks, and a time-dependent diffusivity, which were 

successfully simulated. It was deduced that the intrinsic Be diffusivity was described by: 

D(cm 2 /s) = 0.17 exp[-3.39(eV)/kT] 


[446-123/124-161] 


Heavily-doped polycrystalline material which had been grown by molecular beam epitaxy 

and metalorganic chemical vapor deposition was studied by using secondary ion mass 

spectrometry and Hall measurements. It was found that Be rapidly diffused into the 

undoped buffer layer at a growth temperature of 450C. The concentration-depth profiles 

of Be in AlGaAs/GaAs heterojunction bipolar transistor layers indicated that Be diffused 

mainly along grain boundaries. 

K.Mochizuki, T.Nakamura: Applied Physics Letters, 1994, 65[16], 2066-8 

[446-119/120-191] 


Implantation (1.5 x 10 14 /cm 2 ) of 30keV Be into (x11)A-oriented semi-insulating GaAs 

substrates (where x took values of up to 4) was carried out. For comparison, (110)- and 

(100)-oriented substrates were also implanted. It was found that the in-diffusion of Be in 

(311)A-oriented substrates was lower than that in (100) material. 

M.V.Rao, H.B.Dietrich, P.B.Klein, A.Fathimulla, D.S.Simons, P.H.Chi: Journal of 


[446-117/118-166] 


A study was made of the diffusion of Be in δ-doped layers of (111)A- or (100)-type, and 

of the evaporation of As atoms from the surfaces. It was found that the diffusion of 

dopants in (111)A layers was slower than in (100), regardless of the presence of As 

vacancies. On the other hand, diffusion in (100) layers was enhanced by the presence of 

As vacancies. It was noted that As atoms on the (111)A surface did not evaporate easily, 

as compared with those on the (100) surface. 

A.Shinoda, T.Yamamoto, M.Inai, T.Takebe, T.Watanabe: Japanese Journal of Applied 

Physics, 1993, 32[2-10A], L1374-6 

[446-115/116-116] 


Abnormal out-diffusion from heavily Be-doped material, prepared by molecular beam 

epitaxy, was found to be initiated by a decrease in the lattice constant of the p + epilayer. 

From double-crystal X-ray spectra, Van der Pauw measurements, photoluminescence data, 

and infra-red absorption spectra for Be-doped material with various dopant concentrations, 

it was deduced that there existed a critical doping concentration (2.6 x 10 19 /cm 3 ) beyond 

which the lattice constant of the epilayer began to decrease, and Be out- 

286


diffusion into the substrate was significantly enhanced. It was suggested that the tensile 

stress on the epilayer resulted in abnormal Be out-diffusion. The absorption coefficient, 

in the 8 to 10µ region, of Be-doped material with a carrier concentration of 8.3 x 

10 19 /cm 3 was found to be about 10 4 /cm. 

B.D.Liu, T.H.Shieh, M.Y.Wu, T.C.Chang, S.C.Lee, H.H.Lin: Journal of Applied Physics, 

1992, 72[7], 2767-72 

[446-106/107-033] 


Samples of n-type and p-type d-doped material, grown by molecular beam epitaxy and 

with quite high doses of Be, were investigated by means of transmission electron 

microscopy. The magnitude of the doses ranged from half a monolayer to 2 monolayers. 

The microscopic structures of the d-doped regions and of the adjacent epilayers were 

observed directly. The effect of impurity spreading upon the hetero-interfaces and 

superlattices was studied. The Be atoms which were present in Be d-doped samples 

spread over a wide region and caused rough hetero-interfaces and wavy superlattices to 

form. The spreading of Be was attributed to segregation and diffusion which occurred 

during growth. Stacking faults were found in d-doped samples when they were grown at 

low temperatures. Their presence was attributed to local strains that were caused by 

heavy doping. 

D.G.Liu, J.C.Fan, C.P.Lee, K.H.Chang, D.C.Liou: Journal of Applied Physics, 1993, 

73[2], 608-14 

[446-106/107-034] 


A study was made of recombination-enhanced impurity diffusion in Be-doped material. 

An investigation of tunnel diodes revealed that Be diffusion under forward bias was 

enhanced by a factor of about 10 15 at room temperature, and that the activation energy for 

diffusion was reduced from 1.8eV for thermal diffusion to 0.6eV for recombinationenhanced 

impurity diffusion. In the latter process for Be, the energy which was related to 

minority carrier injection at the recombination center encouraged the annihilation of the 

recombination center, where a point defect which enhanced the Be diffusion was 

generated. 

M.Uematsu, K.Wada: Materials Science Forum, 1992, 83-87, 1551-6 

[446-99/100-063] 

GaAs/AlGaAs: Be Diffusion 

The active layers of single quantum-well separate confinement heterostructure lasers 

which had been grown by means of molecular beam epitaxy were investigated by using 

photoluminescence absorption spectroscopy, secondary ion mass spectroscopy, 

capacitance-voltage profiling and laser threshold current measurements. It was found that 

a significant amount of Be diffusion occurred under normal growth conditions. Large 

concentrations of Be in the quantum well were correlated with the lack of an exciton 

feature in the absorption spectrum. The amount of Be in the active region was reduced by 

287

Be GaAs C 

reducing the Be concentration and by decreasing the growth temperature in the upper 

cladding region of the laser. 

G.E.Kohnke, M.W.Koch, C.E.C.Wood, G.W.Wicks: Applied Physics Letters, 1995, 

66[21], 2786-8 

[446-121/122-062] 

GaAs/GaAlAs: Be Diffusion 

Heterostructures of GaAs/Ga 0.7 Al 0.3 As, which contained Zn and Se as intrinsic p and n 

dopants, were subjected to combined Be and O implantation. Rapid thermal annealing 

then resulted in the redistribution of Be. The Se dopant profile remained essentially 

unchanged. The atomic profile Be could be related to the microscopic defect 

distributions. A change in the photoluminescence spectrum, due to over-compensation of 

the n-doped GaAs and GaAlAs layers, was observed and the corresponding signals which 

were associated with Be were identified. 

T.Humer-Hager, R.Treichler, P.Wurzinger, H.Tews, P.Zwicknagl: Journal of Applied 

Physics, 1989, 66[1], 181-6 

[446-74-026] 

C 

GaAs: C Diffusion 

The effects of background doping, surface encapsulation, and an As 4 over-pressure upon 

C diffusion were studied by annealing samples which had 100nm p-type C doping spikes 

within 0.00lmm layers of undoped (n - ), Se-doped (n + ) and Mg-doped (p + ) material. The 

layers were grown via low-pressure metalorganic chemical vapor deposition, using CCl 4 

as the dopant source. Two different As 4 over-pressure conditions were investigated. 

These were those of an equilibrium P As (no excess As), and of a pressure, P As , of about 

2.5atm. For each As 4 over-pressure condition, both capless and Si 3 N 4 -capped samples of 

the n - , n + and p + crystals were simultaneously annealed (825C, 24h). Secondary-ion mass 

spectroscopy was used to measure atomic C depth profiles. The C diffusion coefficient 

was always low, but depended upon the background doping. It was highest in Mg-doped 

(p + ) material and was lowest in Se-doped (n + ) material. The effect of Si 3 N 4 surface 

encapsulation and P As upon C diffusion was minimal. 

B.T.Cunningham, L.J.Guido, J.E.Baker, J.S.Major, N.Holonyak, G.E.Stillman: Applied 


[446-70/71-106] 


Atomic and carrier concentration profiles in C-implanted material were measured. The 

300keV C-ion implantation was carried out to a dose of 1.0 x 10 14 /cm 2 . The C 

concentration profiles which were revealed by secondary ion mass spectrometric 

measurements were found to be in good agreement with profiles which were predicted by 

288

C GaAs Cd 

Monte Carlo simulations. The implanted C did not diffuse greatly during annealing at 

900C because the diffusion coefficient was less than 4 x 10 -16 cm 2 /s for ion-implanted C. 

Therefore, a shallow carrier concentration profile was found after annealing. The 

activation efficiency was equal to 17% at the surface (with a depth of less than 0.47µ). 

However, the efficiency was as low as 4% in deeper regions. This was attributed to the 

suppression of activation by the precipitation of C after annealing. 

T.Hara, S.Takeda, A.Mochizuki, H.Oikawa, A.Higashisaka, H.Kohzu: Japanese Journal 

of Applied Physics, 1995, 34[2-8B], L1020-3 

[446-125/126-120] 


Heavily-doped polycrystalline material which had been grown by molecular beam epitaxy 

and metalorganic chemical vapor deposition was studied by using secondary ion mass 

spectrometry and Hall measurements. It was found that C diffusion was negligible, even 

during post-growth annealing at 800C. However, annealing increased the resistivity of C- 

doped GaAs, and this was suggested to be due to a change in the occupation site 

preference of C atoms from As sites. 

K.Mochizuki, T.Nakamura: Applied Physics Letters, 1994, 65[16], 2066-8 

[446-119/120-191] 


First-principles estimates were made of the doping efficiency and diffusion mechanism of 

C. The C acceptor which occupied an As site was found to be the most stable, and was 

responsible for a high doping efficiency. However, the hole concentration saturated at 

about 10 20 /cm 3 , due to compensation by donors such as [100] split interstitial (CC) [100] 

complexes. A mechanism was proposed, for C diffusion that was accompanied by the 

formation and dissociation of the (CC) [100] complex, in which the activation energy was 

lower than that for atom diffusion. 

B.H.Cheong, K.J.Chang: Physical Review B, 1994, 49[24], 17436-9 

Cd 

[446-115/116-116] 

GaAs: Cd Diffusion 

Low p-type surface concentrations were introduced at high temperatures by using a Ga- 

Cd alloy as a diffusion source. Concentration profiles were determined by using 

electrochemical profiling techniques. The resultant profiles were of erfc-type. It was 

found that the surface concentration of the carriers was reduced if diffusion was carried 

out by using Ga-Cd alloys which contained less than 1at%Cd. Results which could be 

described by: 

D(cm 2 /s) = 1.10 x 10 -13 exp[-2.12(eV)/kT] 

were found when pure Cd was used as a diffusion source, together with a 5nm SiO 2 overlayer. 

When Ga-1at%Cd alloy was used as a source, at temperatures of 800 and 850C, the 

289

Cd GaAs Co 

diffusivities were 8.2 x 10 -15 and 2.54 x 10 -14 cm 2 /s, respectively; thus suggesting that the 

diffusivity could be described by: 

D(cm 2 /s) = 1.29 x 10 -14 exp[-2.17(eV)/kT] 

The surface concentration could be further reduced to 10 17 /cm 3 if an 0.1at%Cd alloy was 

used. In this case, the diffusion coefficient at 850C was 1.45 x 10 -14 cm 2 /s. 

D.K.Gautam, Y.Nakano, K.Tada: Japanese Journal of Applied Physics, 1991, 30[6], 

1176-80 

[446-84/85-013] 

GaAs: Cd Diffusion 

The removal of damage after heavy implantation with 111m Cd and 111 In was investigated 

by using perturbed angular correlation and Hall techniques. After implantation at 90K, 

and subsequent annealing, the removal of structural disorder in the vicinity of the 111 In 

probe atom was observed at about 300K. The annealing behavior, at temperatures ranging 

from 500 to 1100K, of GaAs which had been implanted with 111m Cd and 111 In was 

investigated as a function of total implantation dose. After annealing at 600K, some of the 

Cd probe atoms were located in a slightly perturbed environment while the remainder 

were in a heavily perturbed one. Annealing at temperatures above 900K led to the outdiffusion 

of Cd which was located in heavily perturbed sites, and electrical activation 

occurred. In contrast to Cd, all of the In probe atoms were located in a slightly perturbed 

environment and no In was lost by out-diffusion. These differences were explained in 

terms of extended defects and their interactions with probe atoms. 

W.Pfeiffer, M.Deicher, R.Kalish, R.Keller, R.Magerle, N.Moriya, P.Pross, H.Skudlik, 

T.Wichert, H.Wolf: Materials Science Forum, 1992, 83-87, 1481-6 

[446-99/100-063] 

Co 

GaAs: Co Diffusion 

Interfacial reactions between thin Co films and monocrystalline GaAs substrates were 

studied by using Auger electron spectroscopic, transmission electron microscopic, and X- 

ray diffraction methods. The interaction began at about 325C, with the formation of a 

ternary phase (probably Co 2 GaAs) which grew in a highly oriented manner with respect 

to the (001) substrate; with a lattice mismatch of about -10%. The reaction kinetics were 

studied and were found to be diffusion-controlled; with an activation energy of 0.7eV. 

The Co was deduced to be the predominant diffusing species. The oriented ternary phase 

coexisted with randomly oriented CoGa and CoAs at temperatures of between 325 and 

500C while, at higher temperatures, only the binary compounds prevailed. 

M.Genut, M.Eizenberg: Applied Physics Letters, 1987, 50[19], 1358-60 

[446-51/52-116] 

290

Cr GaAs Cr 

Cr 

311 GaAs: Cr Diffusion 

The diffusion of Cr S resulted from the rapid migration of Cr i , and their subsequent 

change-over so as to occupy Ga sites (or vice versa); a typical substitutional-interstitial 

diffusion process. It was noted that there were 2 ways in which the Cr i -Cr s change-over 

could occur. One involved a kick-out mechanism in which Ga self-interstitials took part, 

and the other involved a dissociative mechanism in which Ga vacancies took part. It was 

observed that the Cr s in-diffusion profiles had a characteristic shape which revealed the 

predominance of a kick-out mechanism, whereas the Cr s out-diffusion profiles were erfshaped; 

thus reflecting the predominance of the dissociative mechanism. An integrated 

substitutional-interstitial diffusion mechanism, which took account of kick-out and 

dissociative mechanisms, was here used to analyze Cr diffusion results. It was confirmed 

that the kick-out mechanism governed Cr in-diffusion, while the dissociative mechanism 

governed Cr out-diffusion. The parameters which were used to fit existing experimental 

results provided quantitative information on the Ga self-interstitial contribution to the Ga 

self-diffusion coefficient. The values which were obtained (table 11) were consistent with 

those of a study of Zn diffusion in GaAs, and with available experimental data on Al-Ga 

interdiffusion coefficients. 


[446-93/94-008] 

Table 11 

Diffusivity of Cr in GaAs 


990 8.1 x 10 -18 

890 2.6 x 10 -18 

790 1.1 x 10 -19 

GaAs: Cr Diffusion 

It was pointed out that one of the potential advantages of rapid thermal annealing, as 

compared with conventional furnace annealing, was a reduced implanted dopant and 

background impurity diffusion. Here, the migration of Cr during the annealing of Crdoped 

semi-insulating material implanted with 100keV Si + ions to a dose of 7 x 10 12 /cm 2 

was measured using secondary ion mass spectrometry. Un-capped rapid thermal 

annealing (860 or 930C, 1 to 60s) was investigated and its effect was compared with that 

of capless furnace annealing (0.5h). It was deduced that, during rapid thermal annealing, 

Cr migration was marked and exhibited a strong time-temperature dependence. 

H.Kanber, J.M.Whelan: Journal of the Electrochemical Society, 1987, 134[10], 2596-9 

[446-55/56-005] 

291

Cr GaAs Cu 

GaAs: Cr Diffusion 

The migration of Cr was studied by using a new photoluminescence method which 

involved measurement of the Cr-related luminescence intensity. The form of the intensity 

profiles after thermal annealing was explained in terms of the substitutional-interstitial 

dissociative mechanism. It was observed that the redistributed profiles had an abnormal 

peak, in the near-surface region of wafers, which was annealed out by heat treatment at 

temperatures above 900C for 6h. This was attributed to the in-diffusion of As vacancies 

from the surface. The Cr-related luminescence was studied as a function of the As 

pressure during wafer annealing. The data showed that the in-depth profiles could be 

understood in terms of the diffusion of Cr and As vacancies. The luminescence center 

was deduced to be a Cr Ga -V As complex. 

J.T.Hsu, T.Nishino, Y.Hamakawa: Japanese Journal of Applied Physics, 1987, 26[5], 

685-9 

[446-61-066] 

Cu 

GaAs: Cu Diffusion 

The diffusion of Cu in Si-doped material was studied. Photo-induced current transient 

spectroscopic techniques were used to identify deep levels in Si-doped Cu-compensated 

material. The effect of capping the deposited Cu layer during the diffusion of Cu in Sidoped 

material was studied by obtaining photo-induced current transient spectra at 

various depths from the sample surface. A concentration gradient of the energy levels of 

Cu-associated complexes was found to exist into the depth of the diffused sample. This 

was to be expected, because the formation of the V As Cu Ga V As complex, which was 

considered to be responsible for the Cu B level, was favored by an increase in the 

concentration of V As . 

L.M.Thomas, V.K.Lakdawala: Defect and Diffusion Forum, 1993, 95-98, 931-6 

[446-95/98-931] 


Experiments were performed on polished plates of Te-doped material. Irradiation with 15 

to 150keV protons was carried out at 300K to doses of between 10 16 and 10 17 /cm 2 . It was 

found that the impurity profiles did not depend upon whether the diffusion source was 

deposited before or after irradiation. The penetration depth in samples which were 

irradiated with 15keV protons was greater than that in samples which were irradiated with 

150keV protons. It was suggested that this was because the low-energy ions generated 

more defects at depths of between 50 and 100nm. 

V.N.Abrosimova, V.V.Kozlovskii, N.N.Korobkov, V.N.Lomasov: Izvestiya Akademii 

Nauk SSSR - Neorganicheskie Materialy, 1990, 26[3], 488-91. (Inorganic Materials, 

1990, 26[3], 411-4) 

[446-84/85-013] 

292

Cu GaAs D 


It was recalled that a Cu-related peak at about 1.35eV was generally observed in the lowtemperature 

photoluminescence spectra of epitaxial layers. Samples were treated in Cusaturated 

aqueous KOH solutions at room temperature, and it was shown that Cu could be 

deposited onto the semiconductor surface when a source of metallic Cu was present in the 

KOH solution. The results suggested that Cu could diffuse into the semiconductor, even 

at room temperature. 

K.Somogyi, D.N.Korbutyak, L.N.Lashkevich, I.Pozsgai: Physica Status Solidi A, 1989, 

114[2], 635-42 

[446-74-009] 


The diffusion of Cu in semi-insulating liquid-encapsulated Czochralski-type material at 

800C was studied by using photoluminescence, photo-etching, secondary ion mass 

spectroscopic, and temperature-dependent Hall techniques. The results indicated a 

diffusivity of 4.5 x 10 -6 cm 2 /s. It was deduced that diffusion occurred mainly by 

substitution on lattice sites. The Cu migrated preferentially along the walls of the 

dislocation cells. 

S.Griehl, M.Herms, J.Klöber, J.R.Niklas, W.Siegel: Applied Physics Letters, 1996, 

69[12], 1767-9 

[446-138/139-076] 


The diffusion of Cu into undoped material reduced the concentration of the EL6 and EL2 

deep-donors and created a deep-donor level at about 0.66eV, in addition to the wellknown 

acceptor levels at 0.15 and 0.44eV; as revealed by deep-level transient 

spectroscopic and temperature-dependent Hall measurements. An analysis which was 

based upon these observations, and a charge-balance equation, was carried out in order to 

understand the compensation process at various stages. The possible identity of the deepdonor 

which was responsible for the semi-insulating properties was considered, as was 

the anomalous transport behavior in the highly-compensated samples. 

B.H.Yang, H.P.Gislason: Materials Science Forum, 1995, 196-201, 713-8 

[446-127/128-118] 

D 

GaAs: D Diffusion 

The diffusion of D in Si-doped GaAs was studied. It was found that the diffusion profile 

could be closely fitted by using an erfc function. It was suggested that, in Si-doped 

samples, the D behaved like a deep acceptor with a level, H -/0 , which was slightly 

resonant in the conduction band. 

293

D GaAs D 

J.Chevallier, B.Machayekhi, C.M.Grattepain, R.Rahbi, B.Theys: Physical Review B, 

1992, 45[15], 8803-6 

[446-86/87-001] 


The diffusion depth and total amount of D which was incorporated during exposure to a 

plasma was found to depend markedly upon the conductivity type of the surface. Thus, a 

shallow n + layer inhibited D in-diffusion in a manner which was consistent with the 

suggestion that the species had a level in the upper half of the GaAs band-gap. The deactivation 

of donors and acceptors by D was then the result of various chemical reactions 

which were based upon its differing charge state in n-type and p-type material. Thus, Zn 

acceptors exhibited a re-activation energy of 1.6eV. This was less than the typical donor 

value (2.1eV). 

S.J.Pearton, W.C.Dautremont-Smith, J.Lopata, C.W.Tu, C.R.Abernathy: Physical Review 

B, 1987, 36[8], 4260-4 

[446-55/56-006] 


The dynamics of D in Zn-doped material were investigated by using anelastic relaxation 

techniques. It was suggested that the most likely configuration was for D to be trapped by 

substitutional Zn, although it was also possible that D was trapped at a Ga vacancy. 

Relaxation of D occurred at about 20K in the kHz range, and had the highest rate which 

had yet been found for a H isotope in a semiconductor. The shape of the curves of elastic 

energy loss versus temperature indicated that the D reorientation was strongly quantized. 

G.Cannelli, R.Cantelli, F.Cordero, E.Giovine, F.Trequattrini, M.Capizzi, A.Frova: Solid 

State Communications, 1996, 98[10], 873-7 

[446-136/137-109] 


The D diffusion profiles in material which was doped with various group-II (Mg, Zn, Cd) 

or group-IV (C, Ge) acceptors had similar characteristics; even though the neutralization 

of acceptors at 300K was not always efficient. Conductivity and Hall-effect 

measurements were used to study the electrical characteristics of hydrogenated p-type 

epilayers. The temperature dependences of the free carrier concentration and hole 

mobility, before and after hydrogenation, showed that the neutralization of acceptors by 

atomic H led to the elimination of the shallow acceptor states. Infra-red absorption lines 

that were associated with H-acceptor complexes were observed for all of the acceptors, 

except Mg. It was established that the microscopic structure of H-acceptor complexes 

depended upon the acceptor site in the lattice. 

R.Rahbi, B.Pajot, J.Chevallier, A.Marbeuf, R.C.Logan, M.Gavand: Journal of Applied 

Physics, 1993, 73[4], 1723-31 

[446-106/107-034] 

294

D GaAs Fe 


It was recalled that H diffusion in III-V semiconductors usually led to a reduction in the 

active dopant concentration, and to an increase in the free carrier mobility. It was 

considered that this neutralization of the dopants was a result of the formation of a 

complex which included H and the dopant atom. The microscopic structure was deduced 

from a detailed analysis of the infra-red local vibrational modes of H and the dopants. 

Modelling of D diffusion profiles Zn-doped material indicated the presence of a H donor 

level at 1.1eV below the conduction band minimum. The D diffusion behavior in Sidoped 

GaAs or Si-doped Al x Ga 1-x As indicated that the acceptor level of H was slightly 

resonant in the GaAs conduction band, and became localized in the band gap of alloys 

when x was greater than 0.08. 

J.Chevallier, B.Pajot: Materials Science Forum, 1992, 83-87, 539-50 

Fe 

[446-99/100-064] 

312 GaAs: Fe Diffusion 

The Fe was diffused from a spun-on glass film, and onto n-type wafers, at temperatures 

of between 700 and 900C (table 12). The diffusivities, as determined by using junctiondepth 

and conductivity techniques, could be explained in terms of a model which 

assumed the existence of exhaustible diffusion sources. It was found that the diffusivity 

was described by: 

D(cm 2 /s) = 1000 exp[-2.7(eV)/kT] 

within the above temperature range. 

J.Ohsawa, H.Kakinoki, H.Ikeda, M.Migitaka: Journal of the Electrochemical Society, 

1990, 137[8], 2608-11 

[446-76/77-008] 

Table 12 

Diffusivity of Fe in GaAs 


905 2.2 x 10 -9 

805 1.4 x 10 -10 

705 1.2 x 10 -11 

GaAs: Fe Diffusion 

Secondary ion mass spectroscopic results revealed that the use of spun-on film diffusion 

could produce very flat Fe profiles whose concentrations of 10 15 to 10 17 /cm 3 were 

consistent with the solubility of Fe at diffusion temperatures of 650 to 900C. The surface 

accumulation region was far smaller than that which was produced by the use of 

295

Fe GaAs Ga 

conventional techniques. It was suggested that there was a relationship between the 

solubility and the equilibrium concentration of Ga antisite defects. 

J.Ohsawa, M.Nakamura, Y.Nekado, M.Migitaka, N.Tsuchida: Japanese Journal of 

Applied Physics, 1995, 34[2-5B], L600-2 

[446-123/124-162] 

1.0E-09 

1.0E-10 

1.0E-11 

1.0E-12 

D (cm 2 /s) 

1.0E-13 

1.0E-14 

1.0E-15 

1.0E-16 

1.0E-17 

table 13 

table 14 

table 15 

1.0E-18 

1.0E-19 

7 8 9 10 11 12 

10 4 /T(K) 

Figure 4: Diffusivity of Ga in GaAs 

Ga 

GaAs: Fe Diffusion 

The structural properties of samples which had been implanted with 150 or 400keV Fe, to 

doses of between 10 12 and 10 15 /cm 2 , were studied. The depth distributions of the implants 

were compared before and after annealing with, or without, a Si 3 N 4 cap. Rutherford backscattering, 

X-ray double-crystal diffractometry, and secondary ion mass spectroscopy 

results indicated that Fe was markedly redistributed in all of the materials during 

annealing. On the other hand, Ti did not redistribute at all. The driving force for the 

redistribution of Fe was thought to be not classical diffusion, but reaction with 

implantation-induced defects and stoichiometric imbalances. The defect chemistry of as- 

296

Ga GaAs Ga 

implanted arsenides was found to be fundamentally different to that of as-implanted 

phosphides since, in the latter case, the mass ratio of the constituents was much greater 

and the specific energy for amorphization was much lower. 

H.Ullrich, A.Knecht, D.Bimberg, H.Kräutle, W.Schlaak: Journal of Applied Physics, 

1992, 72[8], 3514-21 

[446-106/107-034] 

GaAs: Ga Diffusion 

A review was presented of progress in the understanding of the mechanisms of Ga selfdiffusion 

and impurity diffusion in GaAs, and of the disordering of GaAs/AlGaAs 

superlattices. The self-diffusion of Ga, and Al-Ga interdiffusion, under intrinsic and n- 

doping conditions were governed by triply negatively charged group-III sub-lattice 

vacancies (V Ga 3- ) while, under heavy p-doping conditions, they were probably governed 

by the doubly positively charged self-interstitial, I Ga 2+ . The GaAs/AlGaAs superlattice 

disordering enhancement which was observed under n-doping by Si or Te was attributed 

to the Fermi-level effect, which increased the V Ga 3- concentration. An elusive disordering 

enhancement under p-doping by Zn or Be was attributed to the combined effects of the 

Fermi level, which increased the I Ga 2+ concentration, and to dopant in-diffusion or outdiffusion 

induced I Ga 2+ supersaturation or undersaturation, respectively. In parallel with 

the Ga self-diffusion mechanism in GaAs, diffusion of the Si donor atoms which 

occupied Ga sites was also governed mainly by V Ga 3- . Meanwhile, Si acceptor atoms 

which occupied As sites (a small fraction of the total) diffused via a negatively charged 

As sub-lattice point-defect species. The interstitial-substitutional p-type dopants, Zn and 

Be, diffused via the kick-out mechanism. Their diffusion induced I Ga 2+ supersaturation 

and undersaturation, respectively, under in-diffusion and out-diffusion conditions. 

T.Y.Tan, U.Gösele, S.Yu: Critical Reviews in Solid State and Materials Science, 1991, 

17[1], 47-106 

[446-157/58-293] 


Self-diffusion was investigated by applying secondary ion mass spectroscopic techniques 

to heterostructures which consisted of molecular beam epitaxial layers that contained one 

or more stable isotopes of host crystal elements. Intermixing of the stable isotopes 

between epilayers of various isotopic compositions constituted near-ideal self-diffusion 

conditions; free from the complications of impurities, strain, electric fields and surfaces. 

When diffusion was investigated by using 69 GaAs/ 71 GaAs isotope heterostructures, the 

Ga self-diffusion coefficient in intrinsic GaAs under As-rich ambient could be described 

by: 

D (cm 2 /s) = 43 exp[-4.24(eV)/kT] 

over 6 orders of magnitude at temperatures between 800 and 1225C. This suggested that a 

single defect mechanism controlled the process. 

E.E.Haller, L.Wang: Defect and Diffusion Forum, 1997, 143-147, 1067-78 

[446-143/147-1067] 

297



The diffusion of Cd into GaAs single crystals was investigated at temperatures ranging 

from 804 to 1201C. The penetration profiles which were measured by using secondary 

ion mass spectroscopy and spreading-resistance profiling techniques were modelled 

numerically on the basis of the kick-out diffusion mechanism. This permitted estimates to 

be made of the Ga self-diffusivity, as mediated by doubly positively charged Ga selfinterstitials, 

I Ga 2+ . The Ga self-diffusivities which were found for As-rich and As-poor 

ambients were consistent. Under an As vapor pressure of 1atm, and under electronically 

intrinsic conditions, the I Ga 2+ -mediated Ga self-diffusion data could be described by: 

D (cm 2 /s) = 3.5 x 10 4 exp[-5.74(eV)/kT] 

G.Bösker, N.A.Stolwijk, U.Södervall, W.Jäger: Defect and Diffusion Forum, 1997, 143- 

147, 1109-16 

[446-143/147-1109] 


The effect of triply negatively charged Ga vacancies (V Ga 3- ) and doubly positively 

charged Ga self-interstitials (I Ga 2+ ) upon the self-diffusivity of Ga was studied. Under 

thermal equilibrium and intrinsic conditions, the contribution of V Ga 3- was characterized 

by an activation enthalpy of 6eV in As-rich crystals and of 7.52eV in Ga-rich crystals. 

The contribution of I Ga 2+ was characterized by an activation enthalpy of 4.89eV in the 

case of As-rich crystals, and of 3.37eV in the case of Ga-rich crystals. 

T.Y.Tan, S.Yu, U.Gösele: Journal of Applied Physics, 1991, 70[9], 4823-6 

[446-91/92-007] 


A review of previous data indicated that the self-diffusion of Ga involved either the 

vacancy or the interstitialcy mechanism. A switch from one to the other could occur, due 

to variations in the Fermi level or the As pressure. The self-diffusion appeared to be 

associated with Ga vacancies with a charge of -3, in the case of n-type material, or with 

Ga interstitials with a charge of 2 in the case of p-type material. In the case of intrinsic 

material, a commonly observed V-form of the Ga diffusivity versus As pressure plot was 

suggested to be consistent with the 2 diffusion mechanisms. It was noted that it was a 

simple matter to predict the near-equilibrium Ga diffusivity on the basis of a single 

intrinsic Ga diffusion coefficient. A model was developed which linked the changeover in 

Ga diffusivity and electron concentration at high donor concentrations. 

R.M.Cohen: Journal of Electronic Materials, 1991, 20[6], 425-30 

[446-88/89-014] 


Calculations were made of the absolute formation energies of native defects in this 

compound. It was found that the formation energy, and thus the equilibrium concentration 

of the defects, depended strongly upon the atomic chemical potentials of As and Ga, as 

298


well as upon the electron chemical potential. Hence, the Ga vacancy concentration 

changed by more than 10 orders of magnitude as the chemical potentials of As and Ga 

varied between the thermodynamically allowed limits. It was deduced that the rate of selfdiffusion 

depended markedly upon the surface annealing conditions. 

S.B.Zhang, J.E.Northrup: Physical Review Letters, 1991, 67[17], 2339-42 

[446-88/89-014] 


The effects of rapid thermal processing, upon material with various thicknesses of SiO 2 

encapsulant, were studied by using capacitance-voltage, secondary ion mass 

spectroscopic, and X-ray photo-electron spectroscopic methods. The processing was 

carried out at 760 or 910C, for a period of 9s. The results indicated that a decrease in 

carrier concentration was related to Ga out-diffusion through the SiO 2 . The decrease in 

carrier concentration was attributed to the formation of V Ga -Si Ga complex defects (selfactivated 

centers). At 760C, the amount of Ga out-diffusion was larger in samples with a 

thick SiO 2 coating. At 910C, the amount of Ga out-diffusion was larger in samples with a 

thin SiO 2 coating. This behavior was explained by assuming the operation of 2 different 

types of driving force. These were interfacial thermal stresses, and an interfacial reaction 

between GaAs and SiO 2 . It was noted that interfacial thermal stresses enhanced Ga outdiffusion 

at 760C, whereas interfacial reactions enhanced such out-diffusion at 910C. 

M.Katayama, Y.Tokuda, Y.Inoue, A.Usami, T.Wada: Journal of Applied Physics, 1991, 

69[6], 3541-6 

[446-86/87-010] 


The absolute formation energies of native defects were calculated. It was found that the 

formation energies, and thus the equilibrium concentrations of the defects, depended 

strongly upon the atomic chemical potentials of As and Ga as well as upon the electron 

chemical potential. Thus, the Ga vacancy concentration changed by more than 10 orders 

of magnitude as the chemical potentials of As and Ga varied over the thermodynamically 

possible range. It was concluded that the rate of self-diffusion depended markedly upon 

the surface annealing conditions, and that impurity-enhanced self-diffusion involved 

V Ga 3- under As-rich n-type conditions, or Ga i 3 under Ga-rich p-type conditions. 

S.B.Zhang, J.E.Northrup: Physical Review Letters, 1991, 67[17], 2339-42 

[446-84/85-014] 


Atomistic thermodynamic calculations were made of the energetics of self-diffusion. An 

assessment of the activation enthalpy of the saddle-point configuration of various modes 

of vacancy self-diffusion indicated that second-nearest neighbor hopping was the most 

energetically favorable mechanism, if vacancies were available in equilibrium 

concentrations. An assessment of the activation entropy indicated that normal diffusion 

pre-exponential factors, of the order of 10 -5 to 0.1cm 2 /s, were consistent with vacancy 

299


self-diffusion via second-nearest neighbor hopping. It was proposed that self-diffusion 

which was characterized by pre-exponential factors of the order of 10 7 to 10 8 cm 2 /s, and 

activation energies of the order of 6eV, involved processes in which surface vacancy 

generation was inhibited and self-diffusion was mediated by Frenkel pair generation. 

J.F.Wager: Journal of Applied Physics, 1991, 69[5], 3022-31 

[446-78/79-011] 


Rutherford back-scattering spectrometry and X-ray photoelectron spectroscopy were used 

to investigate compositional changes, in thin metal-semiconductor systems, which were 

caused by Ar + and N + ion bombardment or by annealing. The investigation was carried 

out on contacts of Ni-Au-Ge on GaAs, as well as on irradiated and non-irradiated Au- 

GaAs. The structures were bombarded with Ar + and N + ions to doses of between 10 14 and 

3 x 10 16 /cm 2 . The experimental results indicated that the interdiffusion of Au and Ga 

atoms depended upon the bombardment dose. Upon annealing the samples, blocking of 

interdiffusion was observed in Au-GaAs structures (which had been deposited on 50keV 

Ar + ion-irradiated and pre-treated GaAs) at certain radiation-defect concentrations. This 

behavior was attributed to Au-Ga bond formation, and appeared to depend upon Ga 

interstitial atoms. 

L.B.Guoba, A.A.Vitkauskas, J.V.Kameneckas, V.R.Sargünas, A.P.Sakalas: Physica 

Status Solidi A, 1989, 111[2], 507-13 

[446-64/65-162] 



disordering. It was concluded that Ga diffusion involved triply negatively charged Ga 

vacancies, under intrinsic and n-doped conditions. With regard to the influence of p-type 

dopants, the Fermi level effect had to be considered; together with dopant diffusioninduced 

Ga self-interstitial supersaturation or undersaturation. The self-diffusion of Ga in 

heavily p-doped material was governed by positively charged Ga self-interstitials. It was 

concluded that dislocations in this material and in other III-V compounds were only 

moderately efficient sinks or sources for point defects. 


[446-62/63-208] 


New interdiffusion data on GaAs/AlAs superlattices led to the conclusion that Ga selfdiffusion 

in GaAs involved triply negatively charged Ga vacancies under intrinsic and n- 

doped conditions. The mechanism of Si-enhanced superlattice disordering was the Fermilevel 

effect, which increased the concentrations of the charged point defect species. With 

respect to the effect of Be and Zn p-type dopants, the Fermi level effect had to be 

considered together with dopant diffusion-induced Ga self-interstitial supersaturation or 

300


undersaturation. The self-diffusion of Ga under heavy p-doping conditions was governed 

by positively charged Ga self-interstitials. 

T.Y.Tan, U.Gösele: Applied Physics Letters, 1988, 52[15], 1240-2 

[446-62/63-209] 


The chemical reactions and Schottky-barrier characteristics of W(200nm)/Si(0 to 

2.5nm)/GaAs contacts when annealed at 800C were investigated. The Si interfacial layers 

and W films were sputter-deposited onto chemically etched GaAs substrates. The 

W/Si/GaAs diodes clearly exhibited the same Schottky-barrier characteristics as those of 

(WSi 0.6 )/GaAs diodes. By using secondary ion mass spectrometry, the Si layer was found 

to suppress Ga atom diffusion from GaAs substrates into W films during annealing 

(800C, 1h). A reduction in natively oxidized GaAs surfaces was also observed in the 

initial stages of Si layer deposition by X-ray photo-emission spectroscopy. These results 

suggested that the Si layer eliminated native oxides from GaAs surfaces, resulting in 

tungsten-silicide/GaAs intimate contact formation at the interface. The Si obstructed the 

diffusion paths of Ga atoms at W grain boundaries with W-Si-O ternary compounds. 

Y.Kuriyama, S.Ohfuji, J.Nagano: Journal of Applied Physics, 1987, 62[4], 1318-23 

[446-55/56-005] 


First-principles molecular dynamics simulations were used to investigate the predominant 

migration mechanism of the Ga vacancy, as well as its free energy of formation and the 

rate constant for Ga self-diffusion. The results suggested that the vacancy migrated via 

second-nearest neighbor hopping. The calculated diffusion constant was in good 

agreement with the experimental value which was deduced by using 69 GaAs/ 71 GaAs 

heterostructures. However, the predictions differed considerably from the results which 

had been obtained by performing interdiffusion experiments on GaAlAs/GaAs 

heterostructures. 

M.Bockstedte, M.Scheffler: Zeitschrift für Physikalische Chemie, 1997, 200[1-2], 195- 

207 

[446-157/159-301] 

313 GaAs: Ga Diffusion 

The diffusion of implanted Zn was studied, at annealing temperatures of between 625 and 

850C, by means of secondary ion mass spectrometry. A substitutional-interstitial 

diffusion mechanism was proposed in order to explain how deviations of the local Ga 

interstitial concentration, from its equilibrium value, regulated Zn diffusion. It was found 

that it was possible to simulate both box-shaped profiles, that resulted from hightemperature 

annealing, and kink-and-tail profiles which resulted from lower-temperature 

annealing. The simulation data permitted the determination of Arrhenius relationships. 

The equilibrium Ga interstitial concentration was described by: 

C (/cm 3 ) = 7.98 x 10 30 exp[-3.47(eV)/kT] 

while the Ga interstitial diffusion coefficient (table 13) was described by: 

D (cm 2 /s) = 0.4384 exp[-2.14(eV)/kT] 

301


M.P.Chase, M.D.Deal, J.D.Plummer: Journal of Applied Physics, 1997, 81[4], 1670-6 

[446-148/149-172] 

Table 13 

Diffusivity of Ga Self-Interstitials in GaAs 


850 9.8 x 10 -11 

800 3.9 x 10 -11 

750 2.2 x 10 -11 

675 1.8 x 10 -12 


Scanning electron microscopy was used to observe the √19 x √19 and (1 x 1) HT 

reconstructions, and their transitions, on (111)B vicinal surfaces under an As pressure. 

These reconstructions were observed in dark and bright contrast, respectively. During the 

transition, √19 x √19 domains began to develop from macro-step edges onto lower (1 x 

1) HT reconstructed terraces, while (1 x 1) HT domains began to develop from the macrostep 

edges onto the upper √19 x √19 reconstructed terraces. The transition diagram for the 

surface coverage of domains exhibited hysteresis. Since Ga diffusion, As incorporation or 

re-evaporation were enhanced during the transitions, marked step-bunching with rough 

macro-step edges was observed. 

H.W.Ren, M.Tanaka, T.Nishinaga: Applied Physics Letters, 1996, 69[4], 565-7 

[446-136/137-109] 


Isotopically controlled heterostructures were used to study Ga self-diffusion by using 

secondary-ion mass spectrometry. This approach produced a near-ideal random walk 

situation that was free of perturbations arising from electric fields, mechanical stresses, or 

chemical potentials. It was found that the Ga self-diffusion coefficient in intrinsic 

material (table 14) could be described by: 

D(cm 2 /s) = 43 exp[-4.24(eV)/kT] 

over 6 orders of magnitude, at temperatures of between 800 and 1225C, under As-rich 

conditions. No significant doping effects were observed in samples with substrates that 

were doped with Te up to 4 x 10 17 /cm 3 or with Zn up to 10 19 /cm 3 . 

L.Wang, L.Hsu, E.E.Haller, J.W.Erickson, A.Fischer, K.Eberl, M.Cardona: Physical 

Review Letters, 1996, 76[13], 2342-5 

[446-134/135-125] 


Impurity and self-diffusion mechanisms, and the nature of the associated point defects, 

were considered with regard to the Ga sub-lattice. Analyses of doping-enhanced 

AlAs/GaAs superlattice disordering data and impurity diffusion data led to the conclusion 

302


that, under thermal equilibrium and intrinsic conditions, the triply negatively-charged Ga 

vacancy (V 3- Ga ) governed Ga self-diffusion and Al-Ga interdiffusion in As-rich crystals, 

while the doubly positively-charged Ga self-interstitial (I 2+ Ga ) predominated in Ga-rich 

3- 

2+ 

crystals. With sufficient doping, V Ga predominated in n-type crystals, while I Ga 

predominated in p-type crystals; regardless of the composition. The V 3- Ga species also 

contributed to the diffusion of the main donor species, Si, while I 2+ Ga also governed the 

diffusion of the main acceptor species, Zn and Be, via the kick-out mechanism. The 

3- 

thermal equilibrium concentration of V Ga was found to exhibit a temperature 

independence, or even a small negative temperature dependence. That is, when the 

temperature was lowered, the equilibrium concentration of V 3- Ga was either unchanged or 

slightly increased. This behavior was consistent with many experimental results. 

T.Y.Tan: Materials Chemistry and Physics, 1995, 40[4], 245-52 

Table 14 

Diffusivity of Ga in GaAs 


1230 3.3 x 10 -13 

1125 2.4 x 10 -14 

1095 1.1 x 10 -14 

1075 6.5 x 10 -15 

1025 2.0 x 10 -15 

975 2.6 x 10 -16 

950 2.4 x 10 -16 

925 3.8 x 10 -17 

900 2.8 x 10 -17 

845 5.3 x 10 -18 

795 6.9 x 10 -19 

[446-125/126-121] 


Diffusion was studied in samples of molecular beam epitaxial material with grown-in Be. 

The diffusion profiles of samples which had been annealed under various conditions were 

determined by using secondary ion mass spectrometry, and a computer simulation was 

used to analyze the experimental results and extract diffusion parameters. It was deduced 

that the Ga interstitial diffusivity was described by: 

D(cm 2 /s) = 6.4 x 10 -5 exp[-1.28(eV)/kT] 

while the equilibrium concentration of Ga interstitials was described by: 

C(/cm 3 ) = 4.7 x 10 28 exp[-3.25(eV)/kT] 


[446-123/124-161] 

303



Samples with a 100nm Co over-layer, which had been subjected to rapid thermal 

annealing (400 to 650C, 60s), were analyzed by using mass and energy dispersive recoil 

spectrometry. Separate characterizations of the C, O, Co, Ga, and As depth distributions 

were carried out. It was found that Ga migrated to the surface at annealing temperatures 

which were higher than 450C. In samples which were annealed at 650C, clear enrichment 

of Ga within the outer 35nm was observed. The composition at various depths was 

determined at a number of temperatures. On the basis of Arrhenius plots, the apparent 

activation energies were estimated to be equal to about 0.6eV for phase formation and 

equal to 1.3eV for diffusion. The X-ray diffraction data indicated that CoGa and CoAs 

were present in all of the annealed samples. Scanning electron microscopy showed that 

the surface was reticulated after heat treatment, and that grain growth occurred at higher 

temperatures. 

M.Hult, H.J.Whitlow, M.Ostling, M.Andersson, Y.Andersson, I.Lindeberg, K.Stähl: 


[446-117/118-165] 


It was recalled that Ga vacancies were believed to mediate Ga self-diffusion, and the 

diffusion of substitutional impurities which resided on the Ga sub-lattice. First-principles 

calculations were presented for the vacancy-mediated diffusion of Ga. It was shown that a 

DX-like mechanism facilitated the migration of lattice-site atoms into the interstitial 

region, and that the dangling bonds of a second-nearest neighbor vacancy assisted 

migration through the interstitial region. Due to these 2 mechanisms, vacancy-assisted 

diffusion of Ga occurred with a low-energy barrier. 

J.Dabrowski, J.E.Northrup: Physical Review B, 1994, 49[20], 14286-9 

[446-115/116-117] 


The thermal equilibrium concentrations of the various negatively charged Ga vacancy 

species were calculated. The triply negatively charged Ga vacancy, V Ga 3- , was studied in 

particular because it dominated Ga self-diffusion and Ga/Al interdiffusion under intrinsic 

and n-doping conditions, as well as the diffusion of Si donor atoms which occupied Ga 

sites. Under strong n-doping conditions, the thermal equilibrium V Ga 3- concentration was 

found to exhibit a temperature independence or a negative temperature dependence. That 

is, the concentration was unchanged, or increased, as the temperature was decreased. This 

was contrary to the normal point defect behavior, in which the point defect thermal 

equilibrium concentration decreased as the temperature was lowered. The observed 

behavior explained a number of known experimental results, and required either that V Ga 

3- 

should attain its thermal equilibrium concentration at the onset of each experiment, or 

304


required the operation of mechanisms which involved point defect non-equilibrium 

phenomena. 

T.Y.Tan, H.M.You, U.M.Gösele: Applied Physics A, 1993, 56[3], 249-58 

[446-111/112-050] 


Undoped 69 GaAs/ 71 GaAs isotope superlattice structures were molecular beam epitaxially 

deposited onto n-type GaAs substrates which had been Si-doped to about 3 x 10 18 /cm 3 . 

They were then used to study Ga self-diffusion in GaAs (table 15). At temperatures 

ranging from 850 to 960C, secondary ion mass spectrometric data indicated an activation 

enthalpy of 4eV for Ga self-diffusion. This value was larger than those previously found 

for Ga self-diffusion and Al-Ga interdiffusion under thermal equilibrium and intrinsic 

conditions; which were characterized by an activation enthalpy of 6eV. Secondary ion 

mass spectroscopic, capacitance-voltage, and transmission electron microscopic data 

showed that the as-grown superlattice layers were intrinsic. They became p-type, with 

hole concentrations of about 2 x 10 17 /cm 3 , after annealing. This occurred because the 

layers contained C. Dislocations were also present, at a density of 10 6 to 10 7 /cm 2 . 

However, the factor which was responsible for the larger Ga self-diffusivity values which 

were observed here appeared to be Si out-diffusion from the substrate. Such out-diffusion 

decreased the electron concentration in the substrate, and caused the release of Ga 

vacancies into the superlattice layers, where they became supersaturated. This Ga 

vacancy supersaturation led to enhanced Ga self-diffusion in the superlattice layers. 

T.Y.Tan, H.M.You, S.Yu, U.M.Gösele, W.Jäger, D.W.Boeringer, F.Zypman, R.Tsu, 

S.T.Lee: Journal of Applied Physics, 1992, 72[11], 5206-12 

[446-106/107-036] 

Table 15 

Diffusivity of Ga in GaAs 


960 4.8 x 10 -16 

930 1.3 x 10 -16 

905 5.6 x 10 -17 

875 2.7 x 10 -17 

850 8.6 x 10 -18 

965 3.1 x 10 -16 

935 1.0 x 10 -16 

905 4.5 x 10 -17 

875 2.0 x 10 -17 

850 6.9 x 10 -18 

305



The lateral diffusion of sources during the selective growth of metalorganic vapor-phase 

epitaxial Si-doped layers was analyzed. The diffusion lengths of Ga species were deduced 

from the carrier concentration profiles which were measured by using Raman 

spectroscopy and thickness profiling. On the basis of these diffusion lengths, it was 

speculated that the effective diffusion material was monomethyl Ga. It was suggested that 

there was no difference between arsine and tertiary butyl arsine, as diffusion sources. 

N.Hara, K.Shiina, T.Ohori, K.Kasai, J.Komeno: Journal of Applied Physics, 1993, 74[2], 

1327-30 

[446-106/107-036] 


The formation energy of Si donors, acceptors, and defect complexes were calculated. The 

equilibrium concentrations of native defects and Si-defect complexes were deduced from 

these energies, as was the total solubility of Si. The calculated equilibrium solubility limit 

for Si was in good agreement with experimental data. The (Si Ga -V Ga ) 2- complex occurred 

at relatively high concentrations under As-rich conditions, and could therefore mediate Si 

and Ga diffusion. It was concluded that the donor-vacancy complex was an important 

compensation mechanism in heavily doped GaAs. 

J.E.Northrup, S.B.Zhang: Physical Review B, 1993, 47[11], 6791-4 

[446-106/107-036] 

GaAs/AlAs: Ga Diffusion 

The implantation of Be ions into heterostructures at room temperature or liquid N 

temperatures was investigated. It was found that room-temperature implantation created 

dislocation loops at the first interface; a distance which was far short of the maximum 

projected range. Implantation at low temperatures caused twinning. The latter could be 

removed by annealing (900C, 1200s), without leading to the interdiffusion of Ga. The 

presence of dislocation networks tended to enhance intermixing. The Be concentrations 

were sufficiently low to prevent Be-induced intermixing. 

S.Mitra: Semiconductor Science and Technology, 1990, 5[11], 1138-40 

[446-76/77-017] 

GaAs/AlGaAs: Ga Diffusion 

The effect of room-temperature electron irradiation upon interdiffusion at quantum-well 

interfaces was investigated by using low-temperature cathodoluminescence spectroscopy. 

It was found that interdiffusion was enhanced by the presence of defects which were 

generated by irradiation with a 400keV electron beam. After irradiation at room 

temperature to doses of between about 1.5 x 10 17 and 2.5 x 10 17 /cm 2 , followed by rapid 

thermal annealing (900C, 60s), an interdiffusion length of 0.3 to 0.5nm was found. The 

resultant damage tended to saturate with increasing irradiation dose. The formation of 

306

Ga GaAs Ge 

defect clusters at high doses limited the degree of defect introduction, and therefore the 

extent of interdiffusion at the interface. 

Y.J.Li, M.Tsuchiya, P.M.Petroff: Applied Physics Letters, 1990, 57[5], 472-4 

[446-76/77-018] 

GaAs/Si(O,N): Ga Diffusion 

The out-diffusion of Ga atoms from a GaAs substrate, and into a SiO x N y encapsulating 

film, during annealing was studied by using secondary ion mass spectrometry. The 

concentration of Ga atoms which was detected within the encapsulant, when annealed at 

850C, was found to increase with an increasing O content in the encapsulant. This 

behavior could be correlated with changes in the concentration of the EL5 electron trap 

(E c - E T = 0.42eV); as detected by using deep-level transient spectroscopy. It was 

concluded that the generation of EL5 traps during annealing was due to Ga out-diffusion. 

M.Kuzuhara, T.Nozaki, T.Kamejima: Journal of Applied Physics, 1989, 66[12], 5833-6 

[446-74-029] 

Ge 

316 GaAs: Ge Diffusion 

This dopant diffused extensively after implantation and long-term annealing. The results 

could be explained by assuming that the diffusivity depended upon the square of the 

electron concentration. The dopant diffusion was affected by the presence of implantation 

damage; the higher the concentration of extended defects, the slower was the diffusivity 

as compared with the values for conventional diffusion from a solid source. If the sample 

was amorphized during implantation, extended defects did not form and the diffusivity of 

the ion was very close to that in material which had been diffused from a solid source. 

When amorphization did not occur, extended defects formed after implantation, and 

diffusion was inhibited; especially after low doses, in the short term, or at low 

temperatures. The higher the density of extended defects, the greater was the suppression 

of diffusion. No time-dependence was observed. It was concluded that the results (table 

16) were consistent with a diffusion mechanism in which the mobile species was the 

donor that was coupled with a charged Ga vacancy. The equilibrium vacancy 

concentration was thought to be suppressed by the presence of extended defects and/or 

excess Ga interstitials which resulted from implantation. 

E.L.Allen, J.J.Murray, M.D.Deal, J.D.Plummer, K.S.Jones, W.S.Rubart: Journal of the 

Electrochemical Society, 1991, 138[11], 3440-9 

[446-84/85-016] 

GaAs: Ge Diffusion 

The behavior of Ge which was pulse-diffused into high-purity epitaxial GaAs from a thin 

elemental source, using rapid thermal processing, was investigated. A comparison of 

secondary ion mass spectrometry and differential Hall effect measurements showed that 

the resultant n + -doped layers were highly compensated. In contrast to the case where Ge 

307

Ge GaAs Ge 

was introduced during crystal growth or by ion implantation, Ge Ga donors were not 

compensated by Ge As acceptors when Ge was pulse-diffused into GaAs. 

Photoluminescence spectroscopy showed that Ge Ga donors were compensated by V Ga 

acceptors rather than by Ge As acceptors. At low diffusion temperatures, Ga vacancies 

were formed as Ga rapidly diffused into the Ge layer. These vacancies suppressed the 

formation of As vacancies and thus Ge As acceptors. At higher diffusion temperatures, 

Ge Ga -V Ga complexes were formed more rapidly than V Ga acceptors. Low-temperature 

Hall effect measurements suggested that these complexes were neutral. The formation of 

complexes, at the expense of isolated Ge Ga donors and V Ga acceptors, was related to 

diffusion temperatures which exceeded 865C (the Ge-GaAs liquidus) and was explained 

by a marked increase in the V Ga concentration in the near-surface region, due to Ga 

dissolution at the Ge-Ga-As liquidus. 

C.W.Farley, T.S.Kim, S.D.Lester, B.G.Streetman, J.M.Anthony: Journal of the 


[446-60-003] 

Table 16 

Diffusivity of Implanted Ge in GaAs 

Dose (/cm 2 ) Temperature (C) D (cm 2 /s) 

1 x 10 14 1000 4.8 x 10 -14 

1 x 10 14 900 7.8 x 10 -15 

1 x 10 14 900 2.9 x 10 -15 

1 x 10 14 900 1.8 x 10 -15 

1 x 10 14 900 8.5 x 10 -16 

1 x 10 14 900 5.0 x 10 -16 

5 x 10 15 850 6.4 x 10 -16 

1 x 10 15 850 4.8 x 10 -16 

1 x 10 14 870 3.0 x 10 -16 

5 x 10 15 750 7.6 x 10 -17 

1 x 10 15 750 1.8 x 10 -17 

1 x 10 14 870 8.9 x 10 -17 

5 x 10 14 750 6.6 x 10 -18 

5 x 10 14 850 4.9 x 10 -17 

5 x 10 14 750 4.8 x 10 -18 

5 x 10 13 850 1.0 x 10 -17 

1 x 10 14 950 2.5 x 10 -17 

1 x 10 14 950 1.5 x 10 -16 

5 x 10 15 950 1.4 x 10 -15 

5 x 10 14 950 2.0 x 10 -15 

1 x 10 15 950 2.9 x 10 -15 

1 x 10 15 950 3.8 x 10 -15 

308

Ge GaAs H 

GaAs: Ge Diffusion 

Implantation (3 x 10 13 /cm 2 ) of 200keV Ge into (x11)A-oriented semi-insulating GaAs 


(100)-oriented substrates were also implanted. No in-diffusion of Ge was observed after 

annealing substrates of any orientation. 



[446-117/118-166] 

1.0E-07 

1.0E-08 

1.0E-09 

D (cm 2 /s) 

1.0E-10 

1.0E-11 

1.0E-12 

table 17 

table 18 

1.0E-13 

13 14 15 16 17 18 19 20 

10 4 /T(K) 

Figure 5: Diffusivity of H in GaAs 

H 

317 GaAs: H Diffusion 

It was recalled that the diffusion of neutral H atoms into a semiconductor was 

accompanied by their binding into molecules. When the thermal dissociation of molecules 

could be ignored, and for times which were sufficiently long to establish an equilibrium 

309

H GaAs H 

state for the binding of H into complexes with impurity atoms, the formation of molecules 

was the main process which determined the steady-state profile of atomic H which 

formed near to the crystal surface. By analyzing secondary ion mass spectrometry data in 

terms of a model for this process, it was estimated that the diffusivity of H was equal to 

1.4 x 10 -10 cm 2 /s at 360C and equal to 5.5 x 10 -11 cm 2 /s at 390C. When combined with 

another result, it was deduced that the overall data (table 17) could be described by: 

D (cm 2 /s) = 0.0385 exp[-1.13(eV)/kT] 

N.S.Rytova: Fizika i Tekhnika Poluprovodnikov, 1991, 25[6], 1078-80 (Soviet Physics - 

Semiconductors, 1991, 25[6], 650-1) 

Table 17 

Diffusivity of H in GaAs 


390 5.2 x 10 -11 

360 1.2 x 10 -10 

250 1.0 x 10 -12 

GaAs: H Diffusion 

Layers of material, which was doped with various group-VI donors (S, Se, Te), were 

exposed to H plasma. Electronic measurements indicated that, after H diffusion, the 

electron concentration systematically decreased while their mobility increased; thus 

demonstrating the passivation of the group-VI donors by H. 

B.Theys, B.Machayekhi, J.Chevallier, K.Somogyi, K.Zahraman, P.Gibart, M.Miloche: 


[446-121/122-052] 


The ability of charged H to drift in this material was used to investigate the behavior of 

H. It was concluded that diffusion within the above temperature range was entirely traplimited, 

and exhibited no dependence upon the diffusivity of free H. By drifting H away 

from the donors in hydrogenated n-type GaAs, reactivation of the passivated donors could 

be studied. Thermal dissociation of the donor-H complex obeyed first-order kinetics, with 

a dissociation energy of 1.52eV. 

A.W.R.Leitch, T.Zundel, T.Prescha, J.Weber: Materials Science Forum, 1992, 83-87, 21- 

6 

[446-93/94-009] 


First-principles calculations were made of the properties of atomic and molecular H in 

pure bulk material. The results indicated that the H penetrated in atomic form. Within the 

sample, atomic H tended to form H 2 molecules on tetrahedral sites. These were deep 

energy wells for H 2 . The H 2 * defect, which consisted of a H atom in a bond-center site 

310


and a H atom in an adjacent tetrahedral position, had a higher energy than H 2 but was a 

lower-energy barrier to diffusion. Isolated H could be present as a metastable species. 

The stable charge state of isolated H was calculated as a function of the Fermi energy. 

The results suggested that H behaved as a negative-U defect. Consequently, isolated H 

was expected to be present only as a charged species (positively charged in p-doped 

samples, negatively charged in undoped and n-doped samples). The conclusions were 

compared with experimental results and with the results of calculations for H in other 

semiconductors. The main features of H behavior in GaAs were quite similar to those for 

Si. 

L.Pavesi, P.Giannozzi: Physical Review B, 1992, 46[8], 4621-9 

[446-93/94-009] 


It was pointed out that Se donors in n-type material were passivated by exposure to a H 

plasma. Thermal reactivation of the passivated donors was investigated at temperatures 

ranging from 154 to 191C. By using the electric field of a reverse-biased Schottky diode 

structure, the drift of thermally dissociated H as a negatively charged species was 

demonstrated. The thermal dissociation of the SeH complex obeyed first-order kinetics, 

with a dissociation energy of 1.52eV. 

A.W.R.Leitch, T.Prescha, J.Weber: Physical Review B, 1991, 44[3], 1375-8 

[446-84/85-016] 


It was shown that the exponential depth profiles which were sometimes observed during 

H diffusion into semiconductors, such as Si and GaAs, could be explained by including a 

term (in the diffusion equation) that described the multiple trapping of H at an impurity. It 

was shown that the effective dimensionality of the random walk in the present case was 

infinite. 

D.A.Tulchinsky, J.W.Corbett, J.T.Borenstein, S.J.Pearton: Physical Review B, 1990, 

42[18], 11881-3 

[446-81/82-010] 


The concentration versus depth profiles of carriers, electrically active defects, and D, 

after exposure to a H plasma (or molecular H), were fitted by using a simple diffusion 

model which involved second-order reactions. It was found that the activation energy for 

H diffusion, and the dissociation energies of H-defect complexes, depended upon the H 

concentration. There was no molecular H formation and there were no fast-diffusing H 

species away from the near-surface region. Atomic H could in-diffuse and passivate EL2 

defects when semi-insulating material was annealed at a high temperature in a molecular 

H ambient. 

R.A.Morrow: Journal of Applied Physics, 1989, 66[7], 2973-9 

[446-74-010] 

311



Wafers of n-type material were exposed to a capacitively coupled radio-frequency H 

plasma, with power densities of between 0.01 and 0.2W/cm 2 at 260C. The properties of 

the layers were investigated by using deep-level transient spectroscopy and capacitancevoltage 

methods. As well as the neutralization of Si donors by in-diffused H atoms, it was 

found that there was a modification of the deep-level transient spectroscopy spectra after 

hydrogenation. At two radio-frequency power densities of less than 0.IW/cm 2 , deep levels 

in the original material were passivated. The results also indicated that the complexes 

which were present were either electrically inactive or were located deep within the 

energy band-gap. At power densities of more than 0.1W/cm 2 , two new deep states 

appeared at 0.41 and 0.55eV below the conduction band. These levels were attributed to a 

large number of defects which were situated in the near-surface region of n-type Si-doped 

material after H plasma exposure. It was suggested that the trapping of H at these defects 

was probably responsible for the observed accumulation of H in the near-surface region. 

A.Jalil, A.Heurtel, Y.Marfaing, J.Chevallier: Journal of Applied Physics, 1989, 66[12], 

5854-61 

[446-74-010] 


The depth profiles of 60 and 100keV protons which were implanted to fluences of 10 16 or 

10 17 /cm 2 at room temperature were determined by using ion beam techniques. The H 

profiles were measured as a function of annealing temperatures of up to 820K. It was 

found that the redistribution of implanted H depended upon the migration of 

implantation-induced defects. The migration of H-defect complexes was described by the 

expression: 

D(cm 2 /s) = 2 x 10 5 exp[-2.16(eV)/kT] 

J.Räisänen, J.Keinonen, V.Karttunen, I.Koponen: Journal of Applied Physics, 1988, 

64[5], 2334-6 

[446-72/73-010] 

318 GaAs: H Diffusion 

Profiles which reflected the passivation of electrically active and recombination-active 

centers by atomic H in single crystals were investigated by using secondary ion mass 

spectrometry and microcathodoluminescence methods. The former data indicated that the 

profiles were the same for all of the samples, and consisted of an 0.0005mm surface 

region with a H content of 10 20 /cm 3 and a low diffusion rate. This was followed by a 

diffusion tail with a much higher diffusion rate. In this tail region, the activation energy 

for H diffusion was 0.83eV at temperatures of between 200 and 500C (table 18), and the 

diffusivity was 6.7 x 10 -9 cm 2 /s at 400C. It was proposed that the results confirmed the 

suggested migration of neutral interstitial H. The introduction of atomic H completely 

suppressed a microcathodoluminescence inhomogeneity which was associated with 

dislocations in the semi-insulating material. 

312


E.M.Omelyanovskii, A.V.Pakhomov, A.J.Polyakov, A.V.Govorkov, O.M.Borodina, 

A.S.Bruk: Fizika i Tekhnika Poluprovodnikov, 1988, 22[7], 1203-7. (Soviet Physics - 

Semiconductors, 1988, 22[7], 763-5) 

[446-62/63-210] 

Table 18 

Diffusivity of H in GaAs 


480 2.3 x 10 -8 

425 5.1 x 10 -8 

445 1.4 x 10 -8 

445 1.0 x 10 -8 

420 1.2 x 10 -8 

410 7.7 x 10 -9 

410 6.2 x 10 -9 

390 2.2 x 10 -9 

365 2.8 x 10 -9 

365 1.9 x 10 -9 

320 2.2 x 10 -9 

320 9.4 x 10 -10 

290 6.0 x 10 -10 

295 2.7 x 10 -10 

255 2.3 x 10 -10 

255 1.3 x 10 -10 


Experiments were performed on buried Si-doped epilayers under hydrostatic pressure. It 

was found that the D diffusion profile in n-type Si-doped samples depended upon the 

hydrostatic pressure, and consisted of a plateau that was followed by a steep progressive 

decrease as the pressure was increased. This was explained as being due to an increasing 

importance of the trapping-detrapping process of H - on Si + donors during H diffusion. 

This increase was attributed to a deepening of the H acceptor level, with respect to the 

bottom of the Γ conduction band, as the hydrostatic pressure was increased. 

D.Machayekhi, J.Chevallier, B.Theys, J.M.Besson, G.Weill, G.Syfosse: Solid State 

Communications, 1996, 100[12], 821-4 

[446-141/142-098] 


The results of investigations of transitions among the sites and charge states of muonium 

were summarized. A model was developed which accounted for all of the major features 

313


which were observed. Its validity for a wide range of dopant concentrations, using a 

single set of parameters, reflected its predictive strength. The near equality of the energy 

parameters for muonium, as compared with those which were available for H, strongly 

implied that the results for muonium dynamic behavior should be applicable to H, with 

very little change. The model could be applied to all tetrahedrally coordinated 

semiconductors, with few modifications, and served as a basis for the understanding of 

muonium dynamics and therefore for the behavior of H in GaAs and other materials. 

Differences in H properties could be understood by examining material-specific 

deviations from the basic model. 

R.L.Lichti, C.Schwab, T.L.Estle: Materials Science Forum, 1995, 196-201, 831-6 

[446-127/128-119] 


The diffusivity of H + was directly determined by means of capacitance transient 

measurements. The H was released from donor-H complexes, in the space-charge layer of 

Schottky diodes on n-type material, by the pulsed-laser injection of minority carriers. 

Capacitance-voltage measurements revealed a recovery of donor dopants after injection. 

This demonstrated minority-carrier enhanced dissociation of a donor-H complex. 

Capacitance transients which were recorded during the migration of H + were analyzed in 

order to obtain diffusivities at near to room temperature. At 320K, the diffusivity was 

equal to 10 -12 cm 2 /s, to within a factor of 2. An Arrhenius plot of the migration timeconstant 

yielded an activation energy for H + diffusion of about 0.66eV. 

N.M.Johnson, C.Herring, D.Bour: Physical Review B, 1993, 48[24], 18308-11 

[446-115/116-117] 


Diffusion experiments were performed on samples of Si-doped Al x Ga 1-x As epitaxial 

layers, with x-values which ranged from 0 to 0.30, as a function of the Si doping level 

and the diffusion temperature. For each composition, calculated H diffusion profiles 

which had been deduced by using Mathiot's model were fitted to the experimental 

profiles. It was assumed that H behaved as a deep acceptor, and that H o and H - were the 

diffusing species. The trapping of H - by Si + donors, and their acceleration by an electric 

field, were incorporated into the model. As well as the diffusion coefficient of H, and the 

dissociation constant of the SiH complexes, the model provided for a compositional 

dependence of the H acceptor level in AlGaAs alloys. Extrapolation of the H acceptor 

level to x = 0 gave its value for GaAs. This level was located at 0.10eV below the Γ 

conduction band of GaAs. 

B.Machayekhi, R.Rahbi, B.Theys, M.Miloche, J.Chevallier: Materials Science Forum, 

1994, 143-147, 951-6 

[446-113/114-001] 


The depth profiles of diffused H in n-type samples were determined by means of 

secondary ion mass spectroscopy. Specimens with Si donor concentrations which ranged 

314

H GaAs In 

from 10 17 to 5 x 10 18 /cm 3 were exposed to monatomic D from a remote microwave 

plasma at temperatures of between 250 and 310C. The profiles clearly revealed that D 

accumulation occurred up to a concentration that was almost equal to that of the donors 

over a depth which depended upon the temperature and time of hydrogenation and upon 

the donor concentration. A plateau in the H concentration was attributed to a trappinglimited 

migration of H which was dominated by the formation of SiH complexes via the 

Coulombic attraction of Si + and H - . The analysis of profiles in the tail region beyond the 

plateau yielded independent estimates of the thermal dissociation rate for the SiH 

complex, and a lower limit of about 0.45eV for the binding of Si + and H - into SiH. 

G.Roos, N.M.Johnson, C.Herring, J.Walker: Materials Science Forum, 1994, 143-147, 

933-8 

[446-113/114-011] 

In 

GaAs: In Diffusion 

An InAs monolayer was grown between GaAs layers by using the migration-enhanced 

epitaxy method. The surface chemical characteristics during growth were investigated by 

means of reflection high-energy electron diffraction. When the substrate temperature was 

equal to about 500C, the oscillation amplitude of the reflected electron beam after the 

growth of one monolayer of InAs vanished during the growth of GaAs over more than 20 

atomic layers. High-resolution secondary ion mass spectroscopic analysis of the 

fabricated structures indicated that an anomalous distribution of In atoms, with a gradual 

slope towards the growth direction, occurred when the substrate temperature was 500C. 

The experimental results were explained in terms of the replacement of In atoms, in the 

InAs monolayer, by newly deposited Ga atoms. 

H.Yamaguchi, Y.Horikoshi: Japanese Journal of Applied Physics, 1989, 28[11], L2010-2 

[446-72/73-011] 

GaAs: In Diffusion 

The diffusion of In markers at 900C was measured in undoped and Te-doped epilayers 

which had been prepared by using organometallic vapor-phase epitaxy. It was found that 

the diffusivity was a linear function of electron concentrations ranging from 2 x 10 17 to 

1.5 x 10 19 /cm 3 . It was concluded that the results were consistent with the interdiffusion of 

AlAs and GaAs superlattices. Also, the In and Al diffusivities at 900C were essentially 

identical; within experimental error. The results strongly suggested that group-III 

interdiffusion in this material was controlled by a Ga vacancy with a charge of -1. 

W.M.Li, R.M.Cohen, D.S.Simons, P.H.Chi: Applied Physics Letters, 1997, 70[25], 3392- 

4 

[446-152-0315] 

315

In GaAs Li 

GaAs/InAs: In Diffusion 

The segregation and interdiffusion of In atoms in GaAs/InAs/GaAs heterostructures were 

investigated by using secondary ion mass spectroscopy. When the InAs was grown in the 

layer-by-layer growth mode, with no dislocations, the segregation of In atoms became 

marked with increasing growth temperature. However, segregation was observed even at 

the relatively low growth temperature of 400C during molecular beam epitaxy. It was 

found that segregation was markedly enhanced by dislocations which were near to the 

hetero-interface when thick InAs layers were grown in a 3-dimensional island growth 

manner. Interdiffusion of In atoms towards the growth direction occurred after thermal 

annealing, and could be assisted by vacancies which propagated from the film surface and 

into the epilayer. It became apparent that interdiffusion was effectively suppressed by 

inserting a thin AlAs layer into the GaAs cap layer. 

T.Kawai, H.Yonezu, Y.Ogasawara, D.Saito, K.Pak: Journal of Applied Physics, 1993, 

74[3], 1770-5 

[446-106/107-085] 

GaAs/Si: In Diffusion 

Samples of GaAs, which were encapsulated with thin films of amorphous Si at 450C, 

were annealed at temperatures of up to 1050C. The resultant poly-Si/GaAs interfaces 

were investigated by using secondary ion mass spectroscopy, Rutherford back-scattering 

spectrometry, and transmission electron microscopy. Little or no interdiffusion was 

detected at undoped Si/GaAs interfaces. The diffusion of dopants such as InP was 

detected. An enhanced diffusivity of In into GaAs was attributed to the diffusion of point 

defects which were created by the diffusion of As and Ga into the encapsulant. It was 

deduced that the In diffusivities in GaAs at doped polycrystalline Si interfaces were 

enhanced by factors of about 10000. 

K.L.Kavanagh, C.W.Magee, J.Sheets, J.W.Meyer: Journal of Applied Physics, 1988, 

64[4], 1845-54 

[446-72/73-027] 

Li 

GaAs: Li Diffusion 

Strong photoluminescence bands, at 1.34 and 1.45eV, were observed after the Li 

diffusion of various samples of n-type material at temperatures of between 400 and 600C. 

The photoluminescence bands were homogeneous throughout the bulk of the samples. 

They were never observed in originally p-type or semi-insulating material. A hole at E v + 

0.17eV was detected by means of deep-level transient spectroscopic measurements, and 

an acceptor level at E v + 0.14eV was revealed by temperature-dependent Hall 

measurements of n-type material which had been converted to p-type by Li diffusion. It 

316

Li GaAs Mg 

was suggested that Li Ga -D acceptors were connected with compensation and 

photoluminescence bands. 

B.H.Yang, T.Egilsson, S.Kristjánsson, J.Pétursson, H.P.Gislason: Materials Science 

Forum, 1994, 143-147, 839-44 

[446-113/114-012] 

GaAs: Li Diffusion 

The correlation between the electrical conductivity and photoluminescence spectra of a 

wide range of n-type, p-type, and semi-insulating material which had been diffused with 

Li was investigated. It was found that Li diffusion compensated both n-type and p-type 

material, and had a marked effect upon the photoluminescence properties of the samples. 

The photoluminescence bands which were observed, between 1.47 and 1.49eV in Zndoped 

GaAs, were related to the compensation of Zn acceptors. The photoluminescence 

bands, which were observed at 1.43 to 1.45eV in all of the samples, were attributed to a 

Li-related acceptor complex which gave rise to the p-type conductivity which was 

observed in all of the samples after full compensation. 

H.P.Gislason, B.Yang, I.S.Haukson, J.T.Gudmundsson, M.Linnarsson, E.Janzén: 

Materials Science Forum, 1992, 83-87, 985-90 

[446-99/100-064] 

Mg 

GaAs: Mg Diffusion 


using metalorganic vapor-phase epitaxy, was studied by using capacitance-voltage etch 

profiling and secondary ion mass spectrometry. It was deduced that the diffusivity of Mg 

in GaAs could be described by: 

D (cm 2 /s) = 6.5 x 10 -5 exp[-1.8(eV)/kT] 

for rapid thermal annealing, while the diffusivity could be described by: 

D (cm 2 /s) = 1.9 x 10 -7 exp[-1.5(eV)/kT] 

for furnace annealing. A model which was based upon an interstitial cum substitutional 

diffusion mechanism, with certain kinetic limitations, was successfully used to simulate 

the observed dopant concentration profiles. 

N.Nordell, P.Ojala, W.H.Van Berlo, G.Landgren, M.K.Linnarsson: Journal of Applied 

Physics, 1990, 67[2], 778-86 

[446-74-004] 


The transient up-hill diffusion of Mg during annealing at 900C was computer-simulated. 

It was assumed that diffusion occurred via the substitutional-interstitial mechanism, with 

excess interstitials and vacancies being produced by implantation; thus causing the 

abnormal diffusion behavior. The substitutional-interstitial mechanism was shown to be 

mathematically equivalent to an existing interstitial-dopant pair diffusion model. This 

317

Mg GaAs Mg 

permitted the programme (a Si process simulator that included dopant/point-defect 

interactions) to be used to model up-hill diffusion if suitable diffusivity and defect 

parameters were included. The profiles of excess interstitials and vacancies which were 

produced by the implantation process were deduced from Boltzmann transport equation 

calculations. It was found that transient up-hill diffusion could be accurately simulated; 

with the dopant diffusing, from regions with excess interstitials, towards the surface or 

towards regions with excess vacancies. When the defect concentrations had returned to 

their steady-state levels, via diffusion, recombination, or capture by sinks, normal 

concentration-dependent diffusion into the substrate occurred. 

H.G.Robinson, M.D.Deal, G.Amaratunga. P.B.Griffin, D.A.Stevenson, J.D.Plummer: 


[446-86/87-010] 


It was noted that Mg implants in this material exhibited 2 types of diffusion during 

annealing. These were up-hill diffusion in the implantation peak, and concentrationdependent 

diffusion into the bulk. The up-hill diffusion predominated at short times and 

low temperatures, while the concentration-dependent diffusion was predominant at long 

times and high temperatures. By studying implants that had been annealed at temperatures 

where no up-hill diffusion occurred, diffused profiles could be modelled and an 

expression for the Mg diffusivity could be obtained. The activation energy for this 

process was 1.77eV. The results of Fermi level experiments showed that the diffusivity 

was hole-dependent rather than concentration dependent. The hole-dependent exponent 

was equal to unity for Mg which was implanted into semi-insulating substrates, but could 

change to 2 at high hole concentrations. 


[446-81/82-010] 


The redistribution of Mg implants during post-implantation annealing was studied in 

order to evaluate the effect of implantation damage upon the diffusion process. Rapid uphill 

diffusion was observed in the peak of Mg implants. This was explained by invoking a 

substitutional-interstitial diffusion mechanism and by performing computer simulations of 

damage-generated point defects. In the up-hill diffusion region, the dopants diffused from 

areas of excess interstitial concentration towards areas of excess vacancy concentration. 

A critical point defect concentration was necessary in order to initiate up-hill diffusion. 


[446-74-009] 


Un-capped wafers were annealed in an AsH 3 -N 2 atmosphere for 10s, at temperatures of 

up to 1100C. No surface decomposition occurred under an AsH 3 partial pressure of 

12.5torr. The present method was used to activate implanted Mg. Measurements of the 

sheet resistance of the annealed layers, as a function of the annealing temperature, 

318

Mg GaAs Mn 

revealed a minimum at a temperature of about 930C. At higher temperatures, the 

diffusion of Mg became significant. Part of the Mg accumulated at the surface and 

diffused out. The internal diffusion of Mg at high temperatures depended upon the AsH 3 

pressure during annealing. 

H.Tews, R.Neumann, A.Hoepfner, S.Gisdakis: Journal of Applied Physics, 1990, 67[6], 

2857-61 

[446-74-010] 


The surface of material which had been implanted with Mg ions (100keV, 10 15 /cm 2 ) was 

encapsulated with As-doped amorphous hydrogenated Si to a thickness of about 80nm. It 

was found that the sheet carrier concentration in thermally annealed samples increased 

with increasing concentration of As in the encapsulant. After annealing (800C, 0.25h), a 

sheet carrier concentration of about 3 x 10 14 /cm 2 was observed in samples which were 

capped with films that were doped with 2 x 10 20 /cm 3 of As. It was noted that the 

diffusivity of the implanted Mg was retarded upon increasing the concentration of As in 

the amorphous hydrogenated Si encapsulant. 

K.Yokota, M.Sakaguchi, H.Inohara, H.Nakanishi, S.Tamura, Y.Horino, A.Chayahara, 

M.Satho, K.Hira, H.Takano, M.Kumagaya: Solid-State Electronics, 1994, 37[1], 9-15 

[446-113/114-012] 

Mn 

GaAs: Mn Diffusion 

The diffusion of Mn was carried out in sealed quartz ampoules, using 4 types of Mn 

source. These were: solid crystalline Mn grains, Mn 3 As, MnAs, and Mn thin films on 

GaAs substrates. It was found that only MnAs led to the formation of a smooth GaAs 

surface and a uniform dopant distribution. In the case of the other sources, interactions 

between the source materials and the substrate gave rise to poor surface morphologies and 

inhomogeneous distributions. In the case of diffusion at 800C, surface p-type carrier 

concentrations of the order of 10 20 /cm 3 were obtained. The diffusion profiles which were 

determined by using capacitance-voltage techniques resembled those which were 

obtained for Zn diffusion. It was suggested that a substitutional-interstitial mechanism 

was the predominant one. 

C.H.Wu, K.C.Hsieh, G.E.Höfler, N.El-Zein, N.Holonyak: Applied Physics Letters, 1991, 

59[10], 1224-6 

[446-84/85-007] 


Plates of Sn-doped n-type material were coated with 54 Mn, annealed (1100C, 4h) and 

analyzed by using alternate grinding and activity measurements. The diffusion profiles 

could be described by an erfc function, and the diffusivity at the annealing temperature 

was estimated to be 4.3 x 10 -10 cm 2 /s. 

319

Mn GaAs Ni 

E.A.Skoryatina, R.S.Malkovich: Fizika i Tekhnika Poluprovodnikov, 1989, 23[1], 164-6. 

(Soviet Physics - Semiconductors, 1989, 23[1], 101-2) 

[446-70/71-106] 


It was pointed out that one of the potential advantages of rapid thermal annealing, as 

compared with conventional furnace annealing, was a reduced implanted dopant and 

background impurity diffusion. Here, the migration of Mn during the annealing of Crdoped 

semi-insulating material implanted with 100keV Si + ions to a dose of 7 x 10 12 /cm 2 

was measured by using secondary ion mass spectrometry. Un-capped rapid thermal 

annealing (860 or 930C, 1 to 60s) was investigated and its effect was compared with that 

of capless furnace annealing (0.5h). The migration of Mn was undetectable for rapid 

thermal annealing times which were shorter than 60s, but dominated 0.5h furnace 

anneals. 

H.Kanber, J.M.Whelan: Journal of the Electrochemical Society, 1987, 134[10], 2596-9 

[446-55/56-005] 


A study was made of the effect of an As over-pressure upon Mn diffusion. The sources of 

Mn included solid Mn thin films, which were deposited directly onto the GaAs substrate, 

and Mn vapor from pure solid Mn, MnAs, or Mn 3 As. When a solid Mn film was used as 

the diffusion source, a non-uniform dopant distribution and a poor surface morphology 

was obtained. This was due to reaction between the Mn film and the GaAs matrix. The 

degraded surface consisted of a layer of polycrystalline cubic alloy with a lattice constant 

of almost 0.84nm, and with a composition that was close to Ga 2 Mn; with a small amount 

of As. Of the remaining diffusion sources, only MnAs consistently produced a uniform 

doping distribution and a smooth surface morphology. By diffusion at 800C, a uniform 

surface hole carrier concentration as high as 10 20 /cm 3 could be obtained by using MnAs 

as a source. The As over-pressure was found to alter drastically the Mn diffusion profile 

and Mn, like Zn, could diffuse interstitial-substitutionally. Vapor from Mn or Mn 3 As 

degraded the surface. However, Mn 3 As degraded the surface more rapidly. A sufficiently 

high As over-pressure completely suppressed surface degradation. 

C.H.Wu, K.C.Hsieh: Journal of Applied Physics, 1992, 72[12], 5642-8 

[446-106/107-037] 

Ni 

GaAs: Ni Diffusion 

The redistribution of Ni in a Ni/GaAs contact specimen was studied by using neutron 

activation and sectioning techniques at temperatures of between 360 and 870K. It was 

concluded that interdiffusion of the components took place in an elastic strain field which 

was generated by differences in the lattice parameters and thermal expansion coefficients 

of Ni, GaAs and intermetallic compounds which formed. At annealing temperatures below 

570K, reactive diffusion of Ni took place, with an activation energy of 0.51eV. The 

320

Ni GaAs O 

formation of micro-cracks in the surface layers of the microcrystalline GaAs led to 

diffusion with an activation energy of 0.25eV. At annealing temperatures that were 

greater than 670K the internal electric field, and cluster formation, markedly affected the 

distribution of the components. 

W.A.Uskov, A.B.Fedotov, E.A.Erofeeva, A.I.Rodionov, D.T.Dzhumakulov: Izvestiya 

Akademii Nauk SSSR - Neorganicheskie Materialy, 1987, 23[2], 186-9. (Inorganic 

Materials, 1987, 23[2], 163-5) 

[446-55/56-006] 

O 

GaAs: O Diffusion 

A new technique was developed in order to study atomic movements during ultra-violet 

laser-enhanced and low-temperature (below 400C) thermal oxidation. This method 

combined a classical marker technique with low-energy ion-scattering spectroscopy. 

During the formation of thin (about 1nm) oxide layers, the marker remained on the oxide 

surface. This indicated that oxidation occurred, at the GaAs/oxide interface, via the 

diffusion of an O species. This differed from the oxidation of metals such as Ni and Cu, 

where the use of the same technique confirmed earlier observations that oxidation 

occurred at the oxide/ambient interface. The diffusion of a metal species resulted in the 

marker being buried during oxidation of the metal surfaces. 

M.T.Schmidt, Z.Wu, C.F.Yu, R.M.Osgood: Surface Science, 1990, 226[1-2], 199-205 

[446-74-011] 


The thermal surface oxidation of samples in dry O 2 and in ambient air was investigated at 

temperatures ranging from 400 to 530C. The studies were carried out by using clean 

polished (111) substrates. It was found that monocrystalline wafers oxidized parabolically 

in dry O 2 at temperatures of 400 to 450C. Linear oxide growth occurred, at temperatures 

ranging from 480 to 530C, in both in dry O 2 and in ambient air. Wagner and Grimley- 

Trapnell models for metal oxidation were used to identify the growth kinetics. The rate 

constants for parabolic and linear oxidation were temperature-dependent, and satisfied 

Arrhenius relationships. 

J.Kucera, K.Navratil: Thin Solid Films, 1990, 191[2], 211-20 

[446-78/79-012] 


Published depth profiles of D in n-type and p-type material, after extended annealing in D 

at 500C, were suggested to reflect the in-diffusion of a native defect (perhaps V As ) and an 

impurity (perhaps O); both of which were tagged with D. It was deduced that the 

diffusivity of the tagged impurity was equal to 4 x 10 -14 cm 2 /s at 500C. The diffusivity of 

321

O GaAs S 

the tagged native defect was deduced to be equal to 3 x 10 -15 cm 2 /s in n-type material, and 

to be equal to about 8 x 10 -15 cm 2 /s in p-type material. 

R.A.Morrow: Applied Physics Letters, 1990, 57[3], 276-8 

[446-76/77-009] 


Infra-red measurements of the concentrations of H-B pairs, which formed in B-doped Si 

that had been heated in H 2 gas and quenched from temperatures of between 900 and 

1300C, led to a new determination of the H solubility: 

[H](/cm 3 ) =9.1 x 10 21 exp[-1.80(eV)/kT] 

There was some evidence that H 2 molecules also formed. The presence of H led to an 

enhancement of the diffusivity of O impurities at temperatures below 500C. Suggestions 

that H was present in as-grown Czochralski Si were supported by the observation of H-C 

complexes, using photoluminescence techniques. The analysis of the structure of a H 

complex, by means of infra-red vibrational spectroscopy, was illustrated for the case of 

the H-C As pair in GaAs. 

R.C.Newman: Philosophical Transactions of the Royal Society A, 1995, 350[1693], 215- 

26 

[446-119/120-192] 

P 

GaAs/Si: P Diffusion 





detected at undoped Si/GaAs interfaces. The diffusion of dopants such P was detected. 

An enhanced diffusivity of P into GaAs was attributed to the diffusion of point defects 

which were created by the diffusion of As and Ga into the encapsulant. It was deduced 

that the P diffusivities in GaAs at doped polycrystalline Si interfaces were enhanced by 

factors of about 10000. 


64[4], 1845-54 

[446-72/73-027] 

S 

GaAs: S Diffusion 

Experiments were performed on polished plates of Te-doped material. Irradiation with 15 

to 150keV protons was carried out at 300K to doses of between 10 16 and 10 17 /cm 2 . It was 

found that the impurity profiles did not depend upon whether the diffusion source was 

322

S GaAs S 

deposited before or after irradiation. The penetration depth in samples which were 

irradiated with 15keV protons was greater than that in samples which were irradiated with 

150keV protons. It was suggested that this was because the low-energy ions generated 

more defects at depths of between 50 and 100nm. 

V.N.Abrosimova, V.V.Kozlovskii, N.N.Korobkov, V.N.Lomasov: Izvestiya Akademii 

Nauk SSSR - Neorganicheskie Materialy, 1990, 26[3], 488-91. (Inorganic Materials, 

1990, 26[3], 411-4) 

[446-84/85-013] 

Table 19 

Diffusivity of Implanted S in GaAs 

120keV S + (/cm 2 ) Temperature (C) D (cm 2 /s) 

1 x 10 15 1000 1.4 x 10 -12 

1 x 10 15 950 9.6 x 10 -13 

1 x 10 15 900 7.1 x 10 -13 

1 x 10 15 850 5.4 x 10 -13 

5 x 10 14 1000 9.9 x 10 -13 

5 x 10 14 950 7.9 x 10 -13 

5 x 10 14 900 4.9 x 10 -13 

5 x 10 14 850 3.0 x 10 -13 

1 x 10 14 1000 5.0 x 10 -13 

1 x 10 14 950 3.4 x 10 -13 

1 x 10 14 900 1.3 x 10 -13 

5 x 10 13 1000 2.6 x 10 -13 

5 x 10 13 950 1.4 x 10 -13 

319,20 GaAs: S Diffusion 

Samples were implanted with 120keV S + ions to doses of between 3 x 10 13 and 10 15 /cm 2 

(table 19). They were capped with an 80nm-thick film of amorphous hydrogenated Si, 

into which As was doped to a concentration of 2 x 10 20 /cm 3 . The samples were then 

annealed in Ar gas (850 to 1000C, 0.25h). It was found that the diffusivity of S (table 20) 

could be described by the expression: 

D = D m [KQ 2 /(1 + KQ 2 )] 

where K was a constant, Q was the implantation dose, and D m was the diffusivity of a 

mobile complex. 

M.Sakaguchi, K.Yokota, A.Shiomi, K.Hirai, H.Takano, M.Kumagai: Japanese Journal of 

Applied Physics, 1996, 35[1-8], 4203-8 

[446-138/139-077] 

323

S GaAs S 


Samples were implanted with 500keV S ions, through 120nm-thick films of amorphous 

hydrogenated Si, to give a concentration of 2 x 10 20 /cm 3 . They were then annealed (700 

to 1000C, 0.25h, Ar). It was found that the diffusivity of the S decreased with increasing 

implantation dose. 

M.Sakaguchi, K.Yokota, A.Shiomi, H.Mori, A.Chayahara, Y.Fujii, K.Hirai, H.Takano, 

M.Kumagai: Japanese Journal of Applied Physics, 1996, 35[1-3], 1624-9 

[446-136/137-110] 

Table 20 

Diffusion Parameters for Implanted S in GaAs 

Dose (/cm 2 ) D o (cm 2 /s) E (eV) 

1 x 10 15 2.0 x 10 -9 0.8 

5 x 10 14 9.0 x 10 -9 1.0 

1 x 10 14 4.2 x 10 -7 1.5 

5 x 10 13 5.5 x 10 -7 1.6 


An investigation was made of the composition and thermal stability of ultra-violet, ozoneoxidized, 

and P 2 S 5 /(NH 4 ) 2 S-treated (100) surfaces. In particular, X-ray photo-electron 

spectroscopy and Auger electron spectroscopy were used to probe the oxide and interface 

at room temperature and after annealing at various temperatures. The room temperature 

data indicated that S was buried between the oxide over-layer and the GaAs substrate. 

This oxide contained various As and Ga bonding configurations which, after moderate 

annealing, were transformed into more thermally stable phases, such as As 2 O 3 and Ga 2 O 3 . 

Complete desorption of the oxide occurred after annealing at 600C. Heating the sample to 

495C caused some S to diffuse towards the oxide surface, while annealing at higher 

temperatures led to S diffusion into the GaAs substrate. Even after complete desorption of 

the O, a small amount of S remained in the GaAs lattice. 

M.J.Chester, T.Jach, J.A.Dagata: Journal of Vacuum Science and Technology A, 1993, 

11[3], 474-80 

[446-111/112-050] 

GaAs/Si: S Diffusion 

The depth distribution of S near to a Si/GaAs(110) interface was measured by using 

particle-induced X-ray emission techniques, together with Rutherford back-scattering 

spectrometry. Ozone oxidation and HF etching were used to remove layers. The 

measurements revealed the presence of a half-monolayer of S on H 2 S x -passivated GaAs 

(110) surfaces. Upon depositing 1.5nm of Si onto S-passivated GaAs (110), the total 

amount of S was found to remain constant as compared to that before Si deposition. 

324

S GaAs Si 

However, no oriented S-Ga bonds were revealed by X-ray absorption measurements, and 

the depth profiles revealed that S atoms diffused into both the GaAs substrate and the Si 

hetero-layer. 

H.Xia, W.N.Lennard, L.J.Huang, W.M.Lau, J.M.Baribeau, D.Landheer: Journal of 


[446-138/139-130] 

Se 

GaAs/Si: Se Diffusion 

The surface of material which had been implanted with 100keV Se ions was encapsulated 

with As-doped amorphous hydrogenated Si to a thickness of about 80nm. Crystallization 

of the encapsulant upon annealing at 1000C was hindered by doping with As atoms to a 

concentration of 2 x10 20 /cm 3 . However, the encapsulant could be crystallized when the 

concentration of the As dopant atoms was lowered. The crystallized encapsulant had 

many grain boundaries, and the diffusion rate of impurities was thereby increased. The 

GaAs surface decomposed markedly via the boundary of the Si/GaAs structure, and As 

and Ga vacancies were produced at the GaAs surface. A number of Si atoms also diffused 

into the GaAs crystal. An As-vacancy rich region formed near to the surface, and the Ga 

vacancy diffused into the GaAs because the diffusivity of the As vacancy in GaAs was 

much lower than that of the Ga vacancy. The Si atoms gave a flat profile, with a steep 

slope at the surface, because of the distribution of the vacancies. That is, the Si atoms 

preferentially occupied vacant As sites in the V As -rich region or vacant Ga sites in the 

V Ga region. A series of reactions, 

Se As + + (Si Ga + + Si As - ) + V Ga - - (Se As + + Si As - ) + (Si Ga + + V Ga - ) → 

(Si As - + Se As + ) + (V Ga - + Si Ga + ) 

increased the diffusion rate of the implanted Se atoms. Thus, the activation efficiency 

improved with increasing concentration of As atoms in the encapsulant, and the diffusion 

rate was reduced because the content of Si atoms and vacancies in GaAs was decreased. 

K.Yokota, K.Nishida, A.Yutani, S.Tamura, Y.Horino, A.Chayahara, M.Satho, K.Hirai, 

H.Takano, M.Kumagaya: Japanese Journal of Applied Physics, 1993, 32[1-10], 4418-24 

[446-115/116-131] 

Si 

321 GaAs: Si Diffusion 

The diffusion of Si was studied in GaAs which had been implanted with 40keV 

30 Si + ions to a dose of 10 16 /cm 2 . The Si concentration profiles were determined by means 

of secondary-ion mass spectrometry and nuclear resonance broadening techniques, and 

the defect distributions were determined by using the Rutherford back-scattering 

spectrometry channelling technique. The implanted samples were subjected to annealing 

in Ar at temperatures ranging from 650 to 850C (table 21). Two independent Si diffusion 

325

Si GaAs Si 

mechanisms were observed. It was found that concentration-independent diffusion, seen 

as a broadening of the initial implanted distribution, was very slow and was attributed to 

Si atoms that diffused interstitially. A concentration-dependent diffusivity with low 

solubility, which extended deeply into the sample, was quantitatively explained in terms 

of diffusion, via vacancies, of Si atoms on the Ga and As sub-lattices. Diffusion 

coefficients, together with the carrier concentration as a function of Si concentration, 

were obtained at various temperatures. The concentration-independent diffusion of Si was 

described by: 

D (cm 2 /s) = 1.23 x 10 -7 exp[-1.72(eV)/kT] 

The solid solubility of Si in GaAs was determined, and an exponential temperature 

dependence was observed. An estimate was made of the numbers of Si atoms which 

resided on the Ga and As sites, and the number of Si Ga + -Si As - pairs was deduced. The 

intrinsic diffusivities via neutral Ga vacancy complexes, triply negatively charged As 

vacancy complexes and triply negatively charged Ga vacancy complexes were: 

D (cm 2 /s) = 3.74 x 10 -3 exp[-2.60(eV)/kT] 

D (cm 2 /s) = 4.67 x 10 -5 exp[-2.74(eV)/kT] 

and 

D (cm 2 /s) = 5.92 x 10 -8 exp[-2.28(eV)/kT] 

respectively. 

T.Ahlgren, J.Likonen, J.Slotte, J.Räisänen, M.Rajatora, J.Keinonen: Physical Review B, 

1997, 56[8], 4597-603 

[446-157/159-326] 

GaAs: Si Diffusion 

The d-doped samples were grown by using molecular beam epitaxial methods, and were 

analyzed by using high-resolution secondary ion mass spectrometry. It was found that 

there was a marked difference between the profiles which were produced from samples 

that had been doped to a surface density of less than 1.3 x 10 13 /cm 2 (where all of the Si 

was incorporated on Ga sites), and highly-doped samples (where auto-compensation 

occurred). All of the samples were grown at a nominal temperature of 580C, and all of 

the doped planes showed some degree of broadening. A computer model for a 2-step 

diffusion process was developed, and this predicted a set of diffusion coefficients for 

lightly-doped samples. The diffusion coefficient which was associated with the postdeposition 

growth of these lightly-doped samples was about 4.2 x 10 -17 cm 2 /s. Because of 

their complicated profiles, more highly-doped samples were modelled by using a 

graphical technique. This revealed the presence of a much larger diffusion coefficient, 

which was tentatively attributed to the diffusion of Si as nearest-neighbor pairs. 

H.C.Nutt, R.S.Smith, M.Towers, P.K.Rees, D.J.James: Journal of Applied Physics, 1991, 

70[2], 821-6 

[446-93/94-010] 


Two Si-doped specimens, which had been grown by using molecular beam epitaxy 

techniques, were used to study Si diffusion at temperatures ranging from 700 to 950C. 

326


Each specimen structure consisted of well-characterized regions which ranged in type 

from undoped semi-insulating to heavily Si-doped (4 x 10 19 /cm 3 ). In one structure, the Si 

doping increased step-wise from the surface to the semi-insulating substrate. The other 

structure was grown in the reverse fashion, with the maximum Si doping situated near to 

the surface. Annealing was carried out after encapsulating the samples with a plasmaenhanced 

chemical vapor deposited nitride or oxide layer. Both structures exhibited 

almost identical diffusion behaviors which were best modelled by using an electrondependent 

diffusion model. A least-squares fit to both sets of data showed that the Si 

diffusivity could be described by: 

D(cm 2 /s) = 60.1 exp[-3.9(eV)/kT](n/n i ) 2 

This diffusion behavior was independent of the encapsulating material which was used, 

and of the proximity of the dopant to the surface region. The results indicated the 

existence of a Fermi level dependent diffusion behavior which was governed by 

compensating acceptor charged point defects, V Ga m- . These were a consequence of Si 

doping during molecular beam epitaxial growth. On the basis of the observed electron 

concentration dependence of the diffusivity, and assuming the operation of a simple Ga 

vacancy diffusion mechanism, these compensating vacancies were suggested to be at least 

doubly negatively charged. 

J.J.Murray, M.D.Deal, E.L.Allen, D.A.Stevenson, S.Nozaki: Journal of the 


[446-93/94-011] 



broadening of d-doped planes of Si. It was found that sharp spikes of Si could be obtained 

for sheet densities which were below 10 13 /cm 2 and for growth temperatures of 500C or 

less. At higher temperatures or densities, segregation or concentration-dependent rapid 

diffusion could occur; thus causing significant spreading even during growth. The codeposition 

of Si and Be markedly reduced this broadening. 



[446-91/92-001] 


The mechanism of Si diffusion was studied by using photoluminescence and secondary 

ion mass spectrometry, and transmission electron microscopy across the corner of a 

wedge-shaped sample. The diffusion source was a grown-in highly Si-doped layer. It was 

deduced that Frenkel defects (column-III vacancies and interstitials) were generated 

within the highly Si-doped region. The column-III interstitials rapidly diffused towards 

the surface, where they reacted with the column-III vacancies which were generated at the 

surface during annealing in a gaseous As ambient. This caused a supersaturation, of 

column-III vacancies in the Si-doped region, which drove Si diffusion. Annealing in 

vacuum reduced the supersaturation of column-III vacancies, and thus decreased Si 

diffusion. A predominant Si-donor plus column-III vacancy complex emission band was 

327


found in spectra from the Si-diffused region. The results supported the concept of a 

vacancy-assisted mechanism for Si diffusion and impurity-induced disordering. 

L.Pavesi, N.H.Ky, J.D.Ganière, F.K.Reinhart, N.Baba-Ali, I.Harrison, B.Tuck, M.Henini: 


[446-86/87-002] 

1.0E-11 

D (cm 2 /s) 

1.0E-12 

1.0E-13 

1.0E-14 

1.0E-15 

1.0E-16 

table 21 

table 22 

table 23 

table 24 

table 25 

1.0E-17 

1.0E-18 

1.0E-19 

7 8 9 10 11 12 13 14 

10 4 /T(K) 

Figure 6: Diffusivity of Si in GaAs 


The diffusion of Si was investigated by using sources from various tie-triangle regions in 

the Si-Ga-As phase diagram. This ensured constant chemical potentials for the 3 

components under isothermal conditions. The Si profiles were determined by using 

secondary ion mass spectrometry, and were found to be very different for the various 

types of source. The Ga-Si-GaAs tie-triangle source produced p-type Si doping with a 

concentration-independent diffusion coefficient. A neutral As or Ga vacancy was thought 

to be the predominant mobile defect under these conditions. The use of As-rich sources 

from 2 tie-triangle regions, or a Si-GaAs tie-line source, produced Si donor diffusion with 

a concentration-dependent diffusion behavior. The concentration-dependent diffusion 

328


coefficients of donor Si for As-rich source diffusion were related to the net ionized donor 

concentration, and exhibited 3 different regions. These were an intrinsic regime, an 

intermediate regime, and a saturation regime. It was proposed that Ga vacancies were 

responsible for donor diffusion; with a V 0 - 

Ga and/or V Ga mechanism for the intrinsic 

regime and a V - Ga -related mechanism for the extrinsic and saturation regimes. The use of 

a Si-GaAs tie-line source produced 2 branch-type profiles which were intermediate 

between the As-rich and Ga-rich diffusion cases. 

K.H.Lee, D.A.Stevenson, M.D.Deal: Journal of Applied Physics, 1990, 68[8], 4008-13 

[446-86/87-011] 

Table 21 

Diffusivity of Si in GaAs 


850 2.5 x 10 -15 

800 7.8 x 10 -16 

750 4.0 x 10 -16 

700 2.3 x 10 -16 

650 3.7 x 10 -17 


The effect of impurities upon the creation of Ga vacancies in Si-doped material, grown on 

a Be-doped epilayer by molecular beam epitaxy, was investigated by means of slow 

positron beam measurements. It was found that doping with Si enhanced the creation of 

Ga vacancies. The results supported a theoretical model in which the creation of Ga 

vacancies was explained in terms of a change in the-Fermi level position due to Si 

doping. It was also suggested that Si atoms diffused as a neutral complex, Si Ga -V Ga , rather 

than as Si Ga -Si As . A change in the S-parameter distribution at the interface between Sidoped 

and Be-doped regions was explained in terms of a so-called Be carry-forward 

effect which occurred during the growth of Si-doped GaAs on a Be-doped epilayer. 

J.L.Lee, L.Wei, S.Tanigawa, M.Kawabe: Journal of Applied Physics, 1990, 68[11], 5571- 

5 

[446-86/87-012] 





damage; the higher the concentration of extended defects, the slower being the diffusivity 





329




of diffusion. The diffusion was time-dependent. It was concluded that the results (table 







[446-84/85-016] 

Table 22 

Diffusivity of Implanted Si in GaAs 


5 x 10 15 950 2.5 x 10 -15 

1 x 10 14 900 2.6 x 10 -15 

1 x 10 15 900 2.0 x 10 -15 

1 x 10 15 900 1.2 x 10 -15 

1 x 10 15 900 8.3 x 10 -16 

1 x 10 15 900 6.2 x 10 -16 

1 x 10 14 900 4.4 x 10 -16 

1 x 10 14 900 3.0 x 10 -16 

1 x 10 14 900 2.2 x 10 -16 

1 x 10 15 900 1.3 x 10 -16 

5 x 10 15 850 1.4 x 10 -16 

1 x 10 15 850 4.2 x 10 -17 

5 x 10 15 750 7.5 x 10 -18 

5 x 10 14 750 3.7 x 10 -18 

5 x 10 14 750 1.8 x 10 -18 

1 x 10 15 750 1.0 x 10 -18 

1 x 10 14 850 1.3 x 10 -17 

1 x 10 14 900 1.4 x 10 -17 

1 x 10 14 900 2.9 x 10 -17 

1 x 10 15 900 4.0 x 10 -17 

1 x 10 14 900 5.4 x 10 -17 

1 x 10 14 950 5.0 x 10 -16 


The behavior of Si, after being initially d-doped into very pure gas-source molecular beam 

epitaxial GaAs layers, was studied by using capacitance-voltage profiling. A non-linear 

330


behavior of the diffusivity of Si as a function of the reciprocal temperature was observed. 

This was explained in terms of a 2-component Arrhenius dependence in which the 

activation energies changed by 1.5eV. When Si diffusion was governed by the lower 

activation energy, the impurity profile grew in width as a linear function of the annealing 

time. Deviations of the measured Si diffusivity from classical impurity diffusion behavior 

were attributed to the existence of a non-equilibrium concentration of vacancies which 

was generated at the d-source position during annealing. 

J.E.Cunningham, T.H.Chui, W.Jan, T.Y.Kuo: Applied Physics Letters, 1991, 59[12], 

1452-4 

[446-84/85-016] 


Samples which had been implanted with 220keV Si ions to doses ranging from 3 x 10 13 to 

10 15 /cm 2 , and annealed at 850C, were studied. By using transmission electron 

microscopy, voids were observed in samples with implanted doses of more than 3 x 

10 14 /cm 2 ; after annealing times which were as short as 5s. In the same region where voids 

were found, capacitance-voltage measurements revealed abnormally low electron 

concentrations. Also in the same region, secondary ion mass spectrometry measurements 

detected anomalies in the Si concentration profiles. It was therefore deduced that Si 

redistribution had occurred. At high Si doses, the onset of void formation, the abnormally 

low electron concentration, and the Si accumulation anomaly were concurrent. On the 

basis of the results, it was concluded that voids inhibited electrical activity and led to the 

Si diffusion anomaly. 

S.Chen, S.T.Lee, G.Braunstein, K.Y.Ko, L.R.Zheng, T.Y.Tan: Japanese Journal of 

Applied Physics, 1990, 29[11], L1950-3 

[446-81/82-010] 


The growth of Si-doped liquid-encapsulated Czochralski material exhibited a significant 

deviation, in Si incorporation, from that which was predicted by a classical segregation 

model. It was usually expected that, for a given impurity segregation coefficient, dopant 

incorporation throughout the crystal could be calculated with fairly good accuracy. 

Profiles for Te-doped liquid-encapsulated Czochralski material gave a close 

approximation to this classical model. However, in the case of Si-doped material, the 

dopant distribution in the crystals deviated significantly from the segregation model. In 

some cases, this deviation amounted to several orders of magnitude. The degree of 

deviation was found to depend upon the growth conditions. The present work was 

performed in order to understand the source of the deviation from the model and to permit 

accurate account to be taken of Si incorporation. Therefore, Si-doped crystals were grown 

by using a low-pressure liquid-encapsulated Czochralski technique and were intended to 

be doped with between 10 16 and 10 18 /cm 3 . The actual dopant incorporation was 

significantly lower than that predicted by the segregation model, and the axial dopant 

variation was much too flat. Also, in spite of the use of intentional Si doping, a significant 

number of the crystals was semi-insulating. An analysis of the experimental data showed 

331


that this doping behavior could not be explained by assuming the existence of an 

unknown compensating impurity. A different model was developed which was consistent 

with all of the observations and which provided an accurate account of Si incorporation. 

An understanding of the Si incorporation anomaly permitted successful process changes 

and improvements to be made. This model explained the observed anomaly by taking 

account of Si diffusion between the GaAs melt and the B 2 O 3 encapsulant, as well as of 

permanent Si trapping in the B 2 O 3 . When the interaction of Si with B 2 O 3 was taken into 

account, the Si incorporation obeyed the segregation model and the anomaly vanished. 

A.Flat: Journal of Crystal Growth, 1991, 109[1-4], 224-7 

[446-81/82-011] 

Table 23 



1000 1.5 x 10 -12 

950 4.1 x 10 -13 

900 3.1 x 10 -14 

800 5.7 x 10 -15 

705 1.0 x 10 -16 


The lateral variation in the emission energy of GaAs which had been masklessly grown 

on V-grooved Si was studied by using cathodoluminescence wavelength imaging. This 

new experimental approach permitted, for the first time, the direct visualization and 

quantification of the extreme homogeneity of this novel growth mode and of the lateral 

variations of Si impurity incorporation into such semiconductor microstructures. It was 

thus an important method for characterizing micro-patterned opto-electronic monolithic 

integrated circuits. 

M.Grundmann, J.Christen, D.Bimberg, A.Hashimoto, T.Fukunaga, N.Watanabe: Applied 


[446-81/82-012] 


Capacitance-voltage methods were used to profile d-doped GaAs layers which had been 

grown on Si substrates via metalorganic chemical vapor deposition. It was found that 

there was a close correlation between dislocation densities, in the epitaxial layers, and the 

associated diffusion coefficients. After rapid thermal annealing (800-1000C, 7s), the 

diffusion data (table 23) could be described by: 

D(cm 2 /s) = 30 exp[-3.4(eV)/kT] 

332


for a relatively thick buffer layer of 0.0033mm. It was concluded that the dislocationenhanced 

diffusion of Si impurities was appreciable, and that the inclusion of an 

0.003mm buffer layer was insufficient to prevent the diffusion of impurities. 

Y.Kim, M.S.Kim, S.K.Min, C.Lee: Journal of Applied Physics, 1991, 69[3], 1355-8 

[446-78/79-014] 


High depth-resolution secondary ion mass spectrometry profiling was used to investigate 

the broadening of d-doped planes of Si plus Be in material which had been prepared by 

using molecular beam epitaxial methods. It was confirmed that concentration-dependent 

diffusion was the predominant broadening process for Be at growth temperatures of less 

than 600C. By incorporating Si atoms into the same plane, it was shown that the 

broadening could be completely inhibited. This suggested that the rapid diffusion process 

resulted from mutual repulsion between the Be Ga - ions, and was prevented by the reverse 

field which arose from Si Ga + ions or by the formation of low-mobility Si Ga + -Be Ga 

- 

complexes. The rapid diffusion of Si as Si Ga -Si As pairs was also reduced. The latter was 

attributed to a Fermi-level effect, with compensation by Be tending to reduce the 

probability of Si As formation. The surface segregation of Si was unaffected, whereas that 

of Be was reduced. This indicated that the surface fields which existed during growth 

contributed to the behavior of Be, but not to that of Si. 

J.J.Harris, J.B.Clegg, R.B.Beall, J.Castagné: Semiconductor Science and Technology, 

1990, 5[7], 785-8 

[446-76/77-007] 


A quantitative model was proposed for the behavior of Si. The model took account of the 

fact that Si was an amphoteric impurity which acted as a shallow donor when it occupied 

the Ga site (Si Ga + ) and as a shallow acceptor when it occupied the As site (Si As - ). Both 

Si Ga + and Si As - existed at high concentrations. The amphoteric behavior of Si could be 

viewed as being an effect of the Fermi level. It was assumed that Si Ga + and Si As - diffused 

on the Ga and As sub-lattices, respectively. Thus, the diffusion of Si Ga + and Si As - was 

governed by the variously charged Ga vacancies and As vacancies (or self-interstitials) 

respectively. The experimentally observed Si diffusivity was concentration dependent, 

and this was attributed to the amphoteric nature of Si as well as to another effect of the 

Fermi level. The latter involved its influence upon the concentrations of charged pointdefect 

species. Satisfactory quantitative descriptions of available experimental data were 

obtained. An analysis of results on the diffusion of Si into a Sn-doped GaAs substrate 

suggested that a previously proposed Si-pair diffusion model was unfavorable. 

S.Yu, U.M.Gösele, T.Y.Tan: Journal of Applied Physics, 1989, 66[7], 2952-61 

[446-74-012] 

333



A model for Si diffusion was developed which was based upon the formation of Si Ga + - 

V Ga - pairs. By using recent data on the diffusivity of Ga vacancies, it was shown that the 

pair diffusion coefficient was approximately equal to 75% of the former value. The 

predictions of the model were found to be in good agreement with Si diffusion data. 

K.B.Kahen: Journal of Applied Physics, 1989, 66[12], 6176-8 

[446-74-012] 


The diffusion of Si was studied after implanting Si into un-doped, Se-, Si- or Zn-doped 

material. The diffusion of grown-in Si in epitaxial layer structures was also studied. No 

diffusion was detected in un-doped or Zn-doped material, while a moderate amount of 

diffusion was detected in Si-doped samples. A significant amount of diffusion occurred in 

Se-doped material and in non-implanted Si-doped epitaxial samples. The results indicated 

that diffusion was controlled by a Fermi-level mechanism which probably involved 

ionized Ga vacancies, and that implantation damage inhibited diffusion by keeping the 

electron concentration and/or the ionized Ga vacancy concentration at a low level. 

J.J.Murray, M.D.Deal, D.A.Stevenson: Applied Physics Letters, 1990, 56[5], 472-4 

[446-74-012] 


The effect of the substrate temperature, during molecular beam epitaxial growth, upon the 

migration of Si atoms in d-doped or planar-doped GaAs was investigated by using 

secondary ion mass spectrometry. The results for d-doped GaAs revealed a measurable 

spread of Si, which increased by about 8nm when the substrate temperature was increased 

from 580 to 640C. For substrate temperatures below 580C, the measured width of the Si 

profile was limited by the resolution of the secondary ion mass spectrometer. 

Magnetotransport measurements were also performed in order to determine dopant 

spreading. The Si migration which was measured by means of secondary ion mass 

spectrometry was in qualitative agreement with the transport results. However, the 

secondary ion mass spectrometry data indicated larger Si areal densities. Two 

mechanisms, auto-compensation and electron localization by a DX center, were believed 

to be responsible for the latter observations. 

A.M.Lanzillotto, M.Santos, M.Shayegan: Applied Physics Letters, 1989, 55[14], 1445-7 

[446-72/73-002] 


Electron-beam deposited films of phosphosilicate were used for the encapsulation of Siimplanted 

material during post-implantation annealing. Depth profiles which were in 

excellent agreement with the Lindhard-Scharff-Schiott curves were obtained by using 

100nm thick films and by annealing for 0.5h at 850C. The diffusion coefficient of 

334


implanted Si was found to be an order of magnitude smaller for phosphosilicate films 

than for conventional plasma-assisted chemical vapor deposited SiO 2 films. 

S.Singh, F.Baiocchi, A.D.Butherus, W.H.Grodkiewicz, B.Schwartz, L.G.Van Uitert, 

L.Yesis, G.J.Zydzik: Journal of Applied Physics, 1988, 64[8], 4194-8 

[446-72/73-011] 


Layers were grown, by molecular beam epitaxy, at a substrate temperature of 520C. The 

layers contained three d-doped planes, with Si concentrations of 4 x 10 12 , 10 13 or 4 x 

10 13 /cm 3 , and were annealed at temperatures of up to 648C. Secondary ion mass 

spectroscopy and capacitance-voltage methods were then used to monitor broadening of 

the profiles. It was found that the most lightly doped plane gave a near-Gaussian profile, 

and the diffusion coefficient was comparable with published data on the simple diffusion 

of isolated Si Ga atoms. The more highly doped planes exhibited complex profiles with 2 

components; in which some of the atoms were confined to the original plane, while there 

was an essentially square-shaped profile of fast-diffusing atoms. A comparison of the 2 

types of experimental data suggested that the formation of Si islands took place during 

deposition of the d-doped plane. This gave rise to electrically inactive atoms which could 

then diffuse into the surrounding material during heat treatment. 

R.B.Beall, J.B.Clegg, J.Castagné, J.J.Harris, R.Murray, R.C.Newman: Semiconductor 

Science and Technology, 1989, 4[12], 1171-5 

[446-72/73-011] 


Quantum oscillations in the magnetoresistance of Si-doped material were analyzed in 

order to obtain the electron densities of the electrical sub-bands. These densities were 

compared with the results for self-consistently calculated sub-band structures of d-doped 

material in which the spread of Si dopant atoms in the growth direction was a fitting 

parameter. The results indicated that there was a negligible spread in structures which 

were grown at substrate temperatures of up to 530C. In structures which were grown at 

higher substrate temperatures, there was a measurable spread. This increased with 

increasing substrate temperature. At a substrate temperature of 640C, the Si spread was 

about 22nm. An examination of the three-dimensional Si densities in these layers 

indicated that the predominant mechanism of Si spreading at substrate temperatures 

greater than 600C was the Si migration which was necessary in order to satisfy the solid 

solubility limit. 

M.Santos, T.Sajoto, A.Zrenner, M.Shayegan: Applied Physics Letters, 1988, 53[25], 

2504-6 

[446-64/65-163] 



disordering. The mechanism of enhanced superlattice disordering due to Si doping was a 

Fermi-level effect which increased the concentrations of charged point defects. The 

335


diffusion of Si appeared to be governed by Ga vacancies, and was well described by 

current Si pair diffusion models. It was concluded that dislocations in this material and in 

other III-V compounds were only moderately efficient sinks or sources for point defects. 


[446-62/63-208] 

Table 24 



995 7.8 x 10 -14 

940 3.4 x 10 -14 

890 1.2 x 10 -14 

800 1.3 x 10 -15 

685 9.5 x 10 -17 

590 1.2 x 10 -17 


The migration of atomic Si in d-doped material was studied by means of capacitancevoltage 

measurements and rapid thermal annealing. It was shown that these methods 

could detect diffusion which occurred at a length scale as small as 1nm. The capacitancevoltage 

profile widths broadened from less than 4nm, to 13.7nm, upon annealing (1000C, 

5s). It was found that the results could be described by: 

D(cm 2 /s) = 0.0004 exp[-2.45(eV)/kT] 

E.F.Schubert, T.H.Chiu, J.E.Cunningham, B.Tell, J.B.Stark: Journal of Electronic 

Materials, 1988, 17[6], 527-31 

[446-62/63-210] 


Impurities, with an initially Dirac d-like distribution profile, were diffused into GaAs by 

using rapid thermal annealing. The diffusion of atomic Si was monitored via a novel 

method which involved comparing the experimental capacitance-voltage profiles with 

predicted self-consistent profiles. The capacitance-voltage profiles broadened during 

rapid thermal annealing (1000C, 5s). It was found that the diffusion data (table 24) could 

be described by: 

D(cm 2 /s) = 0.0004 exp[-2.45(eV)/kT] 

The diffusivity values were 2 orders of magnitude smaller than those for Si-pair diffusion 

in GaAs. 

E.F.Schubert, J.B.Stark, T.H.Chiu, B.Tell: Applied Physics Letters, 1988, 53[4], 293-5 

[446-62/63-210] 

336


Table 25 



960 7.5 x 10 -15 

955 4.7 x 10 -15 

895 4.4 x 10 -15 

900 2.0 x 10 -15 

795 2.0 x 10 -15 

790 1.1 x 10 -15 

695 1.0 x 10 -15 

695 5.1 x 10 -16 

645 1.9 x 10 -16 

590 7.8 x 10 -17 

535 4.9 x 10 -17 

605 3.7 x 10 -17 

485 2.0 x 10 -17 

590 1.9 x 10 -17 

555 8.2 x 10 -18 

515 1.0 x 10 -18 


High-quality samples were prepared by means of chemical beam epitaxy, at a substrate 

temperature of 500C, by using triethylgallium and arsine. Capacitance-voltage 

measurements of Si-doped material revealed profile widths of 2.2nm at 300K and 1.8nm 

at 77K. This indicated that a high degree of Si spatial localization had been achieved. 

Subsequent annealing treatments showed that appreciable Si segregation and diffusion 

occurred at a growth temperature of about 600C. The capacitance-voltage widths of 

annealed doped structures provided a good estimate of the diffusion coefficient of Si in 

GaAs. 

T.H.Chiu, J.E.Cunningham, B.Tell, E.F.Schubert: Journal of Applied Physics, 1988, 

64[3], 1578-80 

[446-61-067] 


Planar confinement, diffusion, and surface segregation results were reported for Si which 

had been d-doped into GaAs. In the case of gas-source molecular beam epitaxy, the Si 

diffusion (table 25) as a function of the reciprocal annealing temperature exhibited an 

337


unique 2-component Arrhenius form in which the activation energies differed by the 

fundamental GaAs band-gap energy of 1.5eV. 

J.E.Cunningham, T.H.Chiu, B.Tell, W.Jan: Journal of Vacuum Science and Technology 

B, 1990, 8[2], 157-9 

[446-74-012] 


Secondary ion mass spectroscopy and carrier concentration measurements were used to 

characterize Si diffusion into GaAs wafers which contained 2 fundamentally different 

types of donor. These were column-IV donors (Si, Sn) and column-VI donors (Se, Te). A 

decrease in the Si diffusion rate was found in material which contained column-VI donors 

as compared with that which contained column-IV donors. This trend was consistent with 

a model in which Si diffused as donor Ga-vacancy complexes. The decrease in the Si 

diffusion coefficient was attributed to the greater binding energy of column-VI donor Gavacancy 

nearest-neighbor complexes. This reduced the numbers of free Ga vacancies 

which were available to complex with the Si. 

D.G.Deppe, N.Holonyak, J.E.Baker: Applied Physics Letters, 1988, 52[2], 129-31 

[446-60-003] 



strongly affected Si. An increase in the As 4 over-pressure resulted in an increase 

in the diffusion depth. No band-edge exciton was observed during absorption on material 

that was diffused with Si, in spite of the high degree of compensation. 



[446-60-004] 


The diffusion profiles of buried Si dopant which had been implanted by using a focussed 

ion beam were determined after annealing. The diffusion coefficient of the Si was 

determined by fitting the results of computer calculations. Concentration-dependent 

diffusion of Si which had been introduced by using a molecular beam was observed. 

However, the diffusion coefficient of Si which had been introduced by using a focussed 

ion beam was undetectably small when compared with that of Be at 850C. 

T.Morita, J.Kobayashi, T.Takamori, A.Takamori, E.Miyauchi, H.Hashimoto: Japanese 


[446-55/56-005] 


Junction-depth measurements, performed using scanning electron microscopy and 

secondary ion mass spectroscopy, were used to characterize Si diffusion in GaAs crystals 

which contained various amounts of Zn background doping. The Zn concentration was 

338


found to control Si diffusion. This behavior was attributed to a shift in the Fermi level 

with increasing n-type doping. Also, the electric field which was due to the p-n junction 

that formed at the Si diffusion front had a large effect upon the Zn background doping 

profile. 

D.G.Deppe, N.Holonyak, F.A.Kish, J.E.Baker: Applied Physics Letters, 1987, 50[15], 

998-1000 

[446-51/52-116] 






profile. 



[446-121/122-045] 


The reported anomalous so-called up-hill diffusion behavior of implanted Si in this 

material was simulated. The up-hill diffusion had been found to be implantation-dose 

dependent. No up-hill diffusion was observed below a threshold dose of 3 x 10 14 /cm 2 . It 

was suggested here that the anomalous behavior was due to the formation of vacancy 

clusters when the implantation dose was sufficiently high. Atoms of Si were assumed to 

migrate through the Ga vacancies, which were comparatively mobile. On the other hand, 

the immobile clusters collected Si atoms around the vacancy peak; thus resulting in uphill 

diffusion. A dip in the carrier concentration profile also occurred at high implantation 

doses. The occurrence of this dip was attributed to the effect of the trapping of Si atoms 

into clusters where the former were not electrically activated. 

V.C.Lo, J.Z.Sun: Modelling and Simulation in Materials Science and Engineering, 1996, 

4[6], 613-21 

[446-157/159-339] 


The diffusion of Si was carried out (900C, 5h, As pressure), from a 50nm sputtered film 

and into undoped semi-insulating material or Te-doped or Zn-doped liquid-encapsulated 

Czochralski material. Secondary ion mass spectroscopy and spreading resistance 

techniques were used to characterize the Si in-diffusion profiles. Lattice defects in highly 

Si-doped diffusion regions were studied as a function of the post-diffusion heat treatment 

(700C, 0.25h; 1000C, 0.5h) via the transmission electron microscopy of plan-view and 

cross-sectional samples. Two types of defect were observed in the diffused region. These 

were perfect prismatic loops of interstitial type on {110} planes, and Frank faulted loops 

on {111}; again of interstitial type. Defect formation, and the role of Si in defect 

generation, were explained in terms of a negative temperature dependence of the thermal 

339


equilibrium concentrations of V Ga 3- ; which were assumed to mediate Si diffusion under 

highly n-doped conditions. Cathodoluminescence spectra at 4 and 77K were obtained 

from the diffusion layer. It was found that the Si diffusion affected the band-gap 

luminescence and generated 2 deep-level emission bands in the 0.9 to 1.3eV spectral 

region. It was suggested that these deep levels were associated with diffusion-induced 

defects and defect complexes. 

L.Herzog, U.Egger, O.Breitenstein, H.G.Hettwer: Materials Science and Engineering B, 

1995, 30[1], 43-53 

[446-134/135-125] 


The migration of Si acceptors (Si on As sites) in δ-doped GaAs which had been grown 

onto GaAs(111)A, was investigated by means of secondary ion mass spectrometry. It was 

found that the diffusion parameters for GaAs(111)A differed from those for GaAs(001). 

The diffusion coefficient in GaAs(111)A was smaller than that in GaAs(001), and the 

activation energy in GaAs(111)A was higher than that in GaAs(001). The diffusion 

mechanism of Si in GaAs(111)A was investigated by means of photoluminescence and it 

was found that, in p-type layers, Si-donors (Si on Ga sites) diffused easily to As sites. The 

data on Si acceptor diffusion could be described by: 

D(cm 2 /s) = 0.0114 exp[-2.74(eV)/kT] 

The results indicated that Si-acceptors were more stable than Si donors. 

M.Hirai, H.Ohnishi, K.Fujita, P.Vaccaro, T.Watanabe: Journal of Crystal Growth, 1995, 

150[1-4], 209-13 

[446-127/128-120] 


An investigation was made of the fast diffusion of Si from deposited surface layers when 

oxidized in an Ar/H 2 O ambient. This revealed the presence of excess As which was 

formed by the oxidation of Ga which originated from the substrate, and was used to 

explain the enhanced diffusion of Si into the substrate. The formation of SiO 2 on the 

surface during oxidation prevented the loss of the excess As, which then accumulated in 

the remaining Si film. The use of higher H 2 O partial pressures during oxidation produced 

higher As/Si ratios; thus resulting in an increase in Si diffusion depth and concentration. 

However, the n-type carrier concentration decreased with increasing As/Si ratios in the 

remaining Si layer. A second non-oxidizing annealing treatment, of samples from which 

the SiO 2 and Si layer had been removed, had differing effects upon the carrier 

concentration; depending upon whether the As was free to escape from the substrate. The 

results indicated that excess As-related defects, such as Ga vacancies, were probably 

responsible for n-type compensation in fast diffused samples. 

R.C.Keller, C.R.Helms: Applied Physics Letters, 1995, 67[3], 398-400 

[446-123/124-162] 

340



The diffusion of Si was studied by using secondary ion mass spectroscopic and 

transmission electron microscopic methods after implanting it using energies ranging 

from 20 to 200keV and doses ranging from 10 13 to 10 14 /cm 2 , followed by furnace 

annealing. It was found that little or no diffusion occurred after implantation at energies 

greater than 100keV. At energies of less than 100keV, there was usually appreciable 

dopant redistribution; regardless of the peak implant concentration. Both concentrationdependent 

and concentration-independent diffusion was observed. The dislocation loop 

density varied inversely with the amount of diffusion as a function of implantation 

energy. A standard Monte Carlo computer program was able to predict the trends in the 

implant energy dependence of diffusion by considering the excess point defect content 

which was produced by implantation. The effect of this excess defect dose and of surface 

effects upon Si diffusion was consistent with vacancy-assisted diffusion. 

C.C.Lee, M.D.Deal, K.S.Jones, H.G.Robinson, J.C.Bravman: Journal of the 


[446-119/120-192] 


Implantation of Si into (x11)A-oriented semi-insulating GaAs substrates (where x took 

values of up to 4) was carried out. For comparison, (110)- and (100)-oriented substrates 

were also implanted. No in-diffusion of Si was observed after annealing substrates with 

any orientation. A similar behavior was observed for Si implants in GaAs and for Si/B coimplants. 



[446-117/118-166] 


Dopant diffusion and defect formation were studied as a function of implantation 

temperature in Si-implanted material. It was found that the diffusion of Si during postimplantation 

annealing decreased by a factor of 2.5 as the implantation temperature was 

increased from -2 to 40C. Within the same temperature range, the maximum depth and 

density of extrinsic dislocation loops increased by factors of 3 and 4, respectively. 

Rutherford back-scattering channelling measurements indicated that Si-implanted GaAs 

underwent an amorphous to crystalline transition at Si implantation temperatures of 

between -51 and 40C. A unified explanation was proposed, for the effects of implantation 

temperature upon diffusion and dislocation formation, which was based upon known 

differences in sputter yield between crystalline and amorphous semiconductors. The 

model assumed that the sputter yield was enhanced by amorphization at lower 

temperatures; which increased the excess vacancy concentration. Estimates of the latter 

341


were obtained by simulating the diffusion profiles, and were found to be quantitatively 

consistent with sputter yield enhancement. 

H.G.Robinson, T.E.Haynes, E.L.Allen, C.C.Lee, M.D.Deal, K.S.Jones: Journal of 


[446-117/118-166] 


A study was made of the diffusion of Si in δ-doped layers of (111)A- or (100)-type, and 

of the evaporation of As atoms from the surfaces. It was found that the diffusion of 

dopants in (111)A layers was slower than in (100), regardless of the presence of As 

vacancies. On the other hand, diffusion in (100) layers was enhanced by the presence of 

As vacancies. It was noted that As atoms on the (111)A surface did not evaporate easily, 

as compared with those on the (100) surface. 

A.Shinoda, T.Yamamoto, M.Inai, T.Takebe, T.Watanabe: Japanese Journal of Applied 

Physics, 1993, 32[2-10A], L1374-6 

[446-115/116-116] 


First-principles calculations were presented for the vacancy-mediated diffusion of Si. It 

was shown that a DX-like mechanism facilitated the migration of lattice-site atoms into 

the interstitial region, and that the dangling bonds of a second-nearest neighbor vacancy 

assisted migration through the interstitial region. Due to these 2 mechanisms, vacancyassisted 

diffusion of Si occurred with a low-energy barrier. 

J.Dabrowski, J.E.Northrup: Physical Review B, 1994, 49[20], 14286-9 

[446-115/116-117] 


A study was made of molecular beam epitaxially grown samples which were δ-doped 

with Si and Al layers. Long-term diffusion annealing was carried out at temperatures 

ranging from 550 to 800C, and the concentration profiles were determined by means of 

secondary ion mass spectrometry. It was found that the results could be described by the 

expression, 

D(cm 2 /s) = 7.9 exp[-2.25(eV)/kT] 

The Si diffusion coefficients which were obtained were in good agreement with data 

which had been obtained by using rapid thermal annealing, capacitance-voltage profiling, 

and sandwiched diffusion sources. They differed from earlier measurements which had 

been based upon the diffusion of implanted dopants that were much more widely spread. 

It was concluded that the more accurate data which resulted from δ-doping showed that 

the diffusion coefficient was an intrinsic parameter, provided that the amount of dopant 

and the dislocation density were kept sufficiently small. 

F.Bénière, R.Chaplain, M.Gauneau, V.Reddy, A.Régrény: Journal de Physique III, 1993, 

3[12], 2165-71 

[446-115/116-119] 

342



Various mechanisms of Si diffusion in GaAs were investigated by using first-principles 

molecular dynamics methods. It was found that the predominant mechanism involved the 

motion of negatively charged Si III -V III pairs via second-nearest neighbor jumps. This 

mechanism explained the ability of Si to disorder superlattices (regardless of whether it 

was introduced during growth or was in-diffused later), and the suppression of 

interdiffusion by compensation doping. The calculated activation energies were in very 

good agreement with experimental data. 

B.Chen, Q.M.Zhang, J.Bernholc: Physical Review B, 1994, 49[4], 2985-8 

[446-115/116-119] 


It was recalled that implantation damage was believed to play a significant role in dopant 

diffusion. An attempt was made here to modify the point defect damage profile of 40keV 

29 Si implants by chemically etching away the top 10nm of the sample after implantation. 

No Si diffusion was observed in these samples after annealing, whereas significant Si 

redistribution occurred in a similar sample which had received no post-implantation 

etching. The results of TRIM simulations predicted an excess Ga vacancy surface layer, 

and the presence of excess Ga interstitials deeper within the sample. It was thought that, 

by removing the vacancy-rich surface layer, the etch altered the implant damage profile 

and thus the diffusion behavior of Si. The surface effects of etching which were related to 

Si diffusion were shown to be consistent with a vacancy-assisted diffusion mechanism. 

There was some evidence that this model might be applicable to B implants in Si. 


[446-115/116-119] 


A study was made of the transport properties of the 2-dimensional electron gas in 

annealed Si δ-doped structures. The diffusion rate of Si atoms was deduced from the 

temperature dependence of the sub-band population. In samples with large selfcompensation, 

it was found that the electron density increased after annealing at 

temperatures below 800C. After annealing at temperatures above 800C, there was a 

reduction in the electron concentration of the 2-dimensional electron gas. The results 

showed that, after annealing, the quantum mobility in the lowest sub-band increased 

slightly, while the quantum mobility in the higher sub-bands markedly decreased. 

P.M.Koenraad, I.Bársony, A.F.W.Van der Stadt, J.A.A.J.Perenboom, J.H.Wolter: 

Materials Science Forum, 1994, 143-147, 663-8 

[446-113/114-012] 


The diffusion of Si out of δ-planes in GaAs was investigated by means of high-resolution 

X-ray diffractometry, infra-red absorption localized vibrational mode spectroscopy and 

secondary ion mass spectroscopy. In the case of a Si δ-doped sample which had been 

343


grown at 400C with a Si areal concentration of 3.4 x 10 14 /cm 2 , the Si was confined to a 

layer which was no more than 0.5nm in thickness. After post-growth annealing (600C, 

3h), 22.6% of the Si remained on the δ-planes, while the remainder had diffused 

homogeneously throughout the epilayer to give a Si concentration of 2.1 x 10 19 /cm 3 . 

Localized vibrational mode spectroscopy indicated that these Si atoms were located 

mainly on Ga sites (Si Ga ). The Si atoms were also found to occupy As sites, or were 

present as Si Ga -Si As pairs, Si-X and Si Ga -V Ga complexes. At 950C, all of the Si had 

diffused away from the δ-planes to form precipitates and dislocation loops near to the 

surface. 

L.Hart, P.F.Fewster, M.J.Ashwin, M.R.Fahy, R.C.Newman: Materials Science Forum, 

1994, 143-147, 647-52 

[446-113/114-013] 


It was recalled that, in n-type material, the diffusion of atoms which resided on the Ga 

sub-lattice occurred mainly via Ga vacancies. In order to elucidate the microscopic details 

and energetics of these processes, local density approximation estimates were made of the 

total energies, electronic structures, and relaxed positions of atoms in microscopic 

configurations which were relevant to the diffusion of Si donors. It was found that a DXlike 

mechanism facilitated the migration of lattice site atoms into the interstitial region. A 

bond-passing mechanism was also identified, in which a second-nearest neighbor vacancy 

assisted migration through the interstitial region. Due to these 2 mechanisms, vacancyassisted 

Si diffusion in n-type material was characterized by an energy barrier of about 

1.0eV. 

J.Dabrowski, J.Northrup: Materials Science Forum, 1994, 143-147, 1263-8 

[446-113/114-013] 


The thermal equilibrium concentrations of the various negatively charged Ga vacancy 

species were calculated. The triply negatively charged Ga vacancy, V Ga 3- , was studied in 

particular because it dominated Ga self-diffusion and Ga/Al interdiffusion under intrinsic 

and n-doping conditions, as well as the diffusion of Si donor atoms which occupied Ga 

sites. 

T.Y.Tan, H.M.You, U.M.Gösele: Applied Physics A, 1993, 56[3], 249-58 

[446-111/112-050] 


Samples were doped with Si to a concentration of about 2.7 x 10 18 /cm 3 and were annealed 

(800 to 1000C, 3 to 20h) under As-rich or As-poor conditions. The Si out-diffusion was 

measured by using the capacitance-voltage method and an electrochemical profiler. It was 

found that the Si diffusivity exhibited a marked dependence upon the As 4 vapor phase 

pressure and upon the electron concentration. When reduced to intrinsic conditions, 

activation enthalpies of 3.91 and 4.19eV were obtained for As-rich and As-poor annealing 

344


conditions, respectively. On the basis of these results, it was concluded that Si outdiffusion 

was governed by the triply negatively charged vacancies, V Ga 3- . 

H.M.You, U.M.Gösele, T.Y.Tan: Materials Science Forum, 1993, 117-118, 399-404 

[446-111/112-051] 


Samples of n-type and p-type d-doped material, grown by molecular beam epitaxy and 

with quite high doses of Si, were investigated by means of transmission electron 

microscopy. The magnitude of the doses ranged from half a monolayer to 2 monolayers. 

The microscopic structures of the d-doped regions and of the adjacent epilayers were 

observed directly. The effect of impurity spreading upon the hetero-interfaces and 

superlattices was studied, and it was found that the Si atoms in Si d-doped samples were 

confined to within a few atomic layers. Stacking faults were found in d-doped samples 

when they were grown at low temperatures. Their presence was attributed to local strains 

that were caused by heavy doping. 

D.G.Liu, J.C.Fan, C.P.Lee, K.H.Chang, D.C.Liou: Journal of Applied Physics, 1993, 

73[2], 608-14 

[446-106/107-034] 


The lateral diffusion of sources during the selective growth of metalorganic vapor-phase 

epitaxial Si-doped layers was analyzed. The diffusion lengths of Si species were deduced 

from the carrier concentration profiles which were measured by using Raman 

spectroscopy and thickness profiling. On the basis of these diffusion lengths, it was 

speculated that the effective diffusion material was silyl arsine. It was suggested that there 

was no difference between arsine and tertiary butyl arsine, as diffusion sources. 

N.Hara, K.Shiina, T.Ohori, K.Kasai, J.Komeno: Journal of Applied Physics, 1993, 74[2], 

1327-30 

[446-106/107-036] 


The formation energy of Si donors, acceptors, and defect complexes were calculated. The 

equilibrium concentrations of native defects and Si-defect complexes were deduced from 

these energies, as was the total solubility of Si. The calculated equilibrium solubility limit 

for Si was in good agreement with experimental data. The (Si Ga -V Ga ) 2- complex occurred 

at relatively high concentrations under As-rich conditions, and could therefore mediate Si 

and Ga diffusion. It was concluded that the donor-vacancy complex was an important 

compensation mechanism in heavily doped GaAs. 

J.E.Northrup, S.B.Zhang: Physical Review B, 1993, 47[11], 6791-4 

[446-106/107-036] 

345



An experimental study was made of Si out-diffusion from GaAs, by using pre-doped 

samples. The results showed that the Si diffusivity depended upon the As 4 vapor-phase 

pressure in the ambient, and upon the electron concentration in the crystal. It was 

concluded that, in GaAs, diffusion of the Si donor species which occupied Ga sites, Si + Ga , 

was governed by the triply negatively charged Ga vacancies, V 3- Ga . However, the present 

V 3- Ga -dominated Si + Ga out-diffusivities were larger, by many orders of magnitude, than 

those which were obtained under Si in-diffusion conditions. A tentative explanation of 

this large difference was given in terms of an undersaturation of V 3- Ga in intrinsic material 

3- 

during in-diffusion experiments, and of a supersaturation of V Ga which developed 

during the out-diffusion of Si from n-type Si-doped material. 

H.M.You, U.M.Gösele, T.Y.Tan: Journal of Applied Physics, 1993, 73[11], 7207-16 

[446-106/107-037] 


Heavily Si-doped (5 x 10 19 /cm 3 ) low-temperature GaAs, sandwiched between undoped 

low-temperature GaAs layers, was grown by using molecular beam epitaxy and was 

annealed at up to 900C. Transmission electron microscopy showed that, within the first 

few minutes of annealing, an accumulation of As precipitates formed near to each 

doped/undoped low-temperature interface. During further annealing, Si segregation to As 

precipitates was detected, using secondary ion mass spectroscopy, in the form of deltalike 

peaks at the As precipitate accumulations. It was found that the Si diffusion 

coefficient was initially independent of concentration, at a value of 2.5 x 10 -13 cm 2 /s, and 

was comparable to diffusion under intrinsic conditions in As-rich material when grown at 

normal temperatures. During annealing for 1h, the Si concentration in the As precipitates 

reached 2.5 x 10 20 /cm 3 . 

K.L.Kavanagh, J.C.P.Chang, P.D.Kirchner, A.C.Warren, J.M.Woodall: Applied Physics 

Letters, 1993, 62[3], 286-8 

[446-106/107-038] 


The diffusion of Si from a novel diffusion source, consisting of an undoped SiO x /SiN 

double-layered film, was studied by rapid thermal annealing at 860 to 940C. The 

characteristics of the Si-diffused layers were investigated by using secondary ion mass 

spectrometric, capacitance-voltage, and Hall methods. The carrier profiles exceeded 2 x 

10 18 /cm 3 , and featured an abrupt diffusion front. A maximum electron concentration of 6 

x 10 18 /cm 3 was obtained at 940C. The diffused Si profiles were consistent with the 

operation of Si Ga + -V Ga - pair diffusion. 

S.Matsushita, S.Terada, E.Fujii, Y.Harada: Applied Physics Letters, 1993, 63[2], 225-7 

[446-106/107-038] 

346


GaAs/AlAs: Si Diffusion 

Various mechanisms of Si-induced interdiffusion in GaAs/AlAs superlattices were 

investigated by using first-principles molecular dynamics methods. It was found that the 

predominant mechanism involved the motion of negatively charged Si III -V III pairs via 

second-nearest neighbor jumps. This mechanism explained the ability of Si to disorder 

superlattices (regardless of whether it was introduced during growth or was in-diffused 

later), and the suppression of interdiffusion by compensation doping. The calculated 

activation energies were in very good agreement with experimental data. 

B.Chen, Q.M.Zhang, J.Bernholc: Physical Review B, 1994, 49[4], 2985-8 

[446-115/116-119] 

GaAs/AlAs: Si Diffusion 

Hall and photo-Hall measurements were performed on GaAs/AlAs short-period 

superlattices which were selectively doped with Si. The dopant was introduced 

selectively into the GaAs or AlAs layers, or at the interface. A superlattice which was 

doped uniformly in both layers was investigated for comparison. It was found that the 

electrical properties were controlled by DX deep donors which lay in the gap of the 

superlattice. The Hall data were explained in terms of a model which took account of the 

existence of two DX deep donors and a shallow donor which were both related to the Si 

impurity. It was found that the Si donor state in AlAs lay 0.06eV below the Si donor state 

in GaAs. The ionization energies of the DX states in GaAs and AlAs were calculated in 

order to account for the experimental results. The interpretation of Hall data in selectively 

doped samples required the assumption of Si segregation during epitaxy. 

P.Sellitto, P.Jeanjean, J.Sicart, J.L.Robert, R.Planel: Journal of Applied Physics, 1993, 

74[12], 7166-72 

[446-109/110-033] 

GaAs/AlGaAs: Si Diffusion 

Laser-assisted disordering was studied by using scanning electron microscopy and 

secondary ion mass spectrometry. This permitted the extent of the layer-disordered region 

to be correlated with the presence of laser-incorporated Si and O. Transmission electron 

microscopic studies permitted the determination of the distribution of Al and Ga at the 

interface between the impurity-disordered alloy and the as-grown crystal. The data 

revealed the occurrence of more rapid Si diffusion in the GaAs layers as compared with 

the Al-rich layers. 

J.E.Epler, F.A.Ponce, F.J.Endicott, T.L.Paoli: Journal of Applied Physics, 1988, 64[7], 

3439-44 

[446-72/73-026] 

GaAs/Si: Si Diffusion 




347

Si GaAs Sn 


detected at undoped Si/GaAs interfaces, whereas Si diffused into the GaAs (from P-doped 

or As-doped Si) to depths as great as 550nm after only 10s of annealing at 1050C. The 

flux of Si into GaAs was related to the fluxes of Ga and As into Si. Both fluxes increased 

with increasing dopant concentration in the Si. An enhanced diffusivity of Si into GaAs 

was attributed to the diffusion of point defects which were created by the diffusion of As 

and Ga into the encapsulant. It was deduced that the Si diffusivities in GaAs at doped 

polycrystalline Si interfaces were enhanced by factors of about 10000. 


64[4], 1845-54 

[446-72/73-027] 

GaAs/Si: Si Diffusion 

A new mechanism was proposed for the incorporation of Si into GaAs/Si hetero-epitaxial 

layers which were grown by metalorganic chemical vapor deposition. The mechanism 

involved gas-phase transport of Si to hetero-epitaxial layers during growth. This mode of 

Si uptake could operate as well as the previously proposed mechanism. The latter 

involved incorporation by enhanced diffusion, from the hetero-interface, via defects in the 

GaAs layer. 

S.Nozaki, J.J.Murray, A.T.Wu, T.George, E.R.Weber, M.Umeno: Applied Physics 

Letters, 1989, 55[16], 1674-6 

[446-72/73-027] 

GaAs/Si: Si Pipe Diffusion 

Preferential diffusion channels for Si were found in GaAs which had been grown, by 

means of metal-organic vapor phase epitaxy, onto Si(100). The density of these diffusion 

channels was consistent with the measured dislocation density. Also, by combining 

scanning electron microscopy and X-ray fluorescence, it was shown that a large amount 

of Si emerged at the surface within small [011]-type overgrowth-oriented defects that 

were present at the surface. 

A.Freundlich, A.Leycuras, J.C.Grenet, C.Grattepain: Applied Physics Letters, 1988, 

53[26], 2635-7 

[446-64/65-168] 

Sn 

GaAs: Sn Diffusion 


broadening of d-doped planes of Sn. It was found that the Sn planes were severely 

broadened by all 3 processes. At higher temperatures or densities, segregation or 

348

Sn GaAs Sn 


even during growth. 



[446-91/92-001] 

Table 26 

Diffusivity of Implanted Sn in GaAs 


1 x 10 14 1000 1.7 x 10 -13 

1 x 10 14 1000 1.1 x 10 -13 

1 x 10 15 1000 7.0 x 10 -14 

1 x 10 13 1000 3.0 x 10 -14 

1 x 10 14 1000 2.3 x 10 -14 

1 x 10 15 900 6.0 x 10 -14 

1 x 10 13 900 2.1 x 10 -14 

1 x 10 14 900 5.2 x 10 -15 

1 x 10 15 900 3.8 x 10 -15 

1 x 10 14 870 1.3 x 10 -15 

1 x 10 15 900 8.2 x 10 -16 

1 x 10 14 870 8.4 x 10 -16 

1 x 10 13 700 7.8 x 10 -19 

1 x 10 15 700 4.8 x 10 -19 

5 x 10 15 750 9.8 x 10 -18 

1 x 10 15 800 5.2 x 10 -17 

5 x 10 15 850 1.5 x 10 -16 

1 x 10 15 850 2.1 x 10 -16 

5 x 10 15 950 9.8 x 10 -16 

5 x 10 14 950 5.1 x 10 -15 

1 x 10 14 1000 1.5 x 10 -14 

326 GaAs: Sn Diffusion 




damage; the higher the concentration of extended defects, the slower was the diffusivity 





349

Sn GaAs Sn 



of diffusion. No time-dependence was observed. It was concluded that the results (table 







[446-84/85-016] 


The migration of implanted Sn was studied, as a function of dose and background doping, 

by means of secondary ion mass spectrometry. It was concluded that the diffusivity of the 

Sn depended upon the electron concentration, and that the Sn diffused via negatively 

charged Ga vacancies. The diffusivity of Sn outside of the implanted region was 

independent of the dose, and was associated with an activation energy of 1.98eV. 

E.L.Allen, M.D.Deal, J.D.Plummer: Journal of Applied Physics, 1990, 67[7], 3311-4 

[446-78/79-014] 


A scanned electron beam was used to diffuse Sn into GaAs from doped emulsions, and 

Rutherford back-scattering was used to investigate the results. It was found that diffusion 

was greatly enhanced by capping the emulsion with evaporated SiO 2 . 

Z.Meglicki, D.D.Cohen, A.G.Nassibian: Journal of Applied Physics, 1987, 62[5], 1778- 

81 

[446-55/56-006] 


The effects of strain upon the diffusion of Sn were studied by using laser Raman and 

photoluminescence spectroscopy. It was found that an increase in compressive strain 

produced an increase in the carrier concentration, while a decrease in compressive strain 

or an increase in tensile strain led to a decrease in the carrier concentration at the surface. 

This behavior was attributed to a decrease in the diffusion coefficient of Sn with 

compressive strain, and an increase with tensile strain. The data showed that the peak 

which was due to Ga antisite defects increased with increasing compressive stress. This 

indicated a decrease, in the Ga vacancy concentration, from the equilibrium concentration 

in an unstressed sample. Photoluminescence data for tensile-stressed samples revealed an 

increase in Ga vacancy concentration with respect to the equilibrium concentration in an 

unstressed sample. It was concluded that the change in diffusion coefficient with strain 

was related to Ga vacancies. It was found that the diffusion coefficient decreased 

350

Sn GaAs Ti 

exponentially with the compressive strain, and increased exponentially with tensile strain. 

The activation energy for Sn diffusion therefore varied linearly as a function of strain. 

A.B.M.Harun-ur Rashid, T.Katoda: Journal of Applied Physics, 1997, 81[4], 1661-9 

[446-148/149-173] 


The stresses which were generated at the GaAs/SiO 2 interface during annealing were 

investigated by means of laser Raman spectroscopy. It was found that compressive 

stresses existed at the surface of the GaAs after annealing, and that these increased with 

increasing SiO 2 cap layer thickness. The compressive stress on the GaAs was generated 

during annealing and was attributed to the difference in the thermal expansion 

coefficients of GaAs and SiO 2 . The increase in compressive stress on the surface of GaAs 

decreased the diffusion coefficient of Sn in the GaAs. This was due to a reduction in the 

number of Ga vacancies in the compressively stressed sample, as compared with the 

equilibrium Ga vacancy content of an unstressed sample. 

A.B.M.Harun-ur Rashid, M.Kishi, T.Katoda: Journal of Applied Physics, 1996, 80[6], 

3540-5 

[446-138/139-078] 


Films of SiO 2 spun-on glass, that were doped with Sn and/or Ga, were used as diffusion 

sources. The diffusion was studied during rapid thermal annealing, with or without any 

As over-pressure. It was found that the diffusivity of Sn decreased as the As over-pressure 

was increased. Modifying the Sn-doped spun-on glass, so as to contain 4mol%Ga, 

slightly reduced the Sn diffusivity. These results were explained in terms of chemical 

reactions between the spun-on glass and the GaAs. Highly doped layers (10 18 to 3 x 

10 18 /cm 3 ) were obtained. 

C.S.Hernandes, J.W.Swart, M.A.A.Pudenzi, G.T.Kraus, Y.Shacham-Diamand, 

E.P.Giannelis: Journal of the Electrochemical Society, 1995, 142[8], 2829-32 

[446-134/135-125] 

Ti 

GaAs: Ti Diffusion 

High-resolution X-ray photo-emission spectroscopy and Ar ion bombardment were used 

to study temperature-dependent chemical reactions and species redistribution in the 

Ti/GaAs(100) system. The results showed that Ti, deposited at room temperature, 

disrupted the GaAs substrate (by reacting with As) and released Ga into the over-layer. 

The As was found to accumulate, near to the buried interface, in the form of a Ti-As 

compound. The Ga was depleted, but accumulated beyond the reaction region. Sputter 

depth profiles indicated that high-temperature annealing caused Ti diffusion into the 

GaAs substrate and enhanced reaction with As. The rejection of Ga from the forming Ti- 

As compound became more severe when the amount of Ti-As increased. Heating 

351

Ti GaAs Ti 

promoted the segregation of rejected Ga atoms to the vacuum surface, but had little effect 

upon As segregation. 

F.Xu, D.M.Hill, Z.Lin, S.G.Anderson, Y.Shapira, J.H.Weaver: Physical Review B, 1988, 

37[17], 10295-300 

[446-62/63-211] 

1.0E-10 

1.0E-11 

1.0E-12 

1.0E-13 

1.0E-14 

D (cm 2 /s) 

1.0E-15 

1.0E-16 

1.0E-17 

1.0E-18 

1.0E-19 

table 27 

table 28 

table 29 

table 30 

table 31 

1.0E-20 

7 8 9 10 11 12 13 

10 4 /T(K) 

Figure 7: Diffusivity of Zn in GaAs 

GaAs: Ti Diffusion 

The structural properties of samples which had been implanted with 150 or 400keV Ti, to 




results indicated that the Ti did not redistribute at all. 


1992, 72[8], 3514-21 

[446-106/107-034] 

352

Zn GaAs Zn 

Zn 

GaAs: Zn Diffusion 


using metalorganic vapor-phase epitaxy, was studied using capacitance-voltage etch 

profiling and secondary ion mass spectrometry. It was deduced that the diffusivity of Zn 

in GaAs could be described by: 

D (cm 2 /s) = 4.6 x 10 -4 exp[-2.1(eV)/kT] 

for rapid thermal annealing, while the diffusivity could be described by: 

D (cm 2 /s) = 1.2 x 10 -6 exp[-1.8(eV)/kT] 

for furnace annealing. A model which was based upon an interstitial cum substitutional 

diffusion mechanism, with certain kinetic limitations, was successfully used to simulate 

the observed dopant concentration profiles. Markedly anomalous diffusion of Zn, from 

GaAs and into highly n-doped GaAlAs, was found. 

N.Nordell, P.Ojala, W.H.Van Berlo, G.Landgren, M.K.Linnarsson: Journal of Applied 

Physics, 1990, 67[2], 778-86 

[446-74-004] 


It was noted that the formation of dislocations and precipitates in monocrystalline 

substrates occurred when elemental gas sources with a sufficiently high Zn partial 

pressure were used to diffuse Zn into the surface regions of the substrate. A study of the 

nature of these defects at various depths in the Zn concentration profiles permitted 

conclusions to be drawn concerning defect formation and evolution during diffusion, and 

revealed valuable information concerning the point defects which were involved in Zn 

diffusion. 

A.Rucki, W.Jäger: Defect and Diffusion Forum, 1997, 143-147, 1095-100 

[446-143/147-1095] 

327 GaAs: Zn Diffusion 

It was recalled that Zn in GaAs exhibited a complicated diffusion behavior, as well as a 

high solid-solubility. The diffusion of Zn, at temperatures of between 700 and 1100C, 

was studied here by using 3 different Zn diffusion sources. In order to compare the 

penetration curves for the various sources, reduced penetration depths and reduced 

concentrations were calculated. Numerical simulation of Zn transport furnished a good 

description for particular sources. Effective diffusivities (table 27) were deduced from the 

experimental data and were normalized to standard vapor pressures and electronically 

intrinsic conditions. The data for normalized Zn diffusion could be described by: 

D (cm 2 /s) = 82.3 exp[-4.03(eV)/kT] 

H.G.Hettwer, N.A.Stolwijk, H.Mehrer: Defect and Diffusion Forum, 1997, 143-147, 

1117-24 

[446-143/147-1117] 

353


Table 27 

Diffusivity of Zn in GaAs as a Function of Diffusion-Source Composition 

Temperature (C) As Zn Ga D (cm 2 /s) 

700 0.054 0.698 0.248 1.08 x 10 -19 

900 0.547 0.414 0.039 2.52 x 10 -16 

900 0.535 0.429 0.036 1.88 x 10 -16 

900 0.246 0.508 0.246 1.20 x 10 -15 

900 0.075 0.130 0.795 1.66 x 10 -15 

1050 0.542 0.323 0.135 2.24 x 10 -14 

1050 0.474 0.366 0.160 2.94 x 10 -14 

1050 0.330 0.333 0.337 5.26 x 10 -14 


It was pointed out that Zn was one of the main p-type dopants which were used for the 

fabrication of devices that were based upon GaAs or related III-V materials. The element 

dissolved substitutionally on the group-III sub-lattice, and diffused via a kick-out 

mechanism which involved group-III self-interstitials. Non-equilibrium concentrations of 

these self-interstitials had a marked effect upon the diffusivity of Zn. Various situations 

were considered in which non-equilibrium point defects played a role in Zn diffusion. 

These included the in-diffusion of such dopants from an external source, the diffusion of 

grown-in dopants, and self-interstitial generation by Fermi-level surface pinning. It was 

noted that the diffusion behavior of C, which was found on the group-V sub-lattice of 

GaAs, was much less sensitive to non-equilibrium point defects. It was therefore used to 

replace Zn as a p-type dopant. 

M.Uematsu, K.Wada, U.Gösele: Applied Physics A, 1992, 55[4], 301-12 

[446-93/94-008] 


The chemical processes which occurred during the diffusion of Zn from spun-on SiO 2 

films and into wafer samples were investigated. It was found that, at 600C, only about 

45% of the film was changed to SiO 2 glass. At 700C, Zn silicates began to form. When 

the Zn was in this form, it was much less able to diffuse into the GaAs. 

E.Nowak, G.Kühn, B.Schumann, R.Hesse, H.Sobotta: Crystal Research and Technology, 

1992, 27[4], 503-8 

[446-91/92-008] 

354



The rapid thermal diffusion of Zn into semi-insulating material, from spun-on silica films, 

was investigated. Depending upon the heating rate, 2 types of secondary ion mass 

spectrometry profile were observed. 

E.Nowak, G.Kühn, T.Morgenstern, B.Schumann: Crystal Research and Technology, 

1991, 26[8], 981-6 

[446-88/89-014] 


The migration of Zn into material which had been grown by means of molecular beam 

epitaxy at 200C was studied. The diffusion coefficient was found, using the sealed 

ampoule technique, to be an order of magnitude higher in so-called low-temperature 

material (1.3 x 10 -11 cm 2 /s) than it was in normal material (1.5 x 10 -12 cm 2 /s). This 

difference was attributed to the large number of defects (including As antisites) which 

were present in the low-temperature material. It was noted that the effectiveness of a 

Si 3 N 4 diffusion mask depended upon whether the mask was deposited directly onto the 

low-temperature material. A failure of such masks to stop Zn diffusion was attributed to 

the effect of As atoms which out-diffused from the As-rich low-temperature material and 

into the Si 3 N 4 . The presence of a 10nm GaAs layer on the low-temperature material was 

effective in conserving the masking properties of the nitride. 

Y.K.Sin, Y.Hwang, T.Zhang, R.M.Kolbas: Journal of Electronic Materials, 1991, 20[6], 

465-70 

[446-88/89-015] 


The diffusion of Zn was studied by using liquid-phase epitaxy methods, and Si-doped n- 

type substrate material. The measurements were carried out at 850C, and dopant 

concentrations which ranged from 10 18 to 10 19 /cm 3 were introduced. It was found that the 

Zn concentration in the solid depended upon the square root of the atomic fraction of Zn 

in the liquid. The diffusivity was dominated by the interstitial-substitutional process, and 

exhibited a cubic dependence upon the Zn content. The Zn interstitial was mainly doublyionized 

Zn i 2+ . 

C.Algora, G.L.Araujo, A.Marti: Journal of Applied Physics, 1990, 68[6], 2723-30 

[446-86/87-002] 


The migration of Zn in AlGaAs at 650C was studied by using the sealed-ampoule method 

and a ZnAs 2 source. It was found that the results for zero Al content (table 28) could be 

described by the expression: 

D (cm 2 /s) = 26 exp[-2.47(eV)/kT] 

V.Quintana, J.J.Clemencon, A.K.Chin: Journal of Applied Physics, 1988, 63[7], 2454-5 

[446-72/73-003] 

355


Table 28 

Bulk Diffusivity of Zn in GaAs 


700 4.08 x 10 -12 

650 1.00 x 10 -12 

600 1.84 x 10 -13 

550 1.83 x 10 -14 


In order to study its diffusion mechanism, Zn was diffused from a ZnAs 2 source into Sidoped 

samples, at temperatures ranging from 575 to 700C, in sealed evacuated quartz 

tubes. The samples were characterized by means of the depth profile of the 

photoluminescence at various temperatures. The photoluminescence spectra contained 

characteristic emissions which were associated with deep levels of Ga and As vacancies. 

A detailed analysis of the spectra revealed the role that was played by vacancies in the Zn 

diffusion process. A spatial correlation between the luminescence spectra, and the Zn 

concentrations deduced from secondary ion mass spectrometric data, was demonstrated. It 

was found that the data (table 29) could be described by the expression: 

D(cm 2 /s) = 2.05 exp[-2.28(eV)/kT] 

N.H.Ky, L.Pavesi, D.Araujo, J.D.Ganière, F.K.Reinhart: Journal of Applied Physics, 

1991, 69[11], 7585-93 

[446-86/87-012] 

Table 29 

Diffusivity of Zn in GaAs 


700 3.33 x 10 -12 

650 7.12 x 10 -13 

600 1.72 x 10 -13 

575 5.43 x 10 -14 


A marked difference between the Be and C distributions in this material was noted after 

Zn diffusion. It was shown that grown-in C was stable, and remained localized even after 

Zn diffusion. On the other hand, Be diffused very rapidly in the presence of diffusing Zn. 

356


This effect was attributed to differences in the crystal lattice sites which the dopant 

occupied. 

E.Tokumitsu, T.H.Chiu, H.S.Luftman, N.T.Ha: Journal of Applied Physics, 1991, 69[12], 

8426-8 

[446-86/87-013] 


The Zn diffusion doping of GaAs, by using metalorganic vapor-phase epitaxy and 

diethylzinc as a dopant source, was examined. Typical Zn concentrations and depths 

which were obtained were 10 19 to 10 21 /cm 3 , and 40 to 200nm. The highest concentration 

gradient which was obtained in this way was 4 orders of magnitude per 50nm, and the 

highest Zn concentration was 2 x 10 21 /cm 3 at the sample surface. 

Z.F.Paska, D.Haga, B.Willén, M.K.Linnarsson: Applied Physics Letters, 1992, 60[13], 

1594-6 

[446-86/87-013] 



and into GaAs was described. The oxide films were deposited by using metal-organic 

chemical vapor deposition. A capping layer of SiO 2 was deposited on top of the source 

films, and diffusion was carried out in flowing N at 650C. Diffusion depths of between 

200nm and several microns could be easily obtained. The diffusion front in n-type 

substrates was sharp. The dependence of the diffusion depth upon the source film 



[446-78/79-002] 


Rapid thermal annealing was used to diffuse Zn into GaAs from a thin-film zinc silicate 

source that was prepared by chemical vapor deposition at atmospheric pressure. 

Comparisons were made with conventional open-tube furnace annealing, for a diffusion 

temperature of 650C. The diffusivities were found to be similar; in contrast to previous 

results. At temperatures ranging from 650 to 750C, sharp Zn diffusion profiles were 

observed. At temperatures above 750C, kinks in the diffusion profiles were found. Such 

kinks were also observed when semi-insulating substrates were used instead of Si-doped 

n + -type substrates. A model for Zn diffusion had already been developed, and this was 

based upon the pairing of interstitial Zn with all of the acceptor species which were 

present during diffusion. The predominant species were found to be substitutional Zn, and 

Ga vacancies. The concentration of the latter was a function of the background dopant 

concentration. The results of the model were shown to agree with all of the present 

experimental evidence, and were also consistent with the experimental observation of 2 

distinct activation energies for Zn diffusion into n + -doped substrates. These energies were 

357


equal to 1.1 and 2.6eV for temperatures which were above or below about 790C, 

respectively. 

G.Rajeswaran, K.B.Kahen, D.J.Lawrence: Journal of Applied Physics, 1991, 69[3], 

1359-65 

[446-78/79-014] 


Samples were diffused with Zn, via a 200 to 300nm protective ZrO 2 layer. The diffusion 

depth exhibited a square-root time dependence. The absolute diffusivity values depended 

slightly upon the diffusion conditions (table 30). The layer had essentially no effect upon 

the carrier concentration profile or the activation energy. 


[446-74-003] 

Table 30 


Protection Source Temperature (C) D (cm 2 /s) 

ZrO 2 Zn 755 3.8 x 10 -11 

ZrO 2 Zn 700 1.2 x 10 -11 

ZrO 2 Zn 650 6.9 x 10 -12 

ZrO 2 Zn 600 1.3 x 10 -12 

ZrO 2 GaAs/Zn 2 As 3 650 2.0 x 10 -12 

- GaAs/Zn 2 As 3 650 1.6 x 10 -12 








The saturation behavior of the free carrier concentrations in p-type GaAs monocrystals 

which had been doped by Zn diffusion was investigated. Free-hole saturation occurred at 

10 20 /cm 3 . The difference in saturation hole concentrations of materials was investigated 

by studying the incorporation and lattice location of Zn. The latter was an acceptor when 

located on a group-III atom site. Zinc was diffused into III-V wafers in a sealed quartz 

ampoule. Particle-induced X-ray emission and ion-channelling techniques were then used 

to determine the exact lattice location of Zn atoms. It was found that more than 90% of 

the Zn atoms occupied Ga sites in diffused GaAs samples. The results were analyzed in 

terms of the amphoteric native defect model. It was shown that differences in the 

358


electrical activities of Zn atoms in various materials were a consequence of the differing 

locations of the Fermi-level stabilization energy. 

L.Y.Chan, K.M.Yu, M.Ben-Tzur, E.E.Haller, J.M.Jaklevic, W.Walukiewicz, 

C.M.Hanson: Journal of Applied Physics, 1991, 69[5], 2998-3006 

[446-78/79-015] 


The use of thin Si films for the selective-area diffusion of Si and Zn was described. It was 

found that Si films behaved as ideal masks for Zn diffusion at temperatures below 750C. 

Ideal lateral Zn diffusion profiles were also observed when using these films; regardless 

of the stress at the interface. 

G.A.Vawter, E.Omura, X.S.Wu, J.L.Merz, L.Coldren, E.Hu: Journal of Applied Physics, 

1988, 63[11], 5541-7 

[446-72/73-003] 


The growth and diffusion of abrupt Zn profiles in un-doped, or Si-doped, material was 

monitored by means of secondary ion mass spectrometry. The sharp diffusion fronts 

which resulted from annealing treatments indicated that the Zn diffusion coefficient was 

concentration-dependent. The data were encompassed by the curves: 

D(cm 2 /s) = 5.5 x 10 -10 exp[-0.8(eV)/kT] 

and 

D(cm 2 /s) = 1.8 x 10 -9 exp[-1.0(eV)/kT] 

However, the diffusion of Zn at high concentrations appeared to be inhibited by crystal 

defect kinetics and resulted in a relatively concentration-independent Zn diffusion 

coefficient. The V/III growth ratio did not have any effect upon Zn diffusion in un-doped 

or Si-doped material. The diffusion of Zn in heterojunction bipolar transistor structures 

was different; in that the diffusion of Zn into a GaAs collector was larger by an order of 

magnitude, and decreased with an increase in the V/III growth ratio. In addition, the 

diffusion of Zn into an AlGaAs emitter was markedly lower and was inhibited by an 

increase in the V/III ratio. These data could be summarized by the expressions: 

GaAs (V/III = 60): D(cm 2 /s) = 2.0 x 10 3 exp[-3.0(eV)/kT] 

GaAs (V/III = 120): D(cm 2 /s) = 1.0 x 10 6 exp[-3.6(eV)/kT] 

Ga 0.78 Al 0.22 As (V/III = 60): D(cm 2 /s) = 6.8 x 10 4 exp[-3.4(eV)/kT] 

Ga 0.78 Al 0.22 As (V/III = 120): D(cm 2 /s) = 1.1 x 10 -5 exp[-1.7(eV)/kT] 

P.Enquist, J.A.Hutchby, T.J.De Lyon. Journal of Applied Physics, 1988, 63[9], 4485-93 

[446-72/73-012] 


The anomalous shape of Zn diffusion profiles in GaAs was quantitatively explained. The 

Frank-Turnbull mechanism was suggested to govern interchanges, between interstitial and 

substitutional Zn, via Ga vacancies. It was proposed that these vacancies were either 

neutral or were singly ionized; depending upon the position of the Fermi level. In 

359


addition, 2 physical phenomena were proposed. Substitutional Zn thermally generated 

interstitial Zn-Ga vacancy pairs, and there was pairing between the donor (interstitial Zn) 

and the acceptor (substitutional Zn). The model furnished good agreement with 

experimental data. 

K.B.Kahen: Applied Physics Letters, 1989, 55[20], 2117-9 

[446-72/73-012] 



disordering. With regard to the influence of Zn p-type dopants, the Fermi level effect had 

to be considered. In accord with its effect upon superlattice disordering, Zn diffusion 

appeared to be governed by the kick-out mechanism. It was concluded that dislocations in 

this material and in other III-V compounds were only moderately efficient sinks or 

sources for point defects. 


[446-62/63-208] 


The diffusion of Zn into GaAs, from diethylzinc and trimethylarsenic, was studied. The 

process produced surface hole concentrations which were greater than 10 20 /cm 3 . Wellcontrolled 

junction depths which could be as shallow as 0.000lmm were obtained, and a 

smooth surface morphology was retained. The profile shape was much more complex 

than those predicted by accepted interstitial cum substitutional Zn diffusion models. In 

order to explain the observed profiles, a new model for Zn diffusion was proposed, and 

used in a computer simulation. 

S.Reynolds, D.W.Vook, J.F.Gibbons: Journal of Applied Physics, 1988, 63[4], 1052-9 

[446-62/63-211] 


Data were presented which showed that low-temperature (680C) Zn diffusion was 

effective in reducing the dislocation density of epitaxial GaAs which was grown onto Si. 

The GaAs/Si system was analyzed by using both cross-sectional and plan-view 

transmission electron microscopy. Simple thermal annealing of GaAs/Si at a higher 

temperature (850C) was also studied. The reduction in dislocation density which occurred 

due to Zn diffusion was suggested to arise from the increased concentrations of point 

defects which were generated during Zn diffusion. This resulted in enhanced dislocation 

climb. The mechanism was consistent with impurity-induced layer disordering, via Zn 

diffusion, in AlGaAs/GaAs heterostructures. 

D.G.Deppe, N.Holonyak, K.C.Hsieh, D.W.Nam, W.E.Plano, R.J.Matyi, H.Shichijo: 

Applied Physics Letters, 1988, 52[21], 1812-4 

[446-62/63-211] 

360



Second diffusion of Zn was observed using low concentrations. The behavior was similar 

to that of double diffusion in InP. The effect of the Zn activity in the vapor phase was 

studied by using a semi-closed box system. The observed Zn profiles were explained in 

terms of a model which involved varying charge transfer by vacancy centers during 

interstitial-substitutional interchanges. 

K.Kazmierski, F.Launay, B.De Cremoux: Japanese Journal of Applied Physics, 1987, 

26[10], 1630-3 

[446-55/56-007] 


A ternary Zn 3 As 2 -ZnAs 2 -GaAs source for the diffusion of Zn into GaAs was developed 

by using a low-temperature sintering technique. It was noted that wafers which were 

diffused by using this source remained free from damage. The dopant concentrations and 

diffusion depths agreed with the results which were obtained by using high-temperature 

Ga/As/Zn diffusion sources. 

J.Werner, H.Melchior: Japanese Journal of Applied Physics, 1987, 26[4], 641-2 

[446-51/52-117] 






profile. 



[446-121/122-045] 


The isothermal diffusion of Zn was described in terms of the Longini reaction and a 

recombination process. It was assumed that, during diffusion, highly mobile Zn 

interstitials recombined with a Ga vacancy and became a relatively immobile site defect. 

It was noted that the long-term profile of the total Zn concentration was governed mainly 

by the vacancy concentration profile. All of the known concentration profiles could be 

obtained for a constant diffusivity, without using fitting parameters. A new method was 

proposed for determining the diffusivities of interstitial Zn, of Zn in Ga lattice sites, and 

of vacancies. These coefficients were deduced from existing experimental data. It was 

shown that the apparent dependence of the Zn diffusivity upon its background 

concentration was due to its recombination with Ga vacancies. 

N.N.Grigorev, T.A.Kudykina: Fizika i Tekhnika Poluprovodnikov, 1997, 31[6], 697-702 

(Semiconductors, 1997, 31[6], 595-9) 

[446-157/159-361] 

361



The diffusion of implanted Zn was studied, at annealing temperatures of between 625 and 

850C, by means of secondary ion mass spectrometry. A substitutional-interstitial 

diffusion mechanism was proposed in order to explain how deviations of the local Ga 

interstitial concentration, from its equilibrium value, regulated Zn diffusion. It was found 

that it was possible to simulate both box-shaped profiles, that resulted from hightemperature 

annealing, and kink-and-tail profiles which resulted from lower-temperature 

annealing. The simulation data permitted the determination of Arrhenius relationships for 

the intrinsic diffusion coefficient of implanted Zn. This (table 31) could be described by: 

D (cm 2 /s) = 0.6075 exp[-3.21(eV)/kT] 

M.P.Chase, M.D.Deal, J.D.Plummer: Journal of Applied Physics, 1997, 81[4], 1670-6 

[446-148/149-172] 

Table 31 



850 1.9 x 10 -15 

800 1.2 x 10 -15 

800 6.0 x 10 -16 

750 1.2 x 10 -16 

750 8.4 x 10 -17 

675 8.4 x 10 -18 

675 3.1 x 10 -18 


The diffusion behavior was investigated by using a spun-on doped silica film. Diffusion 

annealing was carried out, at temperatures which ranged from 650 to 800C, using a SiO 2 

or Si 3 N 4 cap. At higher temperatures, the sample surface after diffusion was found to be 

damaged. The profiles which were obtained at temperatures ranging from 700 to 780C 

were of abrupt or kink-and-tail type. Diffusion annealing which was carried out using a 

nitride cap were termed Ga-rich, and produced abrupt box-like profiles. The use of an 

oxide cap was expected to produce a superposition of shallow As-rich and deeper Ga-rich 

profiles; leading to a characteristic kink-and-tail profile. This behavior was observed at 

higher temperatures, but the junctions were too shallow to exhibit a tail at lower 

temperatures. The junction depths were attributed to the existence of differing activation 

energies for the two types of cap material. The energy that was required for oxide-cap 

diffusion was close to published values for oxide caps. The lower values which were 

found for nitride-cap diffusion were closer to published values for a phosphosilicate glass 

362


cap. These results clearly showed that the nature of the cap material determined the type 

of diffusion profile. 

S.Chatterjee, K.N.Bhat, P.R.S.Rao: Solid-State Electronics, 1997, 41[3], 496-500 

[446-148/149-174] 


The in-diffusion of Zn at high concentrations, under As-deficient conditions, caused the 

generation of dislocation loops, elongated dislocations and Ga precipitates decorated with 

voids within the diffusion zone. Similar treatment under As vapor led to the recovery of 

diffusion-induced damage in the sub-surface region. This was accompanied by the 

appearance of 2 distinct steps in the Zn concentration profile. Previous work had 

suggested that these phenomena were connected with the out-diffusion of Ga from 

precipitates and towards the surface. The present results showed that the replacement, by 

P or Sb, of As in the diffusion ambient produced similar recovery effects. 

G.Boesker, H.G.Hettwer, A.Rucki, N.A.Stolwijk, H.Mehrer, W.Jaeger, K.Urban: 

Materials Chemistry and Physics, 1995, 42[1], 68-71 

[446-136/137-110] 


The diffusion of Zn, at high concentrations, into samples under As-deficient ambient 

conditions led to the generation of defects such as dislocation loops, elongated 

dislocations and Ga precipitates - decorated with voids- throughout the diffusion zone. 

Similar treatments under As vapor led to the recovery of diffusion-induced damage in 

regions beneath the surface. This was accompanied by the appearance of 2 distinct steps 

in the Zn concentration profile. Previous work had suggested that these phenomena were 

associated with the out-diffusion of Ga from the precipitates and towards the surface. It 

was shown here that the replacement of As by P or Sb produced similar recovery effects. 

This indicated that ambient group-V elements made the near-surface of GaAs into an 

effective sink for diffusion-induced excess Ga. 

G.Bösker, H.G.Hettwer, A.Rucki, N.A.Stolwijk, H.Mehrer, W.Jäger, K.Urban: Materials 

Chemistry and Physics, 1995, 42[1], 68-71 

[446-134/135-126] 


The diffusivity of Zn in heavily-doped pnpn structures was measured after growth and 

annealing. During growth at 650C, the Zn diffusivity of about 10 -12 cm 2 /s in the buried p- 

type layer was found to be more than 10000 times the Zn diffusivity in the top p-type 

layer. During annealing at 800C, the Zn diffusivity of about 5 x 10 -14 cm 2 /s in the buried 

layer remained orders of magnitude greater than the Zn diffusivity in the top layer. The 

measurements provided clear experimental evidence that a large flux of Ga interstitials 

was injected from the surface during the growth of n-type layers, and that the Ga 

interstitials were trapped in the buried p-type layer by the electric field of the pn junctions 

(and were therefore positively charged). It was suggested that the resultant large 

concentration of Ga interstitials in the buried layer accounted for the increased Zn 

363


diffusivity via a kick-out mechanism. Finally, it was deduced that the mobile Zn 

interstitial was positively charged. 

C.Y.Chen, R.M.Cohen, D.S.Simons, P.H.Chi: Applied Physics Letters, 1995, 67[10], 

1402-4 

[446-125/126-121] 


It was recalled that the rapid diffusion of Zn into GaAs had recently been attributed to the 

fact that a small fraction of the Zn interstitials changed to Ga sites, thereby producing 

interstitial Ga (I Ga ). This kick-out reaction led to the possibility of determining the I Ga 

transport properties from Zn diffusion experiments on essentially perfect GaAs. However, 

previous attempts had been impeded by the diffusion-induced generation of 

microstructural defects which acted as I Ga sinks. In the present study, this was prevented 

by using Zn-doped GaAs powder as a diffusion source. The measured 2-stage profiles 

showed that, under these conditions, Zn diffusion at 906C was controlled by I Ga 3+ in 

addition to I Ga 2+ . Analysis of the profiles yielded quantitative data on the Ga- and Znrelated 

diffusivities, and on the concentration of I Ga . 

G.Bösker, N.A.Stolwijk, H.G.Hettwer, A.Rucki, W.Jäger, U.Södervall: Physical Review 

B, 1995, 52[16], 11927-31 

[446-125/126-121] 


The Zn was diffused into GaAs through anodic oxide layers with various thicknesses and 

densities. Electrochemical profiling was used to determine the electrically active Zn 

concentration and the diffusion depth. It was found that the depth of the junction varied 

inversely with the thickness and density of the oxide. However, the surface concentration 

appeared to be independent of the oxide thickness or density and attained a value which 

was identical to that which was found for diffusion into a bare GaAs sample. The results 

demonstrated that the most important effect of the oxide was to delay the introduction of 

Zn into the lattice. Thus, the anodic oxide could not be used as a mask or a Zn 

concentration attenuator. 

H.Cutlerywala, R.J.Roedel: Journal of the Electrochemical Society, 1994, 141[6], 1639- 

43 

[446-119/120-192] 


The Zn was introduced by using the electron beam doping method. That is, a Zn sheet 

was sandwiched between GaAs wafers and the surface of the GaAs was irradiated with 

7MeV electrons. The use of secondary ion mass spectroscopy revealed U-shaped 

diffusion profiles for impurities in the substrates. The results could be explained in terms 

of the kick-out mechanism, and surface diffusion processes. 

A.Takeda, T.Wada: Materials Science Forum, 1994, 143-147, 1421-6 

[446-113/114-013] 

364



Secondary ion mass spectroscopy and photoluminescence studies were made of V Ga 

during Zn diffusion into Si-doped material. Photoluminescence spectra were obtained at 

various etching depths below the sample surface. After annealing the samples in excess 

As 4 vapor at 650C, the conversion of n-type material into p-type material was observed 

by making electrical measurements in the region near to the sample surface. The 

importance of the Si Ga -V Ga emission band in photoluminescence spectra from thermally 

converted regions indicated that V Ga which were generated at the surface during 

annealing were responsible for the thermal conversion. The results also showed that the 

main point defects which were generated during annealing under Ga-rich conditions were 

V As and Ga As . In the case of Zn-diffused Si-doped substrates at 600C, the disappearance 

of the V Ga -related band from the photoluminescence spectra of diffused regions furnished 

evidence for the incorporation of Zn interstitials into Ga sites during diffusion. An 

accumulation of V Ga was found ahead of the Zn diffusion front. The Zn-diffused samples 

were also annealed at 800C for 2h in vacuum, or in As 4 vapor with or without a Si 3 N 4 

cap. In the case of samples which were annealed in vacuum, an abrupt diffusion front 

advanced slightly into the bulk; with a supersaturation of V Ga ahead of the front. On the 

other hand, samples which were annealed in As vapor, with or without a cap on the 

surface, exhibited double Zn concentration profiles with an undersaturation of V Ga around 

the tail region. These results revealed the important role which was played by nonequilibrium 

point defects, and were explained in terms of a kick-out mechanism for Zn 

diffusion. 

N.H.Ky, J.D.Ganière, F.K.Reinhart, B.Blanchard: Materials Science Forum, 1994, 143- 

147, 1397-402 

[446-113/114-014] 


Experimental studies were made of Zn diffusion, using high temperatures, open tubes, 

and SiO 2 protective layers. Precisely controlled diffusion depths of less than 0.2 or 0.4µ 

could be obtained by using SiO 2 , doped with 0.1%Zn, as a diffusion source. Under these 

conditions, the diffusion coefficients were equal to 1.31 x 10 -13 and 1.76 x 10 -12 cm 2 /s at 

temperatures of 700 and 1025C, respectively. 

D.K.Gautam, Y.Shimogaki, Y.Nakano, K.Tada: Materials Science Forum, 1993, 117- 

118, 417-22 

[446-111/112-051] 


After Zn diffusion into Si-doped material, the diffused samples were annealed in vacuum, 

in As vapor, or with a Si 3 N 4 mask capping the surface. The Zn concentration profiles 

which were obtained by secondary ion mass spectroscopy, and photoluminescence spectra 

for various depths below the sample surface, were studied in detail. After annealing in 

vacuum, the steep (p + -n) Zn diffusion front advanced into the bulk. It was observed that 

365


the intensity ratio between the Si donor-Ga vacancy complex (Si Ga -V Ga ) related emission 

band and the band-to-band (e-h) transition was enhanced in the region ahead of the Zn 

diffusion front. On the other hand, Zn atoms diffused deeper into the bulk of samples 

which were annealed in As vapor, with or without a capping layer. These samples 

exhibited kink-and-tail (p + -p-n ) Zn concentration profiles with a decrease in the intensity 

ratio around the tail region. Analysis of the photoluminescence data suggested that there 

was a supersaturation of Ga vacancies ahead of the diffusion front of samples which were 

annealed in vacuum, and an under-saturation of this defect around the tail region of 

samples which were annealed in As vapor. The results emphasized the important role 

which was played by non-equilibrium of the defect concentration during post-diffusion 

annealing. This permitted an anomalous Zn double profile to be explained in terms of the 

interstitial-substitutional mechanism. 

N.H.Ky, J.D.Ganière, F.K.Reinhart, B.Blanchard, J.C.Pfister: Journal of Applied Physics, 

1993, 74[9], 5493-500 

[446-111/112-051] 


The microscopic mechanisms of Zn diffusion in GaAs were investigated by using ab 

initio molecular dynamics techniques. It was found that, among the various proposed 

mechanisms for Zn diffusion, kick-out by Ga interstitials had the lowest activation 

energy. The occurrence of Zn in-diffusion generated non-equilibrium group-III 

interstitials which were bound to Zn by Coulomb forces. The interstitials followed the Zn 

diffusion front and disordered the superlattice. The calculated activation energies for 

these processes were in good agreement with experimental data. 

C.Wang, Q.M.Zhang, J.Bernholc: Physical Review Letters, 1992, 69[26], 3789-92 

[446-106/107-039] 


The electrical activity and lattice-site locations of Zn atoms which had been diffused into 

GaAs were studied by using various characterization techniques. Particle-induced X-ray 

emission channelling showed that all of the Zn atoms were substitutional and were 

electrically active acceptors. A difference between the behaviors of Zn in GaAs and InP 

could be understood in terms of the amphoteric native defect model. It was also shown 

that the Fermi level stabilization energy provided a convenient energy reference for the 

treatment of dopant diffusion at semiconductor hetero-interfaces. 

W.Walukiewicz, K.M.Yu, L.Y.Chan, J.Jaklevic, E.E.Haller: Materials Science Forum, 

1992, 83-87, 941-6 

[446-99/100-065] 


A model was developed which could explain the nature of Zn diffusion profiles in n + -type 

material. The model was based upon the effect of Coulomb pairing between interstitial 

and substitutional Zn. By extending the model so as to include the Coulomb pairing of 

interstitial Zn with all of the acceptors which were present during diffusion, the 

366


predictions were caused to be in good agreement with experimental data. Only one 

adjustable parameter was involved. 

K.B.Kahen, J.P.Spence, G.Rajeswaran: Journal of Applied Physics, 1991, 70[4], 2464-6 

[446-91/92-008] 

GaAs/AlAs: Zn Diffusion 

The microscopic mechanisms of Zn-induced interdiffusion in GaAs/AlAs superlattices, 

were investigated by using ab initio molecular dynamics techniques. It was found that, 

among the various proposed mechanisms for Zn diffusion, kick-out by Ga interstitials had 

the lowest activation energy. The occurrence of Zn in-diffusion generated nonequilibrium 

group-III interstitials which were bound to Zn by Coulomb forces. The 

interstitials followed the Zn diffusion front and disordered the superlattice. The calculated 

activation energies for these processes were in good agreement with experimental data. 

C.Wang, Q.M.Zhang, J.Bernholc: Physical Review Letters, 1992, 69[26], 3789-92 

[446-106/107-039] 


A model was proposed for the effect of Zn in-diffusion in enhancing superlattice 

disordering. It combined recently proposed models for Ga self-diffusion and Zn diffusion 

in GaAs. Four coupled partial differential equations, which described the process, were 

solved numerically. Satisfactory agreement was obtained between the simulated results 

and published experimental data. At a given temperature, and for the values which were 

assumed for the diffusion coefficient and thermal equilibrium concentration of point 

defects, doubly positively charged Ga self-interstitials, I Ga 2+ , were deduced to be a 

consistent splitting of the known Ga self-diffusion coefficient which was dominated by 

I Ga 2+ . The superlattice disordering enhancement was due mainly to the Fermi-level effect, 

but I Ga 2+ supersaturation also made a small contribution. Because of p-doping by Zn 

acceptor atoms, the I Ga 2+ concentration was greatly increased via the Fermi-level effect. 

An I Ga 2+ supersaturation also developed because the I Ga 2+ generation rate was higher than 

its removal rate. Enhanced superlattice disordering occurred mainly under Ga-rich 

superlattice conditions. The Zn in-diffusion enhanced Al-Ga interdiffusion coefficient 

exhibited an apparent dependence, upon the Zn s - concentration, which differed slightly 

from a quadratic relationship. 

H.Zimmermann, U.Gösele, T.Y.Tan: Journal of Applied Physics, 1993, 73[1], 150-7 

[446-106/107-078] 


The Car-Parrinello method was used to study Zn-enhanced interdiffusion in superlattices. 

The energetics of several mechanisms for the diffusion of Zn were examined. It was 

found that a pair which consisted of a substitutional Zn acceptor and an interstitial group- 

III atom had a substantially lower formation energy than did an isolated interstitial. The 

low formation energy of this pair resulted in the interstitial kick-out mechanism having a 

much lower activation energy than those which involved vacancies or the dissociative 

367


(Frank-Turnbull or Longini) mechanism. The lowest-energy path for the interchange of 

group-III atoms involved the kick-out of Zn by a group-III interstitial, followed by fast Zn 

interstitial diffusion and the subsequent ejection of another group-III atom into the 

interstitial channel. The activation energies for these processes, as determined by 

following the kick-out trajectories and including full relaxation of all of the atoms, were 

in good agreement with experimental data. 

Q.M.Zhang, C.Wang, J.Bernholc: Materials Science Forum, 1992, 83-87, 1351-6 

[446-99/100-083] 

GaAs/AlGaAs: Zn Diffusion 

Multiple quantum well structures, with the same well thickness but differing Al x Ga 1-x As 

compositions (with x = 0.1, 0.2, 0.45, or 1), were grown by using molecular beam 

epitaxy. After Zn diffusion at 575C (1 to 16h), the structures were studied by means of 

the transmission electron microscopy of cleaved wedges, secondary electron imaging in a 

scanning electron microscope, and by means of secondary ion mass spectroscopy. The 

results showed that totally and partially disordered regions always lay beyond the Zn 

diffusion front. The extent of partial disordering depended upon the value of x. As x 

increased, the disordering rate increased due to an increase in Zn diffusivity. The effect of 

a high Zn concentration was investigated by monitoring the photoluminescence and 

Raman scattering. Analysis of the photoluminescence spectra of structures which had 

been diffused for various times, and of the photoluminescence spectra from various 

depths below the sample surface, made it possible to describe the physical processes 

which occurred during Zn diffusion. Column-III vacancies were created at the sample 

surface and diffused into the bulk of the sample, where they were filled by other defects. 

By using X-ray diffraction techniques, an expansion of the lattice constant in the region 

beyond the Zn diffusion front was detected. This was attributed to a supersaturation of 

column-III interstitials. During the incorporation of Zn into the crystal lattice, column-III 

interstitials were generated. These were suggested to be responsible for the enhancement 

of Al-Ga interdiffusion. An important role was played by the electric field at the p-n 

junction that was formed by Zn diffusion. That is, the negatively charged column-III 

vacancies and the positively charged column-III interstitials were confined to the n and p 

sides of the p-n junction, respectively. These results provided evidence for a selfinterstitial 

mechanism of Zn diffusion-induced disordering in these multiple quantum well 

structures. 

N.H.Ky, J.D.Ganière, M.Gailhanou, B.Blanchard, L.Pavesi, G.Burri, D.Araújo, 

F.K.Reinhart: Journal of Applied Physics, 1993, 73[8], 3769-81 

[446-106/107-079] 

GaAs/AlGaAs: Zn Diffusion 

An investigation was made of the diffusion of Zn in n-GaAs/Zn-AlGaAs/Se-AlGaAs 

structures during the growth of n-GaAs layers which were doped with Se and Si. It was 

found that the diffusion of Zn in these structures depended strongly upon the type of 

dopant as well as upon the carrier concentration in the n-GaAs layer. The amount of Zn 

which diffused into both n-GaAs and Se-AlGaAs layers was much smaller for a Si-doped 

368


GaAs layer than for a Se-doped layer. The slower diffusion which occurred during the 

growth of Si-doped GaAs layers in these structures could be reasonably well explained by 

modifying a model in which interstitial Ga which diffused from the n-GaAs layer and into 

the Zn-AlGaAs layer was supposed to kick out substitutional Zn. The density of 

interstitial Ga in the Si-doped GaAs layer could be lower than in the Se-doped GaAs 

because the interstitial Ga atoms replaced Si which occupied the column-III site. This was 

not the case for Se-doped GaAs, where Se occupied the column-V site. 

N.Fujii, T.Kimura, M.Tsugami, T.Sonoda, S.Takamiya, S.Mitsui: Journal of Crystal 

Growth, 1994, 145[1-4], 808-12 

[446-119/120-201] 

GaAs/GaAlAs: Zn Diffusion 

Heterostructures of GaAs/Ga 0.7 Al 0.3 As, which contained Zn and Se as intrinsic p and n 

dopants, were subjected to combined Be and O implantation. Rapid thermal annealing 

then resulted in the enhanced out-diffusion of Zn. The Se dopant profile remained 

essentially unchanged. The atomic profile of Zn could be related to the microscopic 

defect distributions. A change in the photoluminescence spectrum, due to overcompensation 

of the n-doped GaAs and GaAlAs layers, was observed. Annealing without 

preceding implantation had no effect upon the Zn atomic profile. 

T.Humer-Hager, R.Treichler, P.Wurzinger, H.Tews, P.Zwicknagl: Journal of Applied 

Physics, 1989, 66[1], 181-6 

[446-74-026] 


Multiple quantum-well GaAs/Ga 0.8 Al 0.2 As structures which were uniformly Si-doped, to 

concentrations ranging from 10 17 to 10 19 /cm 3 , were grown by means of molecular-beam 

epitaxy in order to study the effects of the background Si dopant level upon the Zn 

diffusion-induced disordering. After Zn diffusion (575C, 4 or 16h), cleaved wedges of the 

samples were investigated by means of secondary-ion mass spectrometry and 

transmission electron microscopy. The results showed that completely and partially 

disordered regions were always behind the Zn diffusion front. A dependence of the 

effective Zn diffusivity and of the disordering rate of the structure upon the background 

Si dopant level was observed. The effective Zn diffusivity and the disordering rate 

significantly decreased with increasing background Si concentration. Before Zn diffusion, 

the photoluminescence spectra of Si-doped structures exhibited an increase in intensity of 

the Si donor column-III vacancy complex emission band with increasing Si dopant level. 

This indicated that the concentration of column-III vacancies in the structures increased 

as the background Si concentration was increased. After Zn diffusion, a large decrease in 

intensity of the column-III vacancy-related emission band was observed in the 

photoluminescence spectra from Zn-diffused regions. A model that was based upon the 

so-called kick-out mechanism of Zn diffusion was proposed in order to explain the effect 

of the background Si doping level upon the effective Zn diffusivity. The model showed 

that the effective Zn diffusivity was controlled by the concentration of column-III 

interstitials behind the Zn diffusion front, and by the donor concentration in the sample. 

369

Zn GaAs Muons 

Column-III interstitials were generated during the incorporation of Zn into the crystal 

lattice. The supersaturation of these interstitials behind the Zn diffusion front was 

responsible for the enhancement of Al-Ga interdiffusion. Because column-III interstitials 

and column-III vacancies could mutually annihilate, the concentration of column-III 

interstitial and column-III vacancies in Zn-diffused regions decreased with increasing Si 

doping level; thus leading to a retardation of Zn diffusion into the structure. A decrease in 

effective Zn diffusivity, caused by an increase in the donor concentration of the samples, 

was also demonstrated. The results revealed effects of the Fermi level and of interactions 

between point defects during the Zn diffusion-induced disordering of GaAs/AlGaAs 

multi-layered structures. 

N.H.Ky, J.D.Ganière, F.K.Reinhart, B.Blanchard: Journal of Applied Physics, 1996, 

79[8], 4009-16 

[446-134/135-136] 


Secondary-ion mass spectrometry and photoluminescence methods were used to study Zn 

diffusion into GaAs/Ga 0.8 Al 0.2 As multi quantum-well structures which were uniformly 

doped with Si to concentrations of between 10 17 and 10 19 /cm 3 . The secondary-ion mass 

spectrometry profiles which were measured after Zn diffusion at 575C revealed a large 

effect of the background Si doping level upon the Zn diffusion process and upon Zn 

diffusion-induced disordering of the multi quantum-well structures. It was found that the 

Zn diffusivity and the disordering rate significantly decreased with increasing Si 

background concentration. Before Zn diffusion, the photoluminescence spectra of the 

multi quantum-well samples revealed an increased intensity of the Si donor-V III complex 

emission band with increased Si doping level. This indicated that the V III concentration in 

the multi quantum-well structures increased as the Si background concentration increased. 

After Zn diffusion, a large decrease in the intensity of the V III -related emission band was 

detected in the photoluminescence spectra which were obtained from Zn-diffused regions. 

The results were explained in terms of the kick-out mechanism of Zn diffusion. During 

the incorporation of Zn into the crystal lattice, column-III interstitials were generated. The 

supersaturation of I III behind the Zn diffusion front resulted in an enhancement of Al-Ga 

interdiffusion in the Zn-diffused region. Since I III and V III could mutually annihilate, a 

reduction in I III concentration occurred in the Zn-diffused region of Si-doped samples 

which contained a high V III concentration. As a result, Zn diffusion and disordering of the 

multi quantum-well structures were retarded with increasing Si-doping level. 

N.H.Ky: Materials Science Forum, 1995, 196-201, 1643-8 

Muons 

[446-127/128-139] 

GaAs: Muon Diffusion 

Measurements were made of the dynamics of negatively charged muonium in heavily Sidoped 

material at temperatures ranging from 295 to 1000K. The muonium began to 

370

Muons GaAs Muons 

diffuse, at temperatures above 500K, at a hopping rate which was described by an attempt 

frequency of 5.6 x 10 12 /s and an activation energy of 0.73eV. At temperatures above 

700K, relaxation from charge-state fluctuations was observed. The data implied that Mu - 

to Mu o conversion occurred, via the alternating capture of holes and electrons. This 

established that Mu was a deep recombination center. Similar dynamics were expected 

for the isolated H- center in n-type material. 

K.H.Chow, B.Hitti, R.F.Kiefl, S.R.Dunsiger, R.L.Lichti, T.L.Estle: Physical Review 

Letters, 1996, 76[20], 3790-3 

[446-136/137-111] 


The results of investigations of transitions among the sites and charge states of muonium 

were summarized. Energy parameters were determined for the full description of 

muonium dynamics and were found to correlate well with available data for H. A model 

was developed which accounted for all of the major features which were observed. Its 

validity for a wide range of dopant concentrations, using a single set of parameters, 

reflected its predictive strength. The near equality of the energy parameters for muonium 

as compared with those which were available for H, strongly implied that the results for 

muonium dynamic behavior should be applicable to H, with very little change. The model 

could be applied to all tetrahedrally coordinated semiconductors, with few modifications, 

and served as a basis for the understanding of muonium dynamics in GaAs and other 

materials. Differences in H properties could be understood by examining materialspecific 

deviations from the basic model. 

R.L.Lichti, C.Schwab, T.L.Estle: Materials Science Forum, 1995, 196-201, 831-6 

[446-127/128-119] 


It was noted that the presence of Si donors had a marked effect upon the charge state and 

diffusion of muonium at the tetrahedral interstitial site (Mu o T ), while it had a relatively 

weak effect upon bond-center muonium (Mu o BC ). In metallic Si-doped material, the 

highly mobile Mu o T center which was observed in non-metallic material was replaced by 

a charged species (probably Mu - ) which had a diffusion rate that was smaller than that of 

the Mu o T center. The Mu o BC center was metastable, and its transition to Mu o T or Mu - 

centers (depending upon dopant concentration) was observed at 50 to 150K. The results 

o 

also suggested that the Mu T state underwent fast spin-exchange interaction with 

conductive carriers before the transition: Mu o T → Mu - . 

R.Kadono, A.Matsushita, K.Nagamine, K.Nishiyama, K.H.Chow, R.F.Kiefl, 

A.MacFarlane, D.Schumann, S.Fujii, S.Tanigawa: Physical Review B, 1994, 50[3], 1999- 

2002 

[446-115/116-118] 

371

Muons GaAs General 


By measuring the muon spin relaxation rate of muonium as it moved from site to site in a 

host lattice with nuclear spins, it was possible to derive an average correlation time, t, for 

fluctuations in the nuclear hyperfine field which acted upon the unpaired electron. 

Provided that the motion was incoherent (diffusive), t was inversely proportional to the 

diffusion constant. Such measurements were reported here for isotropic muonium at 

temperatures ranging from 20mK to 300K. Large differences were found in the muonium 

motion, and were most marked at low temperature; where l/t became temperatureindependent 

below about 10K in GaAs. A study was also made of the diffusive behavior 

of muonium in various GaAs samples in order to detect the possible effect of slight 

crystalline imperfections. 

W.Schneider, R.F.Kiefl, E.J.Ansaldo, J.H.Brewer, K.Chow, S.F.J.Cox, S.A.Dodds, 

R.C.Duvarney, T.L.Estle, E.E.Haller, R.Kadono, S.R.Kreitzman, R.L.Lichti, 

C.Niedermayer, T.Pfiz, T.M.Riseman, C.Schwab: Materials Science Forum, 1992, 83-87, 

569-74 

[446-99/100-065] 

General 

GaAs: Diffusion 

A review was presented of developments in the understanding of self-diffusion and 

impurity diffusion processes in this material; with particular emphasis being placed on the 

inclusion of recent Ga-isotope diffusion data. Specific diffusion mechanisms were 

suggested for C, P, Sb and S; which were all substitutionally dissolved on the As sublattice. 

U.Gösele, T.Y.Tan, M.Schultz, U.Egger, P.Werner, R.Scholz, O.Breitenstein: Defect and 

Diffusion Forum, 1997, 143-147, 1079-94 

[446-143/147-1079] 

GaAs[l]: Electromigration 

A new experimental technique was proposed which permitted the measurement of the 

effective mobility of solute elements in III-V solutions. It was based upon the concept of 

an electro-epitaxial growth experiment in which the Peltier effect contribution to growth 

was eliminated by adjusting the current flow direction so that it was parallel to the 

substrate surface. Thus, only diffusion and electrotransport contributed to growth. 

Z.R.Zytkiewicz: Journal of Crystal Growth, 1987, 82[4], 647-51 

[446-55/56-168] 


During the fabrication of refractory gate MESFET devices, sputter deposition of a WSi x 

gate and reactive ion etching of the gate pattern could lead to surface damage and 

contamination. In order to study these effects, material with a shallow Si implant was 

372

General GaAs General 

subjected to reactive ion etching alone, or to both WSi x sputter deposition and reactive 

ion etching, before annealing. The surface damage, due to WSi x sputter deposition and 

reactive ion etching at self-bias under 200V, was healed by capped (SiN x ) furnace 

annealing at 800C. It was found that sheet resistance and Hall mobility measurements 

could be correlated with the diffusion of compensating impurities into the bulk. 

Secondary ion mass spectrometry profiles indicated that the major contaminants (Fe, Cr, 

Ni, Cu, V) were already present in the W targets and were thus present in the WSi x 

layers. These contaminants were left on the surface of the GaAs after gate reactive ion 

etching and were driven into the bulk during capped annealing. An HCl etch was found to 

remove the contaminants; thus resulting in lower sheet resistances for implanted and 

processed GaAs. Refractory gate sub-micron MESFET devices which were fabricated by 

using an HCl etch after gate reactive ion etching exhibited a reduced access resistance. 

H.Baratte, A.J.Fleischman, G.J.Scilla, T.N.Jackson, H.J.Hovel, F.Cardone: Journal of the 


[446-78/79-010] 


A long-term reliable spun-on source was developed for open-tube p + diffusion into III/V 

compound semiconductors. The source consisted of an aqueous/alcoholic solution of 

zircon alcoholate, doped with zinc chloride. After spinning and drying in air at 300C, a 

glassy film of ZrO 2 :ZnO on the semiconductor surface acted as a solid-state source for 

subsequent diffusion. This solution exhibited an excellent long-term durability. 

Furthermore, the thermal expansion coefficient of zirconia was well matched to that of 

most III/V compound semiconductors and yielded very stable films, even at high 

temperatures. This permitted essentially stress-free diffusion. Diffusion into GaAs was 

carried out at temperatures ranging from 600 and 700C. The hole concentrations and 

diffusion coefficients which were obtained by using this source were rather close to those 

obtained by using other diffusion techniques. The simpler handling and long-term 

durability of zirconia-based solutions offered significant advantages over other p + - 

diffusion techniques which were used in the preparation of III/V compound 

semiconductors. 

G.Franz, M.C.Amann: Journal of the Electrochemical Society, 1989, 136[8], 2410-3 

[446-70/71-105] 


Deep-level transient Fourier spectroscopic, Hall-effect and DSL etching techniques were 

used to analyze initially semi-insulating bulk samples after annealing at temperatures of 

between 800 and 1100C under As pressures of between 0 and 3bar. It was found that the 

electrical resistivity, electronic mobility and EL2 concentration increased with increasing 

As pressure at all temperatures. Initially n-type samples were converted to p-type at 

pressures below 0.5bar and temperatures above 1000C. It was noted that As precipitates 

disappeared at high temperatures, and reappeared during low temperature annealing. In 

order to explain the observations, it was proposed that long-range rapid As interstitial 

transport occurred between the surface and the bulk, together with short-range Ga 

373

General GaAs Surface 

vacancy migration involving dislocations as sources and sinks. It was suggested that a 

discrepancy, with regard to very small reported tracer diffusion coefficients, could be 

resolved by assuming that rapidly diffusing marked interstitials which entered the lattice 

at the crystal surface tended to exchange sites with unmarked lattice species. They then 

became immobile. 

M.Noack, K.W.Kehr, H.Wenzl: Journal of Crystal Growth, 1997, 178, 438-44 

[446-152-0374] 


Al 

GaAs: Al Surface Diffusion 

During the molecular beam epitaxial growth of GaAs on the vicinal (100) surface of 






of Sn. The presence of Sn increased the surface mobility of Al adatoms on (100) AlAs 

surfaces; as indicated by the annealing behavior of the AlAs surface at 600C. 


[446-93/94-001] 


Misoriented samples which were tilted by 1 or 2º away from the (001) plane and towards 

[110] or (111), or towards [¯110] or (111), were prepared. The critical temperature at 

which a transition in growth mode occurred was studied by using reflection high-energy 

electron diffraction methods during the growth of AlAs on GaAs vicinal surfaces. The 

critical temperature of AlAs was higher than that of GaAs, thus indicating that the stepflow 

growth of AlAs occurred at a higher temperature. By combining these data with a 

theory that took account of 2-dimensional nucleation and surface diffusion, the surface 

diffusion length of Al was deduced. It was found to be greater than that of Ga for both 

types of substrate surface. By taking account of chemical equilibrium at the step edge of 

each surface, it was predicted that the surface diffusion length of Al should be 

anisotropic. 

M.Tanaka, T.Suzuki, T.Nishinaga: Japanese Journal of Applied Physics, 1990, 29[5], 

L706-8 

[446-76/77-006] 


The adsorption and migration characteristics of Al atoms on (100) surfaces were 

investigated by using reflection high-energy electron diffraction and the alternate 

374

Surface GaAs Surface 

deposition of Al and As 4 onto the growing surface. All of the results were attributed to 

the existence of differing migration velocities for various atoms. 

Y.Horikoshi, H.Yamaguchi, M.Kawashima. Japanese Journal of Applied Physics, 1989, 

28[8], 1307-11 

[446-72/73-001] 

As 

GaAs: As Surface Diffusion 

The incorporation diffusion length of surface-migrating adatoms on molecular beam 

epitaxial material was studied theoretically and experimentally. By using microprobe 

reflection high-energy electron diffraction and scanning electron microscopic techniques, 

it was demonstrated that the incorporation diffusion length depended strongly upon the 

As partial pressure, and was of the order of 1µ. This behavior was explained by assuming 

that the Ga flux which entered the step exceeded the As flux. When the Ga flux which 

entered the step was lower than that of As, the incorporation diffusion length decreased to 

that of the step separation; and was then typically of the order of tens of nm. It was 

concluded that the balance of Ga and As fluxes which entered the step edge determined 

the incorporation diffusion length of Ga. 

T.Nishinaga, X.Q.Shen: Applied Surface Science, 1994, 82-83, 141-8 

[446-123/124-161] 

Ga 

GaAs: Ga Surface Diffusion 

During the molecular beam epitaxial growth of GaAs on the vicinal (100) surface of 






of Sn. This indicated that the Ga mobility had increased. 


[446-93/94-001] 


The dependence of Ga adatom surface diffusion upon the As flux during molecular beam 

epitaxial growth was investigated. Variations of the growth rate of GaAs layers, grown 

onto (001) surfaces adjacent to (111) surfaces, were measured by means of scanning 

microprobe reflection high-energy electron diffraction. The surface diffusion length was 

375


deduced from the variations in the growth rate. It was found that the surface diffusion 

length of the Ga adatoms became larger under a lower As flux. 

M.Hata, A.Watanabe, T.Isu: Journal of Crystal Growth, 1991, 111[1-4], 83-7 

[446-91/92-007] 


The microscopic details of Ga adatom diffusion upon an As-stabilized (001) surface were 

investigated by using an ab initio pseudopotential method. The results showed that Ga 

adatoms diffused on the surface by passing through the missing As dimer rows. A 

comparison with scanning tunnelling microscopic experiments during molecular beam 

epitaxial growth suggested that a low pressure of As increased surface Ga adatom 

diffusion due to the creation of a continuous Ga adatom diffusion path. This conclusion 

was consistent with the observation that low-temperature growth was possible via 

migration-enhanced epitaxy in which As and Ga sources were supplied alternately. 

K.Shiraishi: Applied Physics Letters, 1992, 60[11], 1363-5 

[446-86/87-010] 


Misoriented samples which were tilted by 1 or 2º away from the (001) plane and towards 

[110] or (111), or towards [¯110] or (111), were prepared. The critical temperature at 

which a transition in growth mode occurred was studied by using reflection high-energy 

electron diffraction methods during the growth of AlAs on GaAs vicinal surfaces. The 

critical temperature of AlAs was higher than that of GaAs, thus indicating that the stepflow 

growth of AlAs occurred at a higher temperature. The surface diffusion length of Ga 

was found to be shorter than that of Ga for both types of substrate surface. 

M.Tanaka, T.Suzuki, T.Nishinaga: Japanese Journal of Applied Physics, 1990, 29[5], 

L706-8 

[446-76/77-006] 


An experimental study of RHEED intensity oscillations was performed on (111)B 

substrates which were misoriented by 1 or 2º towards [110], [2¯1¯1], or [¯211] orientations. 

The behavior of the RHEED oscillations on (111)B was similar to that on (001). 

However, the temperature at which the RHEED intensity oscillations began to appear on 

(111)B was lower than that on (001). The surface diffusion length of Ga on (111)B was 

evaluated by taking account of the supersaturation ratio of adatoms on the terrace. 

T.Shitara, E.Kondo, T.Nishinaga: Journal of Crystal Growth, 1990, 99, 530-4 

[446-76/77-008] 


Surface diffusion during molecular beam epitaxy was studied. Firstly, the mode transition 

between 2-dimensional nucleation and step flow during molecular beam epitaxial growth 

on vicinal surfaces was studied theoretically and experimentally. The basis of the theory 

was to assume that the transition occurred when the surface supersaturation on the step 

376


terrace became identical to the critical supersaturation for 2-dimensional nucleation. This 

permitted the diffusion length of Ga to be calculated at the experimentally determined 

critical temperature for the mode transition. It was found that the diffusion length 

increased, as the temperature decreased, due to an increased residence time. Also, the 

diffusion length on (111)B was longer than that on (001) when the same formation energy 

for 2-dimensional nuclei was assumed for both surfaces. The theory gave good agreement 

with the experimental data, and it was concluded that surface diffusion was one of the 

most important processes which controlled molecular beam epitaxial growth and impurity 

incorporation. 

T.Nishinaga, T.Shitara, K.Mochizuki, K.I.Cho: Journal of Crystal Growth, 1990, 99, 482- 

90 

[446-76/77-009] 


The adsorption and migration characteristics of Ga atoms on (100) surfaces were 

investigated by using reflection high-energy electron diffraction and the alternate 

deposition of Ga and As 4 onto the growing surface. Excess Ga deposition onto the surface 

produced Ga clusters or droplets on the first Ga layer. These dissolved very quickly after 

As 4 deposition and formed flat GaAs layers when the number of Ga atoms was near to 2 

or 3 times the surface site number. All of the results were attributed to the existence of 

differing migration velocities for various atoms. 

Y.Horikoshi, H.Yamaguchi, M.Kawashima. Japanese Journal of Applied Physics, 1989, 

28[8], 1307-11 

[446-72/73-001] 


The generation of extra facets on ridge-type triangles, with (001)-, (110)- and (201)- 

related equivalent slopes on GaAs (111) A substrates, and stripes running in [¯110], [110] 

and [100] directions on (001) substrates, was investigated during the molecular beam 

epitaxy of GaAs/AlGaAs multi-layers. By investigating local variations in the layer 

thickness in regions adjacent to extra (114)A, (110) and (¯1¯1¯1)B facets which were 

common to the (111)A and (001) patterned substrates, and extra facets which were related 

to the respective substrates and growth rates of the facets relative to the growth rate on 

the substrate plane, the orientation-dependent Ga surface diffusion lengths were 

determined. They increased in the order: (001), (¯1¯1¯1)B-related, (111)A-related and (110). 

Or, diagrammatically, 

(001) < (¯1¯1¯1)B < (159) < (110) 

(¯1¯1¯3)B (¯3¯3¯1)B (114)A 

(013)B (111)A 

(113)B 

T.Takebe, M.Fujii, T.Yamamoto, K.Fujita, T.Watanabe: Journal of Applied Physics, 

1997, 81[11], 7273-8 

[446-152-0377] 

377



The migration potentials of Ga adatoms near to step edges on the c(4 x 4) surface were 

investigated by using an empirical interatomic potential and an energy term that 

accounted for charge redistribution on the surface. The energy term, as a function of the 

number of electrons which remained in the Ga dangling bonds, was deduced from firstprinciples 

calculations. The latter results implied that lattice sites along the A-type step 

edges were stable for Ga adatoms, whereas no preferential adsorption site was found near 

to B-type step edges. This was because the number of electrons which remained in the Ga 

dangling bond was reduced by Ga adatoms that occupied lattice sites along A-type step 

edges, rather than being unchanged by those which occupied lattice sites near to B-type 

step edges. 

T.Ito, K.Shiraishi: Japanese Journal of Applied Physics, 1996, 35[2-8B], L1016-8 

[446-138/139-076] 


The migration potentials of Ga adatoms near to kink and step edges were qualitatively 

investigated by using empirical inter-atomic potentials and an energy term. The latter 

term, as a function of the number of electrons that remained in the Ga dangling bond, was 

deduced from first-principles pseudopotential calculations. The calculated results implied 

that the lattice sites in the missing dimer row were favorable for Ga adatoms on the 

GaAs(001)-(2 x 4)β2 surface. This was because the formation of Ga dimers reduced the 

number of electrons that remained in Ga dangling bonds. Lattice sites in the missing 

dimer row, near to a kink and a B-type step edge, were stable locations for a Ga adatom. 

On the other hand, no preferential adsorption site was found near to an A-type step edge. 

This was because a Ga adatom in the missing dimer row near to a kink and a B-type step 

edge was slightly stretched by an As atom and As-dimer on the plane that was 1 layer 

below, rather than being strongly stretched by two As-dimers near to an A-type step edge. 

T.Ito, K.Shiraishi: Japanese Journal of Applied Physics, 1996, 35[2-8A], L949-52 

[446-138/139-077] 


A valence force field was optimized in order to reproduce the phonon dispersion curves 

of crystalline GaAs and derived interaction energies of Ga adatoms on the (001) surface. 

Calculations of the diffusion constant of isolated Ga atoms on the GaAs surface were 

performed by means of molecular dynamics simulations. All of the bulk Ga and As 

atoms, and the adsorbed Ga atoms, were completely free to move and no normalization of 

the velocity was performed after the trajectory had begun. Averages were taken of the 

results of hundreds of such trajectories for each temperature. Surface diffusion constants 

were then obtained from the (001) in-plane components. It was found that the data could 


378


D(cm 2 /s) = 2.41 x 10 -5 exp[-0.0971(eV)/kT] 

A.Palma, E.Semprini, A.Talamo, N.Tomassini: Journal of Crystal Growth, 1995, 150[1- 

4], 180-4 

[446-127/128-118] 


Facets of (110)-type were formed, on mesa edges which defined (100)-(110) facet 

structures, by the molecular beam epitaxial growth of GaAs onto [001]-mesa stripes on 

(100) GaAs substrates. The surface diffusion length of Ga adatoms along the [010] 

direction on the mesa stripes was estimated for various growth conditions by means of in 

situ scanning microprobe reflection high-energy electron diffraction. By using these 

values, and the corresponding growth rate on the (110) GaAs facets, the diffusion length 

on the (110) plane was deduced. It was found that the Ga diffusion length on the (110) 

plane was greater than that on the (100) and (111)B planes. The long diffusion length on 

the (110) plane was explained in terms of the particular surface reconstruction on this 

plane. 

M.López, Y.Nomura: Journal of Crystal Growth, 1995, 150[1-4], 68-72 

[446-127/128-118] 


Systematic measurements were made of the surface diffusion lengths of Ga adatoms 

during the molecular beam epitaxy of this material in the presence of H or H 2 . The spatial 

variation in the growth rate on the (100) surface adjacent to the (111)A surface was 

deduced from the period of reflection high-energy electron diffraction intensity 

oscillations. The surface diffusion length of Ga adatoms, as estimated from the spatial 

variation in the growth rate, increased with increasing H or H 2 pressure. It also increased 

as the substrate temperature was increased at a given H or H 2 pressure. The diffusion 

length in the case of H was greater than that in the case of H 2 . 

Y.Morishita, Y.Nomura, S.Goto, Y.Katayama: Applied Physics Letters, 1995, 67[17], 

2500-2 

[446-125/126-121] 


The incorporation diffusion length of surface-migrating adatoms on molecular beam 

epitaxial material was studied theoretically and experimentally. By using microprobe 

reflection high-energy electron diffraction and scanning electron microscopic techniques, 

it was demonstrated that the incorporation diffusion length depended strongly upon the 

As partial pressure, and was of the order of 1µ. This behavior was explained by assuming 

that the Ga flux which entered the step exceeded the As flux. When the Ga flux which 

entered the step was lower than that of As, the incorporation diffusion length decreased to 

that of the step separation; and was then typically of the order of tens of nm. It was 

379


concluded that the balance of Ga and As fluxes which entered the step edge determined 

the incorporation diffusion length of Ga. 

T.Nishinaga, X.Q.Shen: Applied Surface Science, 1994, 82-83, 141-8 

[446-123/124-161] 


Two-dimensional nuclei of GaAs which were grown on a singular (001) surface by 

metalorganic chemical vapor deposition were observed by means of high-vacuum 

scanning tunnelling microscopy. The nuclei extended in the [110] direction, which was 

opposite to that of molecular beam epitaxial growth. The nucleus number density was 

obtained from scanning tunnelling microscopic images, and the relationship between the 

density and the surface diffusion coefficient of Ga species was estimated by simulating 

growth at the surface. The surface diffusion coefficient was estimated to be equal to about 

10 -7 cm 2 /s at 530C. 

M.Kasu, N.Kobayashi: Journal of Crystal Growth, 1994, 145[1-4], 120-5 

[446-119/120-191] 


Quantitative features of Ga atom surface migration (migration length, migration time, 

hopping frequency) were obtained by the Monte Carlo simulation of molecular beam 

epitaxial growth of GaAs on flat and stepped (001) surfaces under various growth 

conditions. Changes in these parameters during growth, and the relationship between 

migration, surface roughness and bulk defect concentration, were considered. 

N.V.Peskov: Surface Science, 1994, 306[1-2], 227-32 

[446-119/120-191] 


Spatial variations in the growth rate on mesa-etched GaAs (¯1¯1¯1)B substrates during the 

molecular beam epitaxy of GaAs were deduced from the period of reflection high-energy 

electron diffraction intensity oscillations by using in situ scanning microprobe methods. 

The surface diffusion length of Ga adatoms on the (¯1¯1¯1)B surface was deduced from the 

spatial variation in the growth rate. The surface diffusion length on the (¯1¯1¯1)B surface 

increased as the substrate temperature was increased or the As pressure was decreased. A 

typical value of the diffusion length was about 10µ, at a substrate temperature of 580C 

and an As pressure of 5.7 x 10 -4 Pa. This was an order of magnitude larger than that on the 

(100) surface, in the [011] direction. The activation energy of the surface diffusion length 

changed with surface reconstruction. Anisotropic diffusion, which had been reported for 

the (100) surface, was not observed on the (¯1¯1¯1)B surface. 

Y.Nomura, Y.Morishita, S.Goto, Y.Katayama, T.Isu: Applied Physics Letters, 1994, 

64[9], 1123-5 

[446-115/116-117] 

380



The faceted surface morphologies of homo-epitaxial films which had been grown onto 

exactly (¯1¯1¯1)-oriented GaAs substrates in the √19 x √19 regime were studied with the aid 

of an atomic force microscope. The facets were composed of 3 vicinal surfaces which 

were tilted by about 2° towards the [2¯1¯1], [¯12¯1], and [¯1¯12] directions. The diffusion 

length was deduced from the surface morphologies, and was found to be equal to some 

hundreds of nm. It was comparable to the diffusion length on (100) surfaces which were 

grown under the same conditions. It was concluded that facet formation on GaAs (¯1¯1¯1) 

films was unlikely to be caused by a lower surface mobility. 

K.Yang, L.J.Schowalter, T.Thundat: Applied Physics Letters, 1994, 64[13], 1641-3 

[446-115/116-117] 


The As pressure dependence of Ga adatom surface diffusion during molecular beam 

epitaxy onto non-planar substrates was investigated. By using in situ scanning 

microprobe reflection high-energy electron diffraction techniques, the distribution of the 

growth rate of GaAs on the (001) surface near to the edge of the (111)A or (111)B sidewall 

was measured under various As pressures. It was found that the surface diffusion 

length of Ga adatom incorporation on the (001) surface, as deduced from the growth rate 

distribution, was of the order of µ and exhibited a marked dependence upon the As 

pressure. A simple model which was based upon 1-dimensional surface diffusion was 

used to estimate the lifetime for Ga adatom incorporation on other surfaces. 

X.Q.Shen, D.Kishimoto, T.Nishinaga: Japanese Journal of Applied Physics, 1994, 33[1- 

1A], 11-7 

[446-113/114-011] 


The As pressure dependence of the surface diffusion of Ga adatoms on molecular beam 

epitaxial (111) B-(001) mesa-etched substrates was investigated by using in situ scanning 

microprobe reflection high-energy electron diffraction techniques. It was observed, for 

the first time, that the direction of Ga adatom migration from, or to, the (111)B side-wall 

changed; depending upon the As pressure. Moreover, the diffusion length of Ga adatoms 

on the (001) surface, in the [¯110] direction, was found to depend exponentially upon the 

As pressure. However, it was independent of the direction of lateral migration. The 

diffusion length of Ga adatoms on the (001) surface, in the [¯110] direction, varied from 

about 0.25 to 1.2µ at 600C, within the present range of As pressures. It was suggested 

that the lifetime of Ga adatoms, before incorporation into the crystal, depended strongly 

upon the As pressure. 

X.Q.Shen, T.Nishinaga: Japanese Journal of Applied Physics, 1993, 32[2-8B], L1117-9 

[446-109/110-029] 

381


GaAs/AlGaAs: Ga Surface Diffusion 

Thickness variations in quantum wells which had been grown on patterned substrates, by 

means of molecular beam epitaxy, were analyzed by using spatially and spectrally 

resolved low-temperature cathodoluminescence methods. In the case of lower and upper 

(100) facets which were joined by an angled (311)A facet, relative increases in the 

quantum well thickness, of up to about 6% and 20% respectively, were observed in the 

vicinity of the facet intersection. The Ga adatom migration length obeyed an exponential 

behavior, and ranged from 0.001 to 0.002mm on both the lower and upper (100) facets. It 

was independent of the quantum well thickness. The present migration length was some 

orders of magnitude greater than that which had previously been reported for Ga adatoms 

during molecular beam epitaxial growth. 

S.Nilsson, E.Van Gieson, D.J.Arent, H.P.Meier, W.Walter, T.Forster: Applied Physics 

Letters, 1989, 55[10], 972-4 

[446-72/73-026] 

K 

GaAs: K Surface Diffusion 

A self-consistent semi-empirical molecular orbital method was used to determine whether 

the adsorption properties of K atoms, and the formation of K adsorbate chains or clusters 

in the low-coverage regime, could be affected by the nature of the semiconductor surface 

(that is, perfect or stepped). It was found to be possible to determine the microscopic 

structures of monatomic and diatomic K molecules on perfect and stepped (110) GaAs 

surfaces. The results for K adsorption on the perfect GaAs(110) surface were consistent 

with scanning tunnelling microscopic observations of Na on (110) GaAs; with the stable 

site for K being the bridge site which encompassed one Ga and two As surface atoms. 

The equilibrium geometry for diatomic K involved the second K atom occupying the 

next-nearest neighbor bridge site; thus strongly supporting the formation of an open linear 

structure parallel to the zig-zag surface atomic chains. The calculated K-K distance in this 

equilibrium configuration was 0.802nm. This was similar to the Na-Na distance (0.8nm) 

which was deduced from scanning tunnelling microscopy experiments. The results for the 

stepped (110) GaAs surface suggested that a step was unlikely to assist the clustering of 

K atoms. However, the formation of the linear adsorbate chain appeared to be influenced 

more by the orientation of the steps. On the perfect surface, the K adsorbates were bound 

more strongly at steps than at bridge sites. 

G.S.Khoo, C.K.Ong: Journal of Physics - Condensed Matter, 1993, 5[36], 6507-14 

[446-106/107-037] 

382


Ti 

GaAs: Ti Surface Diffusion 

Scanning tunnelling microscopy was used to study the morphology of the Ti/(110)GaAs 

interface after atom deposition at 300K. The results indicated that thermally activated 

surface diffusion was minimal during the growth process. This was because of the high 

activation energy which was imposed by chemical bonding. On the other hand, surface 

hopping was observed because of the dynamics which were associated with the cooling of 

impinging atoms. It was noted that single Ti atoms formed stable bonds with the 

substrate, but unique bonding sites were not distinguishable and the areas around these 

sites were not disrupted. 

Y.N.Yang, B.M.Trafas, Y.S.Luo, R.L.Siefert, J.H.Weaver: Physical Review B, 1991, 

44[11], 5720-5 

[446-84/85-017] 


GaAs: Surface Diffusion 

The bombardment of (110) samples with Ar + ions of normal incidence, at temperatures of 

300 to 775K, created surface layer defects that usually spanned 1 or 2 unit cells (as 

revealed by scanning tunnelling microscopy). Vacancies which were produced in this way 

diffused via thermal activation to form single-layer vacancy islands. The diffusion of divacancies 

favored [1¯10], and accommodation at islands produced roughly isotropic 

islands. Modelling of this growth process revealed an overall Arrhenius behavior of the 

diffusion, with an activation energy of 1.3eV. Investigations of the surface morphology 

during multi-layer erosion revealed deviations from layer-by-layer removal, with scaling 

exponents of between 0.4 and 0.5 at temperatures of between 626 and 775K. 

R.J.Pechman, X.S.Wang, J.H.Weaver: Physical Review B, 1995, 51[16], 10929-36 

[446-121/122-053] 


It was found that surface migration was effectively enhanced by evaporating Ga or Al 

atoms onto a clean GaAs surface under an As-free atmosphere or low As pressure. This 

characteristic was exploited by alternately supplying Ga and/or Al and As to the substrate 

surface in order to grow atomically-flat GaAs-AlGaAs hetero-interfaces, and also to grow 

high-quality GaAs layers at very low temperatures. The migration characteristics of 

surface adatoms were investigated by using reflection high-energy electron diffraction 

measurements. It was found that differing growth mechanisms operated at high and low 

temperatures. Both mechanisms were expected to yield flat heterojunction interfaces. By 

383


applying this method, GaAs layers could be grown at substrate temperatures of 200 and 

300C, respectively. 

Y.Horikoshi, M.Kawashima, H.Yamaguchi: Japanese Journal of Applied. Physics, 1988, 

27[2], 169-79 

[446-60-001] 


A 90° double-reflection high-energy electron diffraction method was used to carry out a 

study of the morphology of vicinal (001) surfaces during molecular beam epitaxy. The 

technique permitted the simultaneous recording of reflection high-energy electron 

diffraction intensifies in the [¯110] and [110] azimuths. Comparative measurements of 

surfaces with 2° misorientations towards (111)Ga (A surface) or (1¯11)As (B surface) 

showed that, regardless of the step-type and reconstruction anisotropy, recordings of the 

specular beam intensity in the azimuth perpendicular to the steps were dominated by 

changes in the staircase order whereas intensity recordings in the azimuth parallel to the 

steps revealed changes in the step-edge roughness. Measurements which were performed 

over a wide range of substrate temperatures clarified the competition between kinetic 

processes and thermodynamic equilibrium at the length scale which was accessible to 

reflection high-energy electron diffraction techniques. In the case of the A surface, the 

transition between 2-dimensional growth and step-flow growth not only occurred at a 

higher temperature than it did on the B surface, but the disappearance of intensity 

oscillations also occurred at differing substrate temperatures for different azimuths. An 

approximately 20C higher disappearance temperature for the [¯110] azimuth was 

explained in terms of a model that was based upon previous scanning tunnelling 

microscopy results which had revealed an increasing elongation of islands, in the [¯110] 

direction, with increasing substrate temperature. The B surface was more isotropic, and 

therefore no difference in the transition temperature for the 2 azimuths could be detected. 

During growth in the transition range between 2-dimensional and step-flow growth, 

increased terrace-width fluctuations were observed on the B surface whereas the A 

surface became more uniformly stepped. It was concluded that, in the kinetically 

controlled regime, the anisotropic barrier height for downward diffusion of adatoms over 

step edges played an important role in the evolution of surface morphology. At high 

temperatures, the barrier height permitted downward jumps of the adatoms over B-type 

steps but not over A-type steps. Under conditions that were close to thermodynamic 

equilibrium, kinetic smoothing was observed on the A surface as well as on the B surface. 

This indicated that another mechanism became operative upon a change in the energetics 

due to ordering of the steps and disordering of the reconstruction on the terraces. This 

surface was metastable and rapidly recovered (within less than 1s) to give the equilibrium 

bunched surface after interruptions in growth at substrate temperatures above 580C. 

H.Nörenberg, L.Däweritz, P.Schützendübe, K.Ploog: Journal of Applied Physics, 1997, 

81[6], 2611-20 

[446-148/149-175] 

384



An investigation was made of surface kinetics, during metalorganic vapor-phase epitaxial 

growth, by means of high-vacuum scanning tunnelling microscopic observations of 2- 

dimensional nuclei and denuded zones. Monte Carlo simulations were carried out which 

were based upon the solid-on-solid model. Two-dimensional nucleus densities were used 

to deduce that the surface diffusion coefficient of GaAs was equal to 2 x 10 -6 cm 2 /s at 

530C. The activation energy for migration was estimated to be 0.62eV. The 2- 

dimensional nucleus size in the [110] direction was about twice that in the [¯110] 

direction. This anisotropy was attributed mainly to a difference in the lateral sticking 

probabilities between steps along [¯110] and those along [110]. The ratio of the sticking 

probabilities was estimated to be greater than 3:1. The denuded zone widths on the upper 

terraces were some 2 times wider than those on the lower terraces. This suggested that the 

sticking probability at descending steps was 10 to 300 times larger than the probability at 

ascending steps. 

M.Kasu, N.Kobayashi: Journal of Crystal Growth, 1997, 170, 246-50 

[446-141/142-093] 


The surface diffusion of group-III atom incorporation during molecular beam epitaxial 

growth was considered. Firstly, the diffusion length for incorporation on the (001) top 

surface, with (111)A or (411)A side surfaces on V grooves, was studied. It was shown 

that the diffusion length took the same value for both cases and was inversely 

proportional to the As pressure. However, the diffusion length of Ga on (111)B exhibited 

an inverse parabolic dependence of the As pressure. It was suggested that, on the (001) 

surface, two As 4 molecules met to furnish active As atoms for growth. On the other hand, 

the behavior of the As 4 molecule on the (111)B surface remained unclear. The ratio of the 

surface diffusion coefficients on (111)B and (001) was calculated. It was found that the 

ratio took a value of about 140. Using this ratio, the incorporation lifetimes on (111)B 

and (001) surfaces were calculated as functions of the As pressure. It was found that the 

curves of incorporation lifetime intersected at the As pressure where flow inversion 

occurred. 

T.Nishinaga, X.Q.Shen, D.Kishimoto: Journal of Crystal Growth, 1996, 163[1], 60-6 

[446-138/139-078] 


The migration of anion and cation vacancies on the (110) surface was studied by means 

of scanning tunnelling microscopy. Marked asymmetries were found in the direction and 

bias-polarity dependence of the migration probability. This indicated the importance of 

the bonding topology at the surface. The asymmetry showed that vacancy motion was 

driven by the recombination of carriers, injected by the scanning tunnelling microscope 

tip, with carriers from the bulk. The impact-parameter dependence of the reaction cross- 

385


section showed that this injection occurred via resonant tunnelling into dangling-bond 

defect states. 

G.Lengel, J.Harper, M.Weimer: Physical Review Letters, 1996, 76[25], 4725-8 

[446-136/137-111] 


Reconstruction during chemical beam etching with AsCl 3 was studied. The detection of 

reflection high-energy electron diffraction intensity oscillations indicated the occurrence 

of a planar etching mode in the initial stages. Its change to 3-dimensional etching could 

be understood in terms of suppressed cation diffusion during etching. It was concluded 

that a suitable choice of etching parameters, which would enhance cation diffusion, 

would lead to a smooth etching morphology. The effectiveness of etch cleaning depended 

upon the planarity of the surface during etching, and upon the reactivity of the 

contaminants with the etching gas. This was illustrated by etching Be and Si δ-doped 

structures in GaAs. 

T.H.Chiu, W.T.Tsang, M.D.Williams, C.A.C.Mendonça, K.Dreyer, F.G.Storz: Journal of 


[446-127/128-120] 


The atomic structures of Ga and As atoms on (110) planes were studied by using a firstprinciples 

pseudopotential method. It was found that both Ga and As atoms resided in the 

center of a triangle that consisted of a surface Ga atom and 2 surface As atoms in the 

single-atom chemisorbed state. The adsorption energies for Ga and As atoms were 3.1 

and 3.5eV, respectively. The energy barrier heights for Ga and As atoms which migrated 

along the path through the interstitial channel were found to be 0.6 and 1.0eV, 

respectively. Simulation of the deposition of 2 atoms revealed that pair formation was 

stable with respect to separate single-atom chemisorption. 

J.Y.Yi, J.Y.Koo, S.Lee, J.S.Ha, E.Lee: Physical Review B, 1995, 52[15], 10733-6 

[446-125/126-122] 


A 1/6 monolayer of GaAs was grown onto a very flat GaAs surface by using metalorganic 

vapor-phase epitaxial techniques, and 2-dimensional nuclei were studied by using highvacuum 

scanning tunnelling microscopy. On the basis of the 2-dimensional nuclei densities, 

the surface diffusion coefficient at 530C was estimated to be equal to 2 x 10 -6 cm 2 /s. It was 

found that, during growth, the bunched-step (multi-step) separation saturated and was 

independent of the substrate misorientation angle. The results could be explained in terms of 

a mechanism that took account of 2-dimensional nucleus formation on the wider terraces, 

and their coalescence on ascending steps. A step-bunching simulation which was based 

upon this model revealed that the saturated multi-step separation was proportional to the 2- 

dimensional nucleus separation (that is, to the reciprocal of the square root of the density). 

M.Kasu, N.Kobayashi: Journal of Applied Physics, 1995, 78[5], 3026-35 

[446-123/124-162] 

386



The mechanisms of molecular beam epitaxy were investigated by growing and analyzing 

the shapes of facet structures which consisted of an (001) top surface and two (111)B side 

surfaces. It was found that all of the Ga flux on the 3 facet planes was incorporated into 

the film, but the growth rates on (111)B and (001) depended strongly upon the As flux 

and were governed mainly by the diffusion of Ga adatoms between the 2 planes. By 

analyzing the shape of the facet, the diffusion length of Ga on a (001) surface was 

estimated to be about 1µ at 580C. On (111)B, the diffusion length of Ga was found to be 

equal to several µ. The reflectivity of diffusing Ga atoms was found to be far less than 

unity for the (001)/(111)B boundary, and was almost equal to unity at facet boundaries 

where the (111)B side surfaces were bounded by the (1¯10) side walls. 

S.Koshiba, Y.Nakamura, M.Tsuchiya, H.Noge, H.Kano, Y.Nagamune, T.Noda, 

H.Sakaki: Journal of Applied Physics, 1994, 76[7], 4138-44 

[446-117/118-159] 


It was shown that Te and Pb which segregated at the surface during epitaxial growth 

decreased and increased, respectively, the surface diffusion length. This indicated that, 

under the generic term of surfactant, there were 2 types of surface-segregating species 

which had opposite effects upon surface diffusion. It was suggested that the key 

parameter which governed the surfactant-induced modification of epitaxial growth 

kinetics was the reactivity of a given pair of surfactant and growing materials. According 

to this theory, it was predicted that surfactants which occupied interstitial surface sites 

(non-reactive surfactants) increased the surface diffusion length, whereas surfactants 

which were in substitutional sites (reactive surfactants) decreased it. 

J.Massies, N.Grandjean: Physical Review B, 1993, 48[11], 8502-5 

[446-106/107-039] 


The growth rates of layers which were grown on a mesa-etched (001) surface were 

measured by using in situ scanning microprobe reflection high-energy electron diffraction 

methods. The diffusion lengths of the surface adatoms of column-III elements were 

deduced from the gradient of the variation of the growth rate. The diffusion lengths were 

of the order of one micron for every source/material combination. When metalorganic 

materials were used as a Ga source, it was found that the diffusion length was larger than 

that of Ga atoms from a metallic Ga source. Because the substrate temperatures which 

were used in the present experiments were high enough to decompose trimethylgallium 

and triethylgallium on the surface, Ga adatoms were considered to be responsible for the 

surface diffusion. It was concluded that derivatives of the metalorganic molecules, such 

as methyl radicals, affected the diffusion of Ga adatoms. 

T.Isu, M.Hata, Y.Morishita, Y.Nomura, S.Goto, Y.Katayama: Journal of Crystal Growth, 

1992, 120[1-4], 45-9 

[446-106/107-039] 

387

Interdiffusion GaAs Interdiffusion 


GaAs: Interdiffusion 

Diffusion and interdiffusion in GaAs was shown to be consistent with a charged point 

defect model. Charged Ga vacancies, V Ga 3- , and interstitials, I Ga 2+ , appeared to control the 

diffusion of group-II, group-III and (probably) group-V elements. After adjusting for 

carrier concentration and As pressure, these elements were found to have an almost 

identical intrinsic diffusivity and activation energy over a wide range of temperatures. A 

natural consequence, of Ga diffusion via negative or positive point defects, was that 

enhanced group-III interdiffusion was expected under either n-type or p-type doping. 

Anomalous enhancements of group-II dopant diffusivity were related to the 

supersaturation of Ga interstitials. 

R.M.Cohen: Journal of Applied Physics, 1990, 67[12], 7268-73 

[446-78/79-011] 

GaAs: Interdiffusion 

The interdiffusion between silicides (W, Ta, Mo, Ti) and GaAs, at temperatures of 825 to 

975C, was studied by using Rutherford back-scattering spectrometry, particle-induced X- 

ray emission, and X-ray diffraction techniques. The specimens were furnace-annealed at 

825 or 875C, or rapid thermally annealed at 975C. Almost no interdiffusion was observed 

in the cases of WSi 2 , TaSi 2 and TiSi 2 at 825C. It was not observed at all at 975C. In the 

case of MoSi 2 , there was marked atomic migration at temperatures as low as 825C. It was 

concluded that rapid thermal annealing was more beneficial from the point of view of 

interface stability. 

J.Osvald, R.Sandrik: Thin Solid Films, 1989, 169[2], 223-8 

[446-64/65-163] 

GaAs/Al/GaAs: Interdiffusion 

It was shown that the presence of low-temperature-grown GaAs, in GaAs/AlAs on Sidoped 

GaAs heterostructures, increased Al/Ga interdiffusion at the heterostructure 

interfaces. The interdiffusivity enhancement was attributed to the presence of Ga 

vacancies, V Ga , in the As-rich low-temperature GaAs, which diffused from a 

supersaturation of V Ga which was frozen in during growth. Chemical mapping, which 

distinguished between the AlAs and GaAs lattices at the atomic scale, was used to 

measure the Al concentration gradient in adjacent Si-doped GaAs layers. 

C.Kisielowski, A.R.Calawa, Z.Liliental-Weber: Journal of Applied Physics, 1996, 80[1], 

156-60 

[446-136/137-115] 

GaAs/AlAs: Interdiffusion 

The results of ab initio molecular dynamics simulations of impurity-induced disordering 

in superlattices were presented. Two typical impurities, the Zn acceptor and the Si donor, 

388


were studied. It was found that Zn-induced interdiffusion was due to the formation of 

non-equilibrium group-III interstitials during Zn in-diffusion. The interstitials, which 

were bound to Zn acceptors via Coulomb forces, disordered the superlattice via kick-out 

processes on the group-III sub-lattice. On the other hand, Si-induced interdiffusion 

occurred via Si Ga -V Ga pairs. Their motion, via second-nearest neighbor jumps, disordered 

the group-III sub-lattice. The calculated activation energies for these processes were in 

good agreement with experiment. 

J.Bernholc, B.Chen, Q.Zhang, C.Wang, B.Yakobson: Materials Science Forum, 1994, 

143-147, 593-8 

[446-113/114-028] 


Photoluminescence measurements were used to investigate C impurity effects upon the 

intermixing behavior of multiple quantum wells which had been grown by means of 

molecular beam epitaxy. The wells were furnace annealed, with a C source. The 

photoluminescence spectra revealed that the degree of intermixing of Al and Ga, which 

was induced by thermal annealing, increased with depth. This behavior did not agree with 

intermixing mechanisms which considered vacancy injection of the surface. The nonuniformity 

of intermixing as a function of depth was attributed to the effect of C 

impurities which were injected during heat treatment. 

Y.T.Oh, S.K.Kim, Y.H.Kim, T.W.Kang, C.Y.Hong, T.W.Kim: Journal of Applied 

Physics, 1995, 77[6], 2415-8 

[446-121/122-062] 


The relative stability of ideal-geometry versus intermixed buried semiconductor interfaces 

was studied by using first-principles density functional methods. It was found that, 

although intermixing of an ideal lattice-matched GaAs/AlAs(001) interface required 

energy, intermixing was an energy-lowering process at the coherent strained-layer (001) 

interfaces of lattice-mismatched materials. Intermixing of an ideal strained-layer (001) 

interface lowered the energy mainly because it partially relieved the strain in a local 

region near to the interface. It was predicted that lattice-mismatched interfaces would 

have a greater degree of atomic-scale micro-roughness than would the analogous 

interfaces of lattice-matched materials. 

R.G.Dandrea, C.B.Duke: Physical Review B, 1992, 45[24), 14065-8 

[446-93/94-027] 


Annealing experiments were performed on undoped single-well heterostructures which 

had been grown by means of molecular beam epitaxy. The annealing was carried out in 

evacuated and sealed silica ampoules. The annealing temperature and time were 1000C 

and 4h, respectively, while the As vapor pressure in the ampoules was varied from the 

dissociation pressure to about 108kPa. Compositional profiles were obtained by using 

dynamic secondary ion mass spectrometry. It was found that the amount of intermixing in 

389


the layers depended upon both the As pressure and the distance from the sample surface. 

In contrast with previous studies, the complex variation in interdiffusion as a function of 

As pressure which was observed here could not be explained in terms of interdiffusion 

via group-III vacancies and interstitials alone. 

N.Baba-Ali, I.Harrison, B.Tuck: Journal of Materials Science - Materials in Electronics, 

1995, 6[3], 127-34 

[446-127/128-138] 


Combined Raman and lattice dynamics studies of ultra-thin (GaAs) 4 (AlAs) 4 superlattices 

showed that a marked degree of interdiffusion occurred in these samples when they were 

grown by using conventional molecular beam epitaxial temperatures of between 580 and 

640C. At these temperatures, an interruption of growth had little effect upon the structural 

quality of the superlattices. 

J.Grant, J.Menéndez, L.N.Pfeiffer, K.W.West, E.Molinari, S.Baroni: Applied Physics - 

Letters, 1991, 59[22], 2859-61 

[446-88/89-030] 


A new method was proposed for the estimation of interdiffusion coefficients in 

superlattices. This was based upon measurements of the thickness of layers which 

remained, without forming an alloy, after an annealing process which caused 

interdiffusion. The measurements were based upon the frequency of phonons from the 

Raman spectra. The interdiffusion coefficient values which were found in this way were 

almost the same as those previously published. It was noted that Ga atoms in the present 

superlattices diffused more rapidly into AlAs layers than Al atoms diffused into GaAs 

layers. The interdiffusion coefficients first decreased with annealing time and increased 

slightly when annealing was performed for more than 1.5h at 860C. 

N.Hara, T.Katoda: Journal of Applied Physics, 1991, 69[4], 2112-6 

[446-78/79-030] 

GaAs/AlAs/AlGaAs: Interdiffusion 

The intermixing of double-barrier quantum wells by 50keV Ga + implantation was studied 

experimentally and theoretically. It was found that, even at low doses (less than 

10 12 /cm 2 ), a considerably broadened emission peak with an appreciable luminescence 

blue-shift could be obtained. At medium doses (about 10 13 /cm 2 ), very large blue-shifts of 

the order of 0.2eV were observed which retained a reasonable emission intensity. At high 

doses (above 3 x 10 14 /cm 2 ), total intermixing occurred and no photoluminescence could 

be recovered. The photoluminescence blue-shifts after low-dose implantation were not 

affected by the annealing temperature, whereas the blue-shift at high doses depended 

greatly upon the annealing conditions. The data indicated heterogeneously enhanced 

390


interdiffusion that was based upon a defect cluster model. It was noted that close control 

of the annealing ambient and sample surface was important. 

R.K.Kupka, Y.Chen: Journal of Applied Physics, 1995, 78[4], 2355-61 

[446-123/124-172] 

GaAs/AlGaAs: Interdiffusion 

The occurrence of interdiffusion during annealing was characterized by Raman scattering 

studies of a GaAs/Al x Ga 1-x As superlattice into which various degrees of damage had been 

introduced by the implantation of electrically inactive iso-electronic 31 P + . This process 

eliminated impurity charge associated effects. Auger and secondary ion mass 

spectrometry methods were used to determine the amount of mixing beyond the damage 

zone in the superlattice. As a result, it was possible to distinguish between the 

contributions of impurities and implantation-induced defects to Ga/Al intermixing. 

J.Sapriel, E.V.K.Rao, F.Brillouet, J.Chavignon, P.Ossart, Y.Gao, P.Krauz.: Superlattices 

and Microstructures, 1988, 4[1], 115-20 

[446-61-076] 


The intermixing of hetero-interfaces, by Ga + implantation and annealing, was 

investigated. Damage accumulation in a GaAs/AlAs superlattice was found to be less 

rapid than that in a GaAs/GaAlAs quantum-well structure. Low-temperature 

photoluminescence spectroscopy of a GaAs/AlAs superlattice was performed for doses as 

high as 10 16 /cm 2 . The photoluminescence spectra exhibited several emission bands on the 

high-energy side. The number and energy of these blue-shifted peaks were found to 

depend upon the implanted dose. Secondary ion mass spectrometric data suggested that 

they could be interpreted as being due to emission from several quantum wells, of the 

superlattice, which disordered at differing mixing rates. Two regimes were revealed. 

Thus, while the depth extension of the disordering was directly related to the postimplantation 

defect distribution in the high-dose regime, some diffusion of these defects 

during annealing occurred in the low-dose regime. Cross-sectional transmission electron 

microscopy confirmed that there was an effect of the structure, of the implanted sample, 

upon damage accumulation. A decrease in the photoluminescence intensity after 

annealing was attributed to the presence of extended residual defects in the implanted 

layers. A study of the effect of annealing time at 760C showed that the 

photoluminescence intensity progressively recovered, whereas the intermixing rapidly 

saturated. 

C.Vieu, M.Schneider, R.Planel, H.Launois, B.Descouts, Y.Gao: Journal of Applied 

Physics, 1991, 70[3], 1433-43 

[446-93/94-028] 


The shape of the confinement potential which resulted from interdiffusion of a GaAs 

quantum well, locally enhanced by defects due to Ga implantation, was computed. The 

simplest model which could take account of lateral diffusion of the defects was used. A 

391


variational calculation of the first 2 electronic levels within this two-dimensional potential 

supported the assignment of recently observed new cathodoluminescence lines to 

electrons which were laterally confined in a graded potential. 

J.Cibert, P.M.Petroff: Physical Review B, 1987, 36[6], 3243-6 

[446-55/56-020] 


Low-energy As + implantation, followed by rapid thermal annealing, was used to modify 

the exciton transition energies of quantum wells. Various structures were irradiated using 

an energy which was sufficiently low that the disordered region was spatially separated 

from the quantum wells. After rapid thermal annealing, the exciton energies exhibited 

large increases which depended upon the quantum-well width and the implantation 

fluence. There was no appreciable increase in the peak line-widths. The observed energy 

shifts were attributed to modifications of the shapes of the quantum wells, due to 

enhanced Ga and Al interdiffusion at hetero-interfaces in irradiated areas. The results 

were consistent with a model which was based upon an enhanced intermixing of Al and 

Ga atoms, in the depth of the material, due to the diffusion of vacancies which were 

generated near to the surface. 

B.Elman, E.S.Koteles, P.Melman, C.A.Armiento: Journal of Applied Physics, 1989, 

66[5], 2104-7 

[446-74-027] 


Implantation with Se + ions, and the annealing of quantum-well samples, were studied by 

using transmission electron microscopy, photoluminescence spectroscopy and Monte 

Carlo simulation. It was concluded that enhanced layer interdiffusion occurred at depths 

which were several times greater than the projected range of the implanted Se. There 

were signs of residual stress at similar depths. 

E.G.Bithell, W.M.Stobbs, C.Phillips, R.Eccleston, R.Gwilliam: Journal of Applied 

Physics, 1990, 67[3], 1279-87 

[446-74-027] 


The transient-enhanced interdiffusion of interfaces during the rapid thermal annealing of 

ion-implanted heterostructures was studied. It was shown that the factors which most 

influenced the degree of interdiffusion were the temperature, the concentration of excess 

vacancies, and the ability of the vacancies to diffuse freely. A model was developed in 

order to explain these observations. It was based upon the solution of coupled diffusion 

equations which involved excess vacancies and the distribution of Al after ion 

implantation. Both initial distributions were deduced from a 3-dimensional Monte Carlo 

simulation of ion implantation into a heterostructure sample. The model was found to give 

392


excellent agreement with experiment. In particular, it was valid for as-implanted vacancy 

concentrations of less than 6 x 10 19 /cm 3 . 

K.B.Kahen, G.Rajeswaran: Journal of Applied Physics, 1989, 66[2], 545-51 

[446-74-026] 


A high-energy (up to 150keV) Ga + focussed ion beam was used to implant quantum-well 

structures, and to interdiffuse heterojunctions so as to create quantum wires and boxes. 

Wires as wide as 80nm were found 200nm below the surface. Optical damage and 

interdiffusion processes were studied as a function of the implantation parameters and the 

(rapid) thermal annealing time. A universal correlation was found between the optical 

damage and the interdiffusion length. 

F.Laruelle, P.Hu, R.Simes, R.Kubena, W.Robinson, J.Merz, P.M.Petroff: Journal of 

Vacuum Science and Technology B, 1989, 7[6], 2034-8 

[446-74-026] 


The occurrence of enhanced interdiffusion at the interface was studied. Ions of Ar were 

implanted to doses which ranged from 2 x 10 13 to 5 x 10 14 /cm 2 , and the samples were 

then rapidly thermally annealed (950C, 30s). It was found that the degree of intermixing 

decreased from the surface towards the projected ion range, and was a function of the 

implantation dose. It was suggested that this variation arose from the coalescence of some 

of the excess vacancies into extended defects, so that they were then unavailable to the 

enhanced diffusion mechanism. By assuming that the concentration of mobile vacancies 

at a given depth was proportional to the electronic energy of the ion, and inversely 

proportional to the nuclear energy loss of the ion, predictions were obtained which were 

in good agreement with the experimental results. 

K.B.Kahen, D.L.Peterson, G.Rajeswaran: Journal of Applied Physics, 1990, 68[5], 2087- 

90 

[446-86/87-031] 


Diffusion and interdiffusion in GaAs/AlGaAs superlattices was shown to be consistent 

with a charged point defect model. Charged Ga vacancies, V Ga 3- , and interstitials, I Ga 2+ , 

appeared to control the diffusion of group-II, group-III and (probably) group-V elements. 

After adjusting for carrier concentration and As pressure, these elements were found to 

have an almost identical intrinsic diffusivity and activation energy over a wide range of 

temperatures. A natural consequence, of Ga diffusion via negative or positive point 

defects, was that enhanced group-III interdiffusion was expected under either n-type or p- 

type doping. Anomalous enhancements of group-II dopant diffusivity were related to the 

supersaturation of Ga interstitials. 


[446-78/79-011] 

393



The superlattice and multi-quantum well structures of various III-V semiconductors were 

remarkably stable to thermal annealing. This was very different to the instability of these 

structures which occurred when an impurity such as Zn or Si was diffused into them. It 

was also recalled that selective area impurity-induced disordering of multi-quantum well 

structures had become a potentially powerful technique for the fabrication of certain 

devices. The current state of understanding of the fundamental mechanisms was reviewed 

here. It was pointed out that, in n-type material, the involvement of group-III vacancies 

was generally accepted. However, in the case of p-type dopants such as Zn, an oversaturation 

of self-interstitials was required for the enhancement of interdiffusion. 

I.Harrison, H.P.Ho, N.Baba-Ali: Journal of Electronic Materials, 1991, 20[6], 449-56 

[446-91/92-016] 


The compositional disordering of GaAs/AlGaAs quantum wells, due to the presence of 

low-temperature molecular beam epitaxially grown GaAs, was studied. It was found that 

Ga vacancy-enhanced interdiffusion was the mechanism which was responsible for the 

observed intermixing. The diffusion equations were solved numerically in order to obtain 

the band profile after intermixing. The transition energies in the quantum wells under 

various annealing conditions were predicted, and were found to agree very well with 

observed photoluminescence emission peaks. The vacancy-induced interdiffusivity was 

found to require an activation energy of 4.08eV. This was smaller than the activation 

energy for the interdiffusion of GaAs/AlGaAs heterostructures which were grown at 

normal temperatures. It was concluded that the present results clearly indicated an 

enhanced interdiffusion that was due to the presence of GaAs which had been grown at 

low temperatures. 

J.S.Tsang, C.P.Lee, S.H.Lee, K.L.Tsai, H.R.Chen: Journal of Applied Physics, 1995, 

77[9], 4302-6 

[446-121/122-062] 


The implantation of Ga + , followed by rapid thermal annealing, was used to enhance 

interdiffusion in single quantum wells. The extent of intermixing was found to depend 

upon the well depth, the number of implanted ions and the annealing time. Very rapid 

interdiffusion occurred in the initial annealing stage. The enhanced diffusion coefficient 

subsequently returned to the non-implanted value. A 2-step model was proposed in order 

to explain the diffusion process as a function of annealing time. This involved a rapid 

diffusion process and a saturated diffusion process. The interdiffusion coefficient for the 

rapid diffusion was found to be well depth-dependent and was estimated to be between 

5.4 x 10 -16 and 1. 5 x 10 -15 cm 2 /s. 

N.Sai, B.Zheng, J.Xu, P.Zhang, X.Yang, Z.Xu: Solid State Communications, 1996, 

98[12], 1039-42 

[446-136/137-116] 

394



A new impurity-free interdiffusion technique was described which involved pulsed 

anodization followed by rapid thermal annealing at temperatures near to 900C. Enhanced 

interdiffusion was observed, in the presence of an anodized GaAs capping layer, in 

GaAs/AlGaAs quantum-well structures. The use of transmission electron microscopy 

revealed evidence of interdiffusion. The photoluminescence spectra from interdiffused 

samples exhibited a large blue-shift, with no appreciable line-width broadening. 

S.Yuan, Y.Kim, C.Jagadish, P.T.Burke, M.Gal, J.Zou, D.Q.Cai, D.J.H.Cockayne, 

R.M.Cohen: Applied Physics Letters, 1997, 70[10], 1269-71 

[446-148/149-178] 


A formula was derived which described the interdiffusion profiles of quantum wells. It 

was shown that it accurately modelled interdiffusion in quantum wells of lattice-matched 

AlGaAs. The formula took account of the differing interdiffusion coefficients between 

layers, and of the interfacial discontinuity of interdiffused species. The formula explained 

how quantum energy shifts due to interdiffusion varied with annealing time and annealing 

temperature in various wide-well layers of both InGaAsP/InP and GaAs/AlGaAs quantum 

wells. The quantitative difference between the interdiffusion profiles in these two 

materials was also demonstrated. 

K.Mukai, M.Sugawara, S.Yamazaki: Physical Review B, 1994, 50[4], 2273-82 

[446-115/116-130] 


The dependence of impurity-free interdiffusion upon the properties of a dielectric cap 

layer was studied in unstrained multi-quantum well structures that had been grown by 

means of molecular beam epitaxy. Electron-beam evaporated SiO 2 films, chemical vapor 

deposited SiO x N y films, and spun-on SiO 2 films were used as cap layers during rapid 

thermal annealing at temperatures of between 850 and 950C. The photoluminescence at 

10K was used to monitor interdiffusion-induced band-gap shifts, and to calculate the 

corresponding Al-Ga interdiffusion coefficients. The latter were found to increase with 

cap layer thickness (electron-beam SiO 2 ) up to a limit which was governed by saturation 

of the out-diffused Ga concentration in the SiO 2 caps. A maximum concentration of 

between 4 x 10 19 and 7 x 10 19 /cm 3 in the SiO 2 caps was found by using secondary ion 

mass spectroscopic profiling. Larger band-edge shifts were also obtained when the O 

content of SiO x N y cap layers was increased, but the differences were insufficient to 

suggest a laterally selective interdiffusion process that was based upon variations in cap 

layer composition alone. Much larger differences were obtained by using various 

deposition techniques for the cap layers. This indicated that the porosity of the cap layer 

395


was a much more important factor than was the film composition in obtaining a laterally 

selective interdiffusion process. 

S.Bürkner, M.Maier, E.C.Larkins, W.Rothemund, E.P.O’Reilly, J.D.Ralston: Journal of 


[446-125/126-131] 

GaAs/AlInP: Interdiffusion 

An analysis was made of the interdiffusion of a discrete GaAs layer, into an Al 0.5 In 0.5 P 

half-space, by using Si doping as an agent for enhanced layer interdiffusion. Enhanced 

interdiffusion was observed on both column-III and column-V sites; but the column-III 

interdiffusion coefficient exceeded the column-V interdiffusion coefficient by 2 orders of 

magnitude. Due to this disparity between the diffusion coefficients, large defectproducing 

strains were introduced by the interdiffusion. It was shown that, by modelling 

the resultant strain profiles and by applying a critical thickness analysis, the instability of 

such interdiffused structures could be understood. 

R.L.Thornton, F.A.Ponce, G.B.Anderson, F.J.Endicott: Applied Physics Letters, 1993, 

62[17], 2060-2 

[446-106/107-082] 

GaAs/GaAlAs: Interdiffusion 

Simple analytical expressions were derived for the approximate estimation of the 

interdiffusion coefficient, of partially disordered quantum-well heterostructures, directly 

from measurements of the photoluminescence phase shift which was associated with 

layer interdiffusion. The phase shift was calculated as a function of the interdiffusion 

length, (Dt) ½ , in the lattice-matched system, GaAs/Ga 0.7 Al 0.3 As. The calculations were 

performed within the framework of the envelope function approximation and Fick's law. 

A simple relationship was derived for the variation in phase shift as a function of the 

dimensionless parameter, (Dt) ½ /L, where L was the quantum-well thickness. This 

satisfactorily accounted for most of the published interdiffusivity values, to within a 

factor of 2. 

M.T.Furtado, M.S.S.Loural: Superlattices and Microstructures, 1993, 14[1], 21-5 

[446-113/114-029] 

332 GaAs/GaAlAs: Interdiffusion 

The effects of Si and Be, at doping levels of up to 10 19 /cm 3 , upon the interdiffusion of 

quantum wells after annealing were studied by using photoluminescence techniques (table 

32). It was found that, for Be concentrations of up to 2.5 x 10 19 /cm 3 , and for Si 

concentrations of up to 10 18 /cm 3 , no change in the interdiffusion coefficients could be 

measured. At a Si dopant concentration of 6 x 10 18 /cm 3 , there was a dramatic degradation 

of the material quality after annealing (750C, 15s). This caused the luminescence from 

the well to disappear, while a deep-level luminescence that was related to donor-Ga 

vacancy complexes and As antisite defects appeared. On the basis of these results, it was 

suggested that the position of the Fermi level played no role in the intermixing of III-V 

396


heterostructures. It was also concluded that most of the enhanced intermixing which was 

observed in Si-doped GaAs/AlGaAs structures was related to Si relocation at very high 

doping levels. 

W.P.Gillin, I.V.Bradley, L.K.Howard, R.Gwilliam, K.P.Homewood: Journal of Applied 

Physics, 1993, 73[11], 7715-9 

[446-106/107-082] 

Table 32 

Interdiffusion Data for Ga 0.8 Al 0.2 As/GaAs 

Dopant Amount (/cm 3 ) Temperature (C) Coefficient (cm 2 /s) 

- - 1000 1.54 x 10 -16 

- - 1050 6.95 x 10 -16 

- - 1100 3.76 x 10 -15 

Si 5 x 10 17 1000 4.50 x 10 -16 

Si 5 x 10 17 1050 5.60 x 10 -16 

Si 5 x 10 17 1100 1.19 x 10 -15 

Si 10 18 1000 8.60 x 10 -17 

Si 10 18 1050 4.50 x 10 -16 

Si 10 18 1100 1.79 x 10 -15 


The Al-Ga interdiffusion which was produced by focussed Si ion-implantation and rapid 

thermal annealing was investigated in a Ga 0.7 Al 0.3 As/GaAs superlattice structure with 

equal (3.5nm) barrier and well widths. Ions of Si 2+ were accelerated to 50 or 100kV and 

were implanted, parallel to the sample normal, to doses which ranged from 10 13 to 

10 15 /cm 2 . The effect of rapid thermal annealing (950C, 10s) was characterized by means of 

secondary ion mass spectrometry. It was found that, in the implanted region, the 

interdiffusion was significantly enhanced by Si implantation. Ion doses which were as low 

as 10 14 /cm 2 led to a 2 orders of magnitude increase, in the interdiffusion coefficient, to a 

value of 4.5 x 10 -14 cm 2 /s. This led to a mixing effectiveness of about 90%. On the other 

hand, the use of rapid thermal annealing alone produced an interdiffusion coefficient of 1.3 

x 10 -16 cm 2 /s; with very little mixing. A marked depth dependence of the mixing process 

was observed at an implantation energy of 100keV, with a more heavily mixed, so-called 

pinch-off, region being formed at a certain depth. It was noted that the depth at which this 

enhancement occurred was not associated with a maximum concentration of Si ions or of 

vacancies. It instead coincided with a positive maximum in the second derivative of the 

vacancy profile. This, in turn, represented a maximum in the vacancy 

397


injection process that was caused by the presence of a transient vacancy concentration 

gradient. 

P.Chen, A.J.Steckl: Journal of Applied Physics, 1995, 77[11], 5616-24 

[446-121/122-065398] 


Electrically inactive iso-electronic 31 P + and 27 Al + were implanted into molecular beam 

epitaxially grown GaAs/Ga 0.7 Al 0.3 As superlattices at 25 or 250C. Evidence was found for 

implantation damage and for Al/Ga interdiffusion which depended upon the annealing 

time. 

E.V.K.Rao, F.Brillouet, P.Ossart, Y.Gao, J.Sapriel, P.Krauz: Journal de Physique - 

Colloque C5, 1987, 48[11], 113-6 

[446-61-076] 


By carrying out various annealing treatments on Sn-doped molecular beam epitaxially 

grown GaAs-Ga 0.72 Al 0.28 As quantum well structures it was shown that Sn, like other 

donor atoms (Si, S), induced disordering by enhancing interdiffusion. Also, the voluntary 

introduction of B atoms into a Sn-doped structure before annealing led to a retardation of 

Sn-enhanced interdiffusion. 

E.V.K.Rao, P.Ossart, F.Alexandre, H.Thibierge: Applied Physics Letters, 1987, 50[10], 

588-91 

[446-51/52-125] 

GaAs/GaAsP, GaAs/GaAsSb: Interdiffusion 

Interdiffusion in superlattices was studied at various temperatures and under various As 

partial pressures. An analysis of the As pressure-dependence of the effective diffusion 

coefficient revealed that a substitutional-interstitial diffusion mechanism governed the 

interdiffusion process. Computer simulations were used to study the profile shapes of 

annealed samples, and the As pressure dependence of the effective diffusion coefficient. 

It was found that the Frank-Turnbull diffusion mechanism governed the interdiffusion of 

these superlattices. The As pressure-dependence of the effective diffusion coefficients, as 

measured in interdiffusion experiments, was opposite to the reported pressure 

dependences which had been measured in As and P in-diffusion experiments. The 

apparently contradictory in-diffusion and out-diffusion behaviors could be reconciled by 

a diffusion model which involved As vacancies, fast-diffusing As-vacancy plus P- 

interstitial complexes, and fast-diffusing P interstitials (or the analogous Sb-related 

defects). 

M.Schultz, U.Egger, R.Scholz, O.Breitenstein, P.Werner, U.Gösele, R.Franzheld, 

M.Uematsu, H.Ito: Defect and Diffusion Forum, 1997, 143-147, 1101-8 

[446-143/147-1101] 

398


GaAs/GaInP: Interdiffusion 

An analytical and experimental investigation was made of the mechanisms of defect 

formation during the interdiffusion of these materials. It was found that the analytical 

model predicted a critical thickness, below which defects were not produced during this 

highly strained interdiffusion process. Transmission electron microscopic analysis of 

diffused buried layers of various thickness revealed a very good qualitative agreement 

with the present model. 

R.L.Thornton, D.P.Bour, D.Treat, F.A.Ponce, J.C.Tramontana, F.J.Endicott: Applied 


[446-119/120-203] 

GaAs/In: Interdiffusion 

The effects of interactions between a thick In layer and heat-treated GaAs at 570C were 

investigated by means of scanning electron microscopy, secondary ion mass 

spectrometry, Rutherford back-scattering spectrometry, X-ray diffraction and Nomarski 

techniques. It was shown that, as well as the usual InGaAs crystallites which grew 

epitaxially upon dissolution of the substrate, an array of In-rich dendrites was present 

whose numbers were related to the dislocation density. The driving force for In to migrate 

along the dislocations and eventually form InGaAs spikes was assumed to be an excess of 

As which had been reported to be present in the vicinity of individual dislocations. It was 

suggested that existing data on the coefficient for the conventional diffusion of In in 

GaAs had been overestimated by a factor of 10 6 . 

A.J.Barcz, J.M.Baranowski, S.Kwiatkowski: Applied Physics A, 1995, 60[3], 321-4 

[446-134/135-136] 

GaAs/InGaAs: Interdiffusion 

The interdiffusion of a multiple quantum-well sample, due to a thin source of vacancies, 

was used as a probe for the simultaneous measurement of the interdiffusion coefficient, 

the diffusivity of group-III vacancies and the background concentration of these 

vacancies. It was shown that interdiffusion at all temperatures was governed by a constant 

background concentration of vacancies, and that this concentration was equal to that of 

the vacancies in the substrate. The measured vacancy concentration was about 2 x 

10 17 /cm 3 . This showed that the vacancy concentrations were not in thermal equilibrium, 

contrary to the usual assumption. It instead had a value which was frozen in; probably at 

the growth temperature. It was shown that the activation energy for the intermixing of 

InGaAs and GaAs was governed only by the activation energy for vacancy diffusion, 

which was estimated to have a value of 3.4eV. 

O.M.Khreis, W.P.Gillin, K.P.Homewood: Physical Review B, 1997, 55[23], 15813-8 

[446-152-0399] 

GaAs/Ir: Interdiffusion 

Interfacial reactions, in both thin-film and bulk samples, were studied at temperatures of 

between 400 and 1000C. The diffusion path for Ir/GaAs was found to be: 

399

Interdiffusion GaAs|Ga(As,P) Zn 

Ir/lrGa/IrAs 2 /GaAs. In the thin-film case, where the Ir supply was limited, the final 

configuration was: Ga 5 Ir 3 /IrAs 2 /GaAs. The activation energies for diffusion in the thinfilm 

and bulk cases were 3.15 and 2.96eV, respectively. 

K.J.Schulz, O.A.Musbah, Y.A.Chang: Journal of Applied Physics, 1990, 67[11], 6798- 

806 

[446-78/79-032] 

Ga(As,P) 

Zn 

GaAsP: Zn Diffusion 


and into GaAs 0.6 P 0.4 , was described. The oxide films were deposited by using metalorganic 

chemical vapor deposition. A capping layer of SiO 2 was deposited on top of the 

source films, and diffusion was carried out in flowing N at 650C. Diffusion depths of 

between 200nm and several microns could be easily obtained. The diffusion front in n- 

type substrates was sharp. The dependence of the diffusion depth upon the source film 



[446-78/79-002] 


Samples of GaAs 0.6 P 0.4 were diffused with Zn, via a 200 to 300nm protective ZrO 2 layer. 

The diffusion depth exhibited a square-root time dependence. The absolute diffusivity 

values depended slightly upon the diffusion conditions. The layer had essentially no 

effect upon the carrier concentration profile or the activation energy. 


[446-74-003] 

400

Zn Ga(As,P)|Ga(As,Sb) Interdiffusion 


The open-tube diffusion of Zn into GaAs 0.8 P 0.2 , from a Zn-doped silica film, was 

investigated by using films of AlN or SiN x as annealing caps. The dependence of the 

diffusion depth upon the thickness of an AlN cap was found to differ from the 

dependence upon SiN x cap thickness. Selective masked diffusion of Zn, using an AlN 

mask, was also studied. It was found that the diffusion depth during selective masked 

diffusion depended upon both the AlN cap thickness and the AlN diffusion-mask 

thickness. The results suggested that the diffusion depth was not necessarily governed by 

either the cap thickness or the diffusion-mask thickness. It was concluded that the total 

film stress was the main factor which determined the diffusion depth under the present 

conditions. 

M.Ogihara, M.Taninaka, Y.Nakamura: Journal of Applied Physics, 1996, 79[6], 2995- 

3002 

[446-131/132-175] 


GaAsP/GaAs: Interdiffusion 

Interdiffusion on the group-V sub-lattice of GaAs was studied. Strained 

GaAs 0.86 P 0.14 /GaAs and GaAs 0.8 P 0.2 /GaAs 0.975 P 0.025 superlattices were annealed at 

temperatures of between 850 and 1100C, under various As vapor pressures. The diffusion 

coefficients were measured by means of secondary ion mass spectroscopy and 

cathodoluminescence spectroscopy. It was found that the interdiffusion coefficients were 

higher under As-rich conditions than under Ga-rich conditions; thus indicating an 

interstitial-substitutional diffusion mechanism. 

U.Egger, M.Schultz, P.Werner, O.Breitenstein, T.Y.Tan, U.Gösele, R.Franzheld, 

M.Uematsu, H.Ito: Journal of Applied Physics, 1997, 81[9], 6056-61 

[446-150/151-144] 

Ga(As,Sb) 

GaAsSb/GaAs: Interdiffusion 

Interdiffusion on the group-V sub-lattice of GaAs was studied. Strained 

GaAs 0.98 Sb 0.02 /GaAs superlattices were annealed at temperatures of between 850 and 

1100C, under various As vapor pressures. The diffusion coefficients were measured by 

means of secondary ion mass spectroscopy and cathodoluminescence spectroscopy. It was 

found that the interdiffusion coefficients were higher under As-rich conditions than under 

Ga-rich conditions; thus indicating an interstitial-substitutional diffusion mechanism. 

U.Egger, M.Schultz, P.Werner, O.Breitenstein, T.Y.Tan, U.Gösele, R.Franzheld, 

M.Uematsu, H.Ito: Journal of Applied Physics, 1997, 81[9], 6056-61 

[446-150/151-144] 

401

(Ga,In)As 

Al 

InGaAs/InAlAs: Al Diffusion 

It was shown that implantation of O (a basically non-dopant impurity), after adequate 

high-temperature furnace annealing (750C, 1h), led to significant interdiffusion of group- 

III atoms in molecular beam epitaxially grown In 0.53 Ga 0.47 As/In 0.52 Al 0.48 As multiquantum 

wells. Photoluminescence and Auger electron spectroscopic measurements 

(coupled with Ar + ion etching) were used to monitor the disordering of multi-quantum 

wells which were implanted with O (5 x 10 13 to 5 x 10 14 /cm 2 ) and then annealed by using 

either rapid thermal annealing or long-term furnace annealing. The role which was played 

by O in enhancing Al interdiffusion was unambiguously established, and a tentative 

explanation for this was based upon the possible migration of O in these multi-quantum 

wells. 

E.V.K.Rao, P.Ossart, H.Thibierge, M.Quillec, P.Krauz: Applied Physics Letters, 1990, 

57[21], 2190-2 

[446-78/79-043] 

As 

InGaAs: As Diffusion 

A theoretical model was proposed for the calculation of compositional variations in III-V 

ternary crystals during growth. In this novel model, phase equilibrium between the crystal 

and the growth solution was maintained; together with a simultaneous constancy of the 

transported and incorporated mass of solute atoms at the crystal/solution interface. This 

model could be applied to the calculation of diffusion-limited growth in a temperaturegradient 

solution; as in the case of the source current-controlled growth method. The 

compositional variations of InGaAs crystals which were grown via diffusion were 

calculated by using this model. The incorporation of As via the crystal/solution interface 

was considered on the basis of phase equilibrium laws and mass-constancy. Upon 

comparing the experimental results with the calculated ones it was found that, in the In-Ga- 

As solution, the diffusion coefficient of Ga was about twice as great as that of As. The 

402

As (Ga,In)As Be 

calculated compositional variations showed that the compositional uniformity of InGaAs 

crystals could be markedly improved by controlling growth parameters such as the 

temperature gradient in the solution. 

K.Nakajima: Journal of Crystal Growth, 1991, 110[4], 781-94 

[446-81/82-039] 

InGaAs/InGaAsP: As Diffusion 

A photoluminescence study was made of the interdiffusion of As in the 

In 0.66 Ga 0.33 As/In 0.66 Ga 0.33 As 0.7 P 0.3 system at temperatures ranging from 950 to 600C. It 

was shown that the diffusion was Fickian, with no dependence of the diffusion coefficient 

upon the substrate dopant-type or etch-pit density. In the case of Sn- or S-doped 

substrates, the diffusion could be described by: 

D(cm 2 /s) = 23 exp[-3.7(eV)/kT] 

at temperatures greater than 675C. This activation energy was the same as that deduced 

for group-III interdiffusion in Ga 0.8 In 0.2 As/GaAs. At lower temperatures, diffusivity in the 

present system could be described by: 

D(cm 2 /s) = 5 x 10 -10 exp[-1.7(eV)/kT] 

W.P.Gillin, S.S.Rao, I.V.Bradley, K.P.Homewood, A.D.Smith, A.T.R.Briggs: Applied 


[446-106/107-116] 

Be 

333 InGaAs: Be Diffusion 

The diffusion mechanisms which operated during post-growth annealing were 

investigated in p-type Be-doped epitaxial layers which had been grown between 2 

undoped InGaAs layers. Annealing (30 to 180s, 500 to 800C) was applied to samples 

with dopant concentrations of 5 x 10 18 , 10 19 or 3 x 10 19 /cm 3 . It was found that no 

appreciable Be diffusion occurred during post-growth rapid thermal annealing, for all of 

the dopant levels, when annealing was carried out at 500 or 600C for times of less than 

60s. The same was true, for a Be concentration of 5 x 10 18 /cm 3 , during annealing at 700C 

for 30s. At a doping level of 3 x 10 19 /cm 3 , significant Be diffusion occurred during 

annealing (>700C, 60s). The resultant curves exhibited a concave kink region. It was 

deduced that the effective Be diffusion coefficient (table 33) was approximately constant 

in one part of the concentration profile, and was proportional to the square root of the 

concentration in another part. By assuming the latter dependence, an effective diffusivity 

could be substituted into Fick’s second law. The resultant differential equation was 

solved numerically by using an explicit finite difference method. Good agreement was 

obtained between the measured depth profiles and the simulated distributions. 

S.Koumetz, J.Marcon, K.Ketata, M.Ketata, P.Launay: Journal of Physics D, 1997, 30[5], 

757-62 

[446-148/149-181] 

403

Be (Ga,In)As Be 

Table 33 

Diffusivity of Be in InGaAs 


800 1.7 x 10 -12 

700 3.4 x 10 -13 

InGaAs: Be Diffusion 

The diffusion of Be during post-growth annealing was studied in epitaxial layers. Two 

models were proposed in order to explain the observed concentration profiles. In the first 

model, the Boltzmann-Matano method was used while taking account of the diffusivity 

time-independence. An observed double profile was explained in terms of a change in 

diffusivity. In a second model, vacancy equilibrium was not assumed, and kinetic terms 

which were related to vacancy production had to be included in the diffusion equations. 

The observed double profile was then explained in terms of a reduction in the vacancy 

concentration in the crystal bulk. By using these two approaches, it was shown that good 

agreement between the experimental and simulated profiles could be obtained. In the case 

of the non-equilibrium model, the vacancy concentration was close to its equilibrium 

value. It was therefore possible to assume that Be diffusion in InGaAs epitaxial layers 

occurred with quasi-equilibrium point defect concentrations. Under these conditions, the 

diffusion depth varied as the square root of the diffusion time. There was no significant 

under-saturation of vacancies or super-saturation of self-interstitials at the diffusion front 

which could change the normalized Be diffusion depth in InGaAs epitaxial layers. 

J.Marcon, S.Gautier, S.Koumetz, K.Ketata, M.Ketata, P.Launay: Solid State 

Communications, 1997, 101[3], 159-62 

[446-141/142-110] 


The diffusion of Be from InGaAs epitaxial layers that had been grown between undoped 

InGaAs layers was investigated during post-growth annealing. A general substitutionalinterstitial 

diffusion mechanism was developed in order to explain the concentration 

profiles which were observed. The possibility of a concentration-dependent diffusivity 

was also considered in order to improve the fit to Be diffusion profiles. 

S.Koumetz, J.Marcon, K.Ketata, M.Ketata, F.Lefebvre, P.Martin, P.Launay: Materials 

Science and Engineering B, 1996, 37[1-3], 208-11 

[446-140-007] 


The behavior of implanted Be + ions was investigated during rapid thermal annealing at 

temperatures of between 600 and 900C. It was found that the apparent activation energy 

for Be was equal to 0.38eV. Higher activation efficiencies were found for the dopant in 

InGaAs, as compared with InAlAs. Anomalously low activation was detected for low- 

404

Be (Ga,In)As Be 

dose Be implants. The latter effect was attributed to a lack of vacant sites for the Be 

atoms to occupy. Extensive redistribution of the Be was observed after annealing (750C, 

10s). 


Physics, 1992, 71[1], 215-20 

[446-86/87-038] 


The occurrence of Be diffusion during post-growth annealing was studied in epitaxial 

InGaAs layers which were grown between 2 undoped InGaAs layers. In order to explain 

the observed concentration profiles and related diffusion mechanisms, a general 

substitutional-interstitial model was proposed. On one hand, simultaneous diffusion by 

dissociative and kick-out mechanisms was suggested and, on the other hand, the Fermilevel 

effect was used to explain changes in the effective diffusion coefficient of Be 

species as a function of concentration. A concentration-dependent diffusivity was also 

used to obtain an improved fit to Be diffusion profiles. 

S.Koumetz, J.Marcon, K.Ketata, M.Ketata, C.Dubon-Chevallier, P.Launay, 

J.L.Benchimol: Applied Physics Letters, 1995, 67[15], 2161-3 

[446-125/126-140] 

GaInAs: Be Diffusion 

The diffusion of Be from buried Be-doped layers was studied at temperatures of between 

600 and 700C. Four types of Be diffusion profile were identified. An interstitial cum 

substitutional model was proposed to be the diffusion mechanism, which depended upon 

the growth conditions. The values of the effective Be diffusion coefficient, for dopant 

levels which were above and below about 5 x 10 16 /cm 3 , were found to be 1.9 x 10 -15 and 

9 x 10 -16 cm 2 /s, respectively. 

E.G.Scott, D.Wake, G.D.T.Spiller, G.J.Davies: Journal of Applied Physics, 1989, 66[11], 

5344-8 

[446-74-030] 

InGaAs/GaAs: Be Diffusion 

The suppression, of Be out-diffusion from a Be-doped GaAs layer, by strained InGaAs 

layers was studied by using secondary ion mass spectroscopy. The experimental 

structures consisted of an 80nm Be-doped (about 10 19 /cm 3 ) GaAs layer that was 

sandwiched between 8nm In x Ga 1-x As layers, where x was equal to 0, 0.1, or 0.25. The 

samples were subjected to rapid thermal annealing (750C, 360s), and it was clearly 

observed that Be diffusion beyond the InGaAs layers was most rapid for a structure with 

x = 0, and was slowest for a structure with x = 0.25. 

K.Zhang, Y.C.Chen, J.Singh, P.Bhattacharya: Applied Physics Letters, 1994, 65[7], 872- 

4 

[446-119/120-214] 

405

Be (Ga,In)As Cd 

InGaAs/InP: Be Diffusion 

Heterojunctions which were grown by using liquid-phase epitaxy were investigated by 

using I-V and I o (T) measurements. It was shown that the injection of electrons across the 

hetero-interface was described well by a thermionic emission model which involved a 

barrier height that was strongly affected by p-type dopant diffusion. Heterojunction 

bipolar transistors with Schottky collectors were used as tools to measure the injection 

current in as-grown p + -In 0.53 Ga 0.47 As/n-InP diodes, and the results were compared with 

the predictions of the thermionic emission model. In this way, the relative diffusivities of 

the three p-type dopants in InP at 600C were evaluated. The results could be summarized 

by: 

D Be /D Mn ˜ 20, D Mg /D Be ˜ 1 

P.S.Whitney, J.C.Vicek, C.G.Fonstad: Journal of Applied Physics, 1987, 62[5], 1920-4 

[446-55/56-028] 

Cd 

InGaAs: Cd Diffusion 

The diffusion of Cd in In 0.53 Ga 0.47 As at 600C was studied by using a closed-ampoule 

technique, secondary ion mass spectrometry, and Hall effect measurements. It was found 

that the diffusivity could be described by an activation energy of 2.6eV. In some cases, a 

slower second diffusion front was found as well as a first steep junction. 

P.Ambreé, B.Gruska: Crystal Research and Technology, 1989, 24[3], 299-305 

[446-70/71-118] 


The Cd was diffused into InGaAs by using Cd 3 P 2 plus P or Cd 3 P 2 plus Cd 3 As 2 as 

diffusion sources. Two diffusion fronts were observed. The diffusion characteristics of 

Cd 3 P 2 plus P sources were explained in terms of the interstitial-substitutional model or 

the vacancy complex model. The charge state of the diffusing interstitial Cd atom was a 

singly ionized donor. Gaseous Cd originated from solid-phase CdP 2 . In the case of Cd 3 P 2 

plus Cd 3 As 2 diffusion sources, the effective diffusion coefficient and the surface acceptor 

concentration decreased with increasing weight fraction of Cd 3 As 2 . The relative depth of 

the deeper diffusion front increased when the supply of vacancies was suppressed. 

K.I.Ohtsuka, T.Matsui, H.Ogata: Japanese Journal of Applied Physics, 1988, 27[2], 253-9 

[446-60-009] 


A new source for Cd diffusion into In 0.53 Ga 0.47 As was developed in which Langmuir- 

Blodgett deposited monolayers of cadmium arachidate were used as the Cd source. 

D.M.Shah, W.K.Chan, R.Bhat, H.M.Cox, N.E.Schlotter, C.C.Chang: Applied Physics 

Letters, 1990, 56[21], 2132-4 

[446-76/77-023] 

406

Cd (Ga,In)As Ga 

InGaAs/InP: Cd Diffusion 

A closed-ampoule technique was used to study the simultaneous diffusion of Cd and Zn 

into material which was lattice-matched to InP. It was found that the Cd atoms, whose 

interstitial diffusion coefficient was small when compared with that of Zn interstitials, 

penetrated into the crystal as deeply as the Zn atoms. These data agreed with the results 

of numerical simulations of simultaneous diffusion, in III-V compounds, which were 

based upon an interstitial-substitutional diffusion mechanism. 

U.Wielsch, P.Ambrée, B.Gruska: Semiconductor Science and Technology, 1990, 5[9], 

923-7 

[446-76/77-024] 

InGaAs/InP: Cd Diffusion 

Results on diffusion across InGaAs/InP and InP/InGaAs hetero-interfaces were described. 

Diffusion from an InP top layer, Cd diffusion, or simple annealing of the samples, had no 

measurable effect upon the stability of the interfaces. The marked interdiffusion of In and 

Ga host atoms, as well as Zn gettering at the interface, were analyzed in terms of kick-out 

and vacancy mechanisms. 

P.Ambrée, A.Hangleiter, M.H.Pilkuhn, K.Wandel: Applied Physics Letters, 1990, 56[10], 

931-3 

[446-74-039] 

D 

InGaAs/AlGaAs: D Diffusion 

The effects of monatomic D diffusion in quantum wells were studied by using 

photoluminescence and secondary ion mass spectroscopic methods. The multiple 

quantum well structures were grown by means of molecular beam epitaxy and were 

hydrogenated with a remote plasma. A significant increase in the 77K photoluminescence 

integrated intensity of bound excitons was observed after hydrogenation. This was 

attributed to the passivation of non-radiative recombination centers within the quantum 

wells. The studies demonstrated that there was an increase in passivation efficiency with 

increasing Al concentration in the barriers, and that the hydrogenation was stable to 

temperatures above 450C. Overall, the results all strongly suggested that the passivated 

non-radiative recombination centers were interface defects. 

S.M.Lord, G.Roos, J.S.Harris, N.M.Johnson: Journal of Applied Physics, 1993, 73[2], 

740-8 

[446-106/107-113] 

Ga 

InGaAs: Ga Diffusion 

A theoretical model was proposed for the calculation of compositional variations in III-V 

ternary crystals during growth. In this novel model, phase equilibrium between the crystal 

407

Ga (Ga,In)As Ga 

and the growth solution was maintained; together with a simultaneous constancy of the 

transported and incorporated mass of solute atoms at the crystal/solution interface. This 

model could be applied to the calculation of diffusion-limited growth in a temperaturegradient 

solution; as in the case of the source current-controlled growth method. The 

compositional variations of InGaAs crystals which were grown via diffusion were 

calculated by using this model. The incorporation of Ga via the crystal/solution interface 

was considered on the basis of phase equilibrium laws and mass constancy. Upon 

comparing the experimental results with the calculated ones it was found that, in the In- 

Ga-As solution, the diffusion coefficient of Ga was about twice as great as that of As. 

The calculated compositional variations showed that the compositional uniformity of 

InGaAs crystals could be markedly improved by controlling growth parameters such as 

the temperature gradient in the solution. 


[446-81/82-039] 

GaInAs/GaAs: Ga Diffusion 

A strained single quantum well of GaAs/Ga 0.77 In 0.23 Ga/GaAs was grown, via lowpressure 

metal-organic vapor phase epitaxy, onto a (100)GaAs substrate at 625C. Samples 

were annealed under AsH 3 /H 2 at temperatures ranging from 750 to 900C. Since the 

quantum well thickness of about 8nm was below the critical value for this latticemismatched 

system, it was assumed that the GaInAs layer was commensurate with the 

GaAs substrate. The low-temperature (2K) photoluminescence of the electron to heavy 

hole transition in the quantum well of these samples was analyzed in order to study In/Ga 

interdiffusion at the GaInAs/GaAs interfaces. The energies of the photoluminescence 

peaks shifted to higher values during annealing. These shifts were quantitatively 

interpreted in terms of changes in the quantum well profile, due to In and Ga 

interdiffusion. The interdiffusion coefficient at 850C was deduced to be 3 x 10 -17 cm 2 /s; 

with an activation energy of 2.07eV. The values which were obtained for the In/Ga 

interdiffusion coefficient were larger than the published values for Al and Ga 

interdiffusion in AlGaAs/GaAs heterojunctions. 

F.Iikawa, P.Motisuke, J.A.Brum, M.A.Sacilotti, A.P.Roth, R.A.Masut: Journal of Crystal 

Growth, 1988, 93, 336-41 

[446-64/65-173] 

334 GaInAs/GaAs: Ga Diffusion 

The enhancement of Ga interdiffusion, due to Zn diffusion, was studied in strained 

In x Ga 1-x As/GaAs single quantum well structures, where x was between 0.20 and 0.24. 

The structures were grown by means of metalorganic vapor base epitaxy, and consisted of 

a 1000 to 2000nm-thick GaAs buffer layer plus an 8 to 10nm GaInAs layer, and a 50 to 

100nm GaAs cap layer. Shallow Zn in-diffusion was carried out at temperatures of 

between 585 and 620C, in order to vary the Zn content of the quantum wells. The 

samples were then annealed at temperatures of between 650 and 785C (in AsH 3 -H 2 

mixtures) in order to determine the enhanced In-Ga interdiffusion coefficient. The results 

for a typical case (table 34) could be described by: 

408

Ga (Ga,In)As Ga 

D(cm 2 /s) = 3 x 10 -6 exp[-2.33(eV)/kT] 

The interdiffusion enhancement was explained in terms of an interstitial migration 

process. 

M.T.Furtado, E.A.Sato, M.A.Sacilotti: Superlattices and Microstructures, 1991, 10[2], 

225-30 

[446-88/89-042] 

Table 34 

Interdiffusivity (In-Ga) in InGaAs/GaAs 


785 2.2 x 10 -17 

750 1.1 x 10 -17 

700 2.5 x 10 -18 

650 6.1 x 10 -19 

InGaAs/GaAs/AlGaAs: Ga Diffusion 

The interdiffusion of In and Ga at an InGaAs/GaAs interface was studied. The 

interdiffusion coefficients and activation energies were determined by correlating shifts in 

photoluminescence peaks with calculated quantum well transition energies that were 

based upon an erf concentration profile. The results indicated that a higher x-value, in 

In x Ga 1-x As single quantum wells, led to a higher interdiffusion coefficient for Ga under 

As over-pressure annealing conditions. Moreover, an increase in the As over-pressure 

increased the tendency to interdiffusion, whereas a Ga over-pressure reduced 

interdiffusion. The thermal activation energies for x-values of 0.057, 0.1 or 0.15 ranged 

from 3.3 to 2.6eV under an As over-pressure, and from 3 to 2.23eV under a Ga overpressure. 

K.Y.Hsieh, Y.L.Hwang, J.H.Lee, R.M.Kolbas: Journal of Electronic Materials, 1990, 

19[12], 1417-23 

[446-84/85-052] 

InGaAs/InAlAs: Ga Diffusion 

It was shown that implantation of O (a basically non-dopant impurity), after adequate hightemperature 

furnace annealing (750C, 1h), led to significant interdiffusion of group-III 

atoms in molecular beam epitaxially grown In 0.53 Ga 0.47 As-In 0.52 Al 0.48 As multi-quantum 

wells. Photoluminescence and Auger electron spectroscopic measurements (coupled with 

Ar + ion etching) were used to monitor the disordering of multi-quantum wells which were 

implanted with O (5 x 10 13 to 5 x 10 14 /cm 2 ) and then annealed by using either rapid 

thermal annealing or long-term furnace annealing. The role which was played by O in 

409

Ga (Ga,In)As H 

enhancing Ga interdiffusion was unambiguously established, and a tentative explanation 

for this was based upon the possible migration of O in these multi-quantum wells. 


57[21], 2190-2 

[446-78/79-043] 

Table 35 

Diffusivity of H in InGaAs 


605 2.1 x 10 -9 

550 7.9 x 10 -10 

500 2.7 x 10 -10 

InGaAs/InP: Ga Diffusion 


Marked interdiffusion of In and Ga host atoms, as well as Zn gettering at the interface, 

were analyzed in terms of kick-out and vacancy mechanisms. The activation energy for 

Zn-stimulated Ga interdiffusion across the InGaAs/InP heterojunction was estimated to be 

3.8eV. 


931-3 

[446-74-039] 

H 

InGaAs: H Diffusion 

The migration of H was studied in In 0.53 Ga 0.47 As. The H diffusivity was much higher in 

p-type than in n-type samples. It was concluded that H was a deep donor in this material. 

E.M.Omeljanovsky, A.V.Pakhomov, A.Y.Polyakov, O.M.Borodina, E.A.Kozhukhova, 

A.Y.Nashelskii, S.V.Yakobson, V.V.Novikova: Solid State Communications, 1989, 

72[5], 409-11 

[446-72/73-035] 

335 InGaAs: H Diffusion 

The diffusion length of H in C-doped material was estimated from resistance variations 

which occurred in samples with n-type cap layers during annealing in an N 2 ambient. It 

was found that the resultant diffusion lengths were proportional to the square root of the 

annealing time; as in a normal diffusion process. The activation energy for H diffusion 

was estimated to be 1.2eV (table 35). The effective diffusion coefficients were smaller 

410

H (Ga,In)As In 

than those in GaAs, thus implying that group-III atoms strongly affected H diffusion in C- 

doped compound semiconductors. 

H.Ito: Japanese Journal of Applied Physics, 1996, 35[2-9B], L1155-7 

[446-141/142-110] 

InGaAs/AlGaAs: H Diffusion 

The effects of monatomic H diffusion in quantum wells were studied by using 

photoluminescence and secondary ion mass spectroscopic methods. The multiple 

quantum well structures were grown by means of molecular beam epitaxy and were 

hydrogenated with a remote plasma. A significant increase in the 77K photoluminescence 

integrated intensity of bound excitons was observed after hydrogenation. This was 

attributed to the passivation of non-radiative recombination centers within the quantum 

wells. The studies demonstrated that there was an increase in passivation efficiency with 

increasing Al concentration in the barriers, and that the hydrogenation was stable to 

temperatures above 450C. Overall, the results all strongly suggested that the passivated 

non-radiative recombination centers were interface defects. 

S.M.Lord, G.Roos, J.S.Harris, N.M.Johnson: Journal of Applied Physics, 1993, 73[2], 

740-8 

[446-106/107-113] 

In 

GaInAs/GaAs: In Diffusion 

The enhancement of In interdiffusion, due to Zn diffusion, was studied in strained In x 

Ga 1-x As/GaAs single quantum well structures, where x was between 0.20 and 0.24. The 

structures were grown by means of metalorganic vapor base epitaxy, and consisted of a 

1000 to 2000nm-thick GaAs buffer layer plus an 8 to 10nm GaInAs layer, and a 50 to 


between 585 and 620C, in order to vary the Zn content of the quantum wells. The 

samples were then annealed at temperatures of between 650 and 785C (in AsH 3 -H 2 

mixtures) in order to determine the enhanced In-Ga interdiffusion coefficient. The results 

for a typical case could be described by: 

D(cm 2 /s) = 3 x 10 -6 exp[-2.33(eV)/kT] 


process. 


225-30 

[446-88/89-042] 

GaInAs/GaAs: In Diffusion 

A strained single quantum well of GaAs/Ga 0.77 In 0.23 Ga/GaAs was grown, via low-pressure 

metal-organic vapor phase epitaxy, onto a (100)GaAs substrate at 625C. Samples were 

annealed under AsH 3 /H 2 at temperatures ranging from 750 to 900C. Since the quantum 

411

In (Ga,In)As In 

well thickness of about 8nm was below the critical value for this lattice-mismatched 

system, it was assumed that the GaInAs layer was commensurate with the GaAs substrate. 

The low-temperature (2K) photoluminescence of the electron to heavy hole transition in 

the quantum well of these samples was analyzed in order to study In/Ga interdiffusion at 

the GaInAs/GaAs interfaces. The energies of the photoluminescence peaks shifted to 

higher values during annealing. These shifts were quantitatively interpreted in terms of 

changes in the quantum well profile, due to In and Ga interdiffusion. The interdiffusion 

coefficient at 850C was deduced to be 3 x 10 -17 cm 2 /s; with an activation energy of 

2.07eV. The values which were obtained for the In/Ga interdiffusion coefficient were 

larger than the published values for Al and Ga interdiffusion in AlGaAs/GaAs 

heterojunctions. 

F.Iikawa, P.Motisuke, J.A.Brum, M.A.Sacilotti, A.P.Roth, R.A.Masut: Journal of Crystal 

Growth, 1988, 93, 336-41 

[446-64/65-173] 

InGaAs/GaAs/AlGaAs: In Diffusion 

The interdiffusion of In and Ga at an InGaAs/GaAs interface was studied. The 

interdiffusion coefficients and activation energies were determined by correlating shifts in 

photoluminescence peaks with calculated quantum well transition energies that were 

based upon an erf concentration profile. The results indicated that a higher x-value, in 

In x Ga 1-x As single quantum wells, led to a higher interdiffusion coefficient for In under As 

over-pressure annealing conditions. Moreover, an increase in the As over-pressure 

increased the tendency to interdiffusion, whereas a Ga over-pressure reduced 

interdiffusion. The thermal activation energies for x-values of 0.057, 0.1 or 0.15 ranged 

from 3.3 to 2.6eV under an As over-pressure, and from 3 to 2.23eV under a Ga overpressure. 

K.Y.Hsieh, Y.L.Hwang, J.H.Lee, R.M.Kolbas: Journal of Electronic Materials, 1990, 

19[12], 1417-23 

[446-84/85-052] 

InGaAs/InAlAs: In Diffusion 

The distribution of group-III metals at In 0.53 Ga 0.47 As/In 0.52 Al 0.48 As interfaces, before and 

after annealing at 1085K, were measured. Little evidence for Al interdiffusion was found, 

but the Ga concentration profiles exhibited some broadening after annealing. The almost 

constant original In profiles developed strong modulations; with near-discontinuities at 

the initial interfaces. This behavior was explained in terms of In diffusion in the chemical 

potential gradient which was established by the disparity in Al and Ga mobilities and by 

the requirement for III-V stoichiometry in the alloys. 

R.J.Baird, T.J.Potter, G.P.Kothiyal, P.K.Bhattacharya: Applied Physics Letters, 1988, 

52[24], 2055-7 

[446-62/63-223] 

412

In (Ga,In)As Mn 

InGaAs/InAlAs: In Diffusion 

It was shown that implantation of O (a basically non-dopant impurity), after adequate 

high-temperature furnace annealing (750C, 1h), led to significant interdiffusion of group- 

III atoms in molecular beam epitaxially grown In 0.53 Ga 0.47 As-In 0.52 Al 0.48 As multiquantum 

wells. Photoluminescence and Auger electron spectroscopic measurements 

(coupled with Ar + ion etching) were used to monitor the disordering of multi-quantum 

wells which were implanted with O (5 x 10 13 to 5 x 10 14 /cm 2 ) and then annealed by using 

either rapid thermal annealing or long-term furnace annealing. The role which was played 

by O in enhancing In interdiffusion was unambiguously established, and a tentative 

explanation for this was based upon the possible migration of O in these multi-quantum 

wells. 


57[21], 2190-2 

[446-78/79-043] 

Mg 

InGaAs/InP: Mg Diffusion 









by: 

D Be /D Mn ˜ 20, D Mg /D Be ˜ 1 


[446-55/56-028] 

Mn 

InGaAs/InP: Mn Diffusion 








413

Mn (Ga,In)As Si 


by: 

D Be /D Mn ˜ 20, D Mg /D Be ˜ 1 


[446-55/56-028] 

P 

InGaAs/InGaAsP: P Diffusion 

A photoluminescence study was made of the interdiffusion of P in the 

In 0.66 Ga 0.33 As/In 0.66 Ga 0.33 As 0.7 P 0.3 system at temperatures ranging from 950 to 600C. It 

was shown that the diffusion was Fickian, with no dependence of the diffusion coefficient 

upon the substrate dopant-type or etch-pit density. In the case of Sn- or S-doped 

substrates, the diffusion could be described by: 

D(cm 2 /s) = 23 exp[-3.7(eV)/kT] 

at temperatures greater than 675C. This activation energy was the same as that deduced 

for group-III interdiffusion in Ga 0.8 In 0.2 As/GaAs. At lower temperatures, diffusivity in the 

present system could be described by: 

D(cm 2 /s) = 5 x 10 -10 exp[-1.7(eV)/kT] 

W.P.Gillin, S.S.Rao, I.V.Bradley, K.P.Homewood, A.D.Smith, A.T.R.Briggs: Applied 


[446-106/107-116] 

Si 

InGaAs: Si Diffusion 

The behavior of implanted Si + ions was investigated during rapid thermal annealing at 

temperatures of between 600 and 900C. The apparent activation energy for Si was equal 

to 0.64eV. Higher activation efficiencies were found for the dopant in InGaAs, as 

compared with InAlAs. The Si underwent no migration, even after annealing at 850C. 


Physics, 1992, 71[1], 215-20 

[446-86/87-038] 

InGaAs/InP: Si Diffusion 

It was found that Si-induced mixing produced comparable anion and cation interdiffusion. 

This was consistent with a di-vacancy mechanism. The mixing depended markedly upon 

the Si content and occurred above the diffusion shoulder, where the vacancy and defect 

pair concentrations were greatly enhanced. The absence of strain-related growth defects in 

414

Si (Ga,In)As Zn 

the un-strained starting material permitted the formation of a high-quality strained-layer 

superlattice by mixing. 

S.A.Schwarz, P.Mei, T.Venkatesan, R.Bhat, D.M.Hwang, C.L.Schwartz, M.Koza, 

L.Nazar, B.J.Skromme: Applied Physics Letters, 1988, 53[12], 1051-3 

[446-62/63-223] 

Ti 

InGaAs: Ti Diffusion 

Samples of n-type In 0.53 Ga 0.47 As were implanted with Co and Fe, and p-type samples of 

the same material were implanted with Ti. In the case of high-temperature single-energy 

Co and Fe implantation, no satellite peaks were observed at locations such as 0.8R, R + 

dR, or 2R, where R was the projected range and dR was the straggle of the implant. 

During high-temperature annealing, out-diffusion of the implant was as severe as that for 

room-temperature implants. In-diffusion of the implant also occurred, but it was not as 

severe as the out-diffusion. High-temperature annealing of Ti-implanted material resulted 

in slight Ti in-diffusion, with minimal redistribution or out-diffusion. In the case of hightemperature 

implants, the lattice quality of the annealed material was close to that of 

virgin material. Regardless of the ion type, resistivities that were close to the intrinsic 

limit were measured in implanted and annealed materials. 

M.V.Rao, S.M.Gulwadi, S.Mulpuri, D.S.Simons, P.H.Chi, C.Caneau, W.P.Hong, 

O.W.Holland, H.B.Dietrich: Journal of Electronic Materials, 1992, 21[9], 923-8 

[446-93/94-038] 

Zn 

InGaAs: Zn Diffusion 

The diffusion of Zn into In 0.57 Ga 0.43 As was studied by using boat diffusion, diffusion 

from As- or P-doped spun-on films, or diffusion from In-doped spun-on films. The depth 

profiles were deduced from junction positions. It was found that the p + /p - junction 

position depended upon the diffusion method which was used, but not upon the sample 

growth technique. The p - /n junction position depended upon both factors. Because the 

amounts of As, Ga and In (or of the respective vacancies) differed, it was possible to 

identify diffusion mechanisms. It was proposed that interstitially diffusing Zn was 

independently trapped by 2 immobile vacancy centers. These consisted of Zn on V In or 

V Ga in the p + region, and Zn on V As ZnV As in the p - region. 

U.König, H.Haspeklo, P.Marschall, M.Kuisl: Journal of Applied Physics, 1989, 65[2], 

548-52 

[446-72/73-035] 

415

Zn (Ga,In)As Zn 


A secondary ion mass spectroscopic study demonstrated that, during the growth of Zndoped 

In 0.53 Ga 0.47 As layers by atmospheric-pressure organometallic vapor-phase epitaxy, 

Zn atoms which were trapped on interstitial sites during growth, rather than interstitial Zn 

defects which were generated by the kick-out mechanism, were probably the main cause 

of the carrying over of Zn into subsequent layers. The interstitial Zn incorporation 

seemed to be due to the saturation of In substitution, rather than of Ga substitution. It was 

possible that the use of pauses in the growth sequence, both before and after the growth 

of a Zn-doped layer, could control these effects. 

S.J.Taylor, B.Beaumont, J.C.Guillaume: Semiconductor Science and Technology, 1993, 

8[12], 2193-6 

[446-115/116-143] 


Secondary ion mass spectrometry and electrochemical profiling studies were made of the 

saturation of Zn doping and diffusion in In 0.53 Ga 0.47 As which had been grown onto InP by 

using organometallic vapor-phase epitaxy. It was found that the results were consistent 

with a so-called kick-out mechanism. It was proposed that the diffusing species was 

probably a neutral Zn interstitial. Accumulation of Zn at the interface with a highly n- 

doped layer indicated the possible formation of Zn-donor complexes. 

S.J.Taylor, B.Beaumont, J.C.Guillaume: Semiconductor Science and Technology, 1993, 

8[5], 643-6 

[446-115/116-143] 

GaInAs: Zn Diffusion 

A systematic study of Zn incorporation showed that the surface Zn concentration in 

metalorganic vapor phase epitaxial material could be increased from a grown-in 

maximum of 2 x 10 19 , to 10 20 /cm 3 , by diffusion from a spun-on glass source. These 

values were found to depend upon the type of p-dopant element which was used in asgrown 

double heterostructures. 

D.L.Murrell: Semiconductor Science and Technology, 1990, 5[5], 414-20 

[446-74-030] 

GaInAs/GaAs: Zn Diffusion 

The enhancement of In and Ga interdiffusion, due to Zn diffusion, was studied in strained 

In x Ga 1-x As/GaAs single quantum well structures, where x was between 0.20 and 0.24. The 

structures were grown by means of metalorganic vapor base epitaxy, and consisted of a 

1000 to 2000nm-thick GaAs buffer layer plus an 8 to 10nm GaInAs layer, and a 50 to 


between 585 and 620C, in order to vary the Zn content of the quantum wells. The samples 

were then annealed at temperatures of between 650 and 785C (in AsH 3 -H 2 mixtures) in 

416


order to determine the enhanced In-Ga interdiffusion coefficient. The results for a typical 

case could be described by: 

D(cm 2 /s) = 3 x 10 -6 exp[-2.33(eV)/kT] 


process. 


225-30 

[446-88/89-042] 

Table 36 

Interdiffusion Coefficients for In x Ga 1-x As/GaAs 

x 

D (cm 2 /s) 

0.23 1.5 x 10 -18 

0.24 7.9 x 10 -17 

0.21 3.2 x 10 -16 

InGaAs/AlGaInAs: Zn Diffusion 

Diffusion-induced disordering of multiple quantum wells was investigated as a possible 

new processing technique for long-wavelength opto-electronic devices. Complete 

disordering of the multiple quantum well structure was confirmed by an observed 

shortening of the photoluminescence peak wavelength and by secondary ion mass 

spectrometry data. Lattice-matched disordering was also detected by using X-ray 

diffraction techniques. The first long-wavelength buried multiple quantum well laser was 

fabricated, in which carrier confinement was obtained by disordering. 

K.Goto, F.Uesugi, S.Takahashi, T.Takiguchi, E.Omura, Y.Mihashi: Japanese Journal of 


[446-117/118-184] 

336 InGaAs/GaAs: Zn Diffusion 

An investigation was made, using a 2-stage Zn diffusion and thermal annealing process, 

of impurity-induced disordering in strained samples of In x Ga 1-x As/GaAs single quantum 

well heterostructures, where x ranged from 0.21 to 0.24. The samples were grown by 

means of metalorganic vapor-phase epitaxy, and various shallow Zn diffusion depths 

were produced in the GaAs cap layer in order to vary the Zn concentration in the quantum 

wells. Thermal annealing (785C, 600s) in an AsH 3 /H 2 atmosphere then produced 

impurity-induced disordering. Partially disordered and completely disordered quantum 

well heterostructures were studied, and In-Ga interdiffusion was monitored via the 

photoluminescence spectroscopy of ground state emissions from the quantum wells. The 

interdiffusion coefficients (table 36) were determined by applying the envelope function 

to shifts in the photoluminescence peak position during annealing. It was found that the 

interdiffusion coefficient was very strongly dependent upon the Zn diffusion depth, and 

417


therefore upon the Zn concentration in the quantum well layer. A model was proposed 

that used an interstitial migration process to explain the enhancement of In-Ga 

interdiffusion by Zn diffusion. 

M.T.Furtado, M.S.S.Loural, E.A.Sato, M.A.Sacilotti: Semiconductor Science and 

Technology, 1992, 7[6], 744-51 

[446-99/100-092] 

InGaAs/InGaAsP: Zn Diffusion 

The stability of the Zn profile in modulation-doped multiple quantum well structures, 

which had been grown by means of low-pressure metalorganic vapor-phase epitaxy, was 

investigated by applying secondary ion mass spectrometric and transmission electron 

microscopic techniques to wedge-shaped samples. Although an excellent stability of the 

Zn profile was observed in as-grown samples with modulation doping (3nm, 10 18 Zn/cm 3 ), 

the modulation-doped structure faded after the second epitaxial re-growth of a p-type InP 

layer (10 18 Zn/cm 3 ) using either liquid-phase epitaxial or metalorganic vapor-phase 

epitaxial techniques. However, the modulation-doping profile was successfully preserved 

even after re-growth of the p-type InP layer for 1.5h (in a sample that comprised an 

undoped InP-clad layer, instead of a p-type InP clad layer, superposed on the modulationdoped 

multiple quantum well structure layers). The Zn diffusion coefficient in the 

modulation-doped region was less than 7 x 10 -18 cm 2 /s. The maximum Zn concentration, 

for obtaining a stable modulation-doping structure in the modulation-doped region of 

barrier layers, was found to be 2 x 10 18 /cm 3 . It was proposed that the suppression of 

interstitial Zn atoms, and of subsequently produced interstitial group-III atoms (which 

were generated in the p-type InP clad layer via a kick-out mechanism and diffused into 

the multiple quantum well region), was important in preserving the modulation-doped 

structure. 

N.Otsuka, M.Ishino, Y.Matsui: Journal of Applied Physics, 1996, 80[3], 1405-13 

[446-138/139-097] 

InGaAs/InP: Zn Diffusion 

Photodiodes were fabricated via selective Zn diffusion using a dimethylzinc source. The 

results showed that this method was a useful process for fabricating such devices. 

M.Wada, M.Seko, K.Sakakibara, Y.Sekiguchi: Japanese Journal of Applied Physics, 

1990, 29[3], L401-4 

[446-76/77-024] 


A closed-ampoule technique was used to study the simultaneous diffusion of Cd and Zn 

into material which was lattice-matched to InP. It was found that the Cd atoms, whose 

interstitial diffusion coefficient was small when compared with that of Zn interstitials, 

penetrated into the crystal as deeply as the Zn atoms. These data agreed with the results of 

418


numerical simulations of simultaneous diffusion, in III-V compounds, which were based 

upon an interstitial-substitutional diffusion mechanism. 

U.Wielsch, P.Ambrée, B.Gruska: Semiconductor Science and Technology, 1990, 5[9], 

923-7 

[446-76/77-024] 


The diffusion of Zn from spun-on films, and into InP/InGaAs/InP heterostructures, was 

studied. Marked segregation occurred at the InGaAs/InP heterojunctions and increased 

the Zn concentration in InGaAs by about an order of magnitude. Diffusion and 

segregation parameters were deduced from the Zn concentration profiles. 

F.Dildey, M.C.Amann, R.Treichler: Japanese Journal of Applied Physics, 1990, 29[5], 

810-2 

[446-76/77-024] 



It was found that marked changes in the group-III sub-lattice occurred, near to the 

interface, when Zn diffused across the heterojunction from an InGaAs top layer. The 

gettering of Zn at the interface was analyzed in terms of kick-out and vacancy 

mechanisms. The activation energy for Zn-stimulated Ga interdiffusion across the 

InGaAs/InP heterojunction was estimated to be 3.8eV. 


931-3 

[446-74-039] 


The Zn was introduced, using spun-on films, into n-InP/p + -InGaAs/n-InP heterostructures 

which had been grown via metalorganic vapor phase epitaxy (with Mg as a p-dopant). 

After diffusion, the Mg was completely replaced by Zn and was enriched in the spun-on 

film. In the presence of Mg, the in-diffusion of Zn was strongly enhanced. By varying the 

dopant level and diffusion conditions, the underlying mechanism was compared with that 

which operated in Be-doped AlGaAs/GaAs heterostructures. 

F.Dildey, R.Treichler, M.C.Amann, M.Schier, G.Ebbinghaus: Applied Physics Letters, 

1989, 55[9], 876-8 

[446-70/71-118] 


It was found that Zn-induced mixing mainly affected the cation sub-lattice of the 

superlattice and was consistent with a so-called interstitial kick-out mechanism. The Zndiffused 

superlattice remained sharply defined. The absence of strain-related growth 

419

Zn (Ga,In)As General 

defects in the un-strained starting material permitted the formation of a high-quality 

strained-layer superlattice by mixing. 

S.A.Schwarz, P.Mei, T.Venkatesan, R.Bhat, D.M.Hwang, C.L.Schwartz, M.Koza, 

L.Nazar, B.J.Skromme: Applied Physics Letters, 1988, 53[12], 1051-3 

[446-62/63-223] 


The intermixing of multiple quantum well structures by Zn diffusion at 550C was 

investigated. Secondary ion mass spectroscopy and X-ray analysis revealed that Zn 

diffusion promoted the intermixing of group-III atoms, but had little effect upon group-V 

profiles. However, the resultant group-III atom profiles were not completely uniform; 

even after Zn diffusion. The results suggested that a large lattice mismatch suppressed the 

Zn diffusion intermixing process. 

K.Nakashima, Y.Kawaguchi, Y.Kawamura, Y.Imamura, H.Asahi: Applied Physics 

Letters, 1999, 52[17], 1383-5 

[446-62/63-223] 


A systematic study was made of the effects of Zn doping and diffusion in capped mesa 

buried heterostructure lasers which had been grown by means of metalorganic chemical 

vapor deposition. It involved varying the Zn content (7 x 10 17 to 3.1 x 10 18 /cm 3 ) of the p- 

type InP cladding layer in the base epitaxial structure, while keeping the growth 

conditions constant during 2 subsequent re-growth steps. Secondary ion mass 

spectrometry was used to make quantitative determinations of the Zn depth profiles, 

following re-growth, by using test sites on 50mm round wafers which contained the 

appropriate epitaxial layers. Clear evidence of Zn diffusion was found, such as the 

penetration of Zn into the active layer and the presence of inflection points (accumulation 

and depletion of Zn near to the p-n heterojunction) in the depth profiles. It was observed 

that the diffusion of Zn during the third growth step dominated the Zn profile in the base 

growth part of the p-type InP layer, and the final amount of Zn in this region was 

independent of the initial dopant level. Above a Zn threshold level of about 2.2 x 

10 18 /cm 3 , the Zn diffusion increased significantly and resulted in the presence of 5 x 

10 18 /cm 3 or more of Zn in the active layer. The threshold for the onset of diffusion was 

found to be in accord with the substitutional-interstitial diffusion of Zn. 

V.Swaminathan, C.L.Reynolds, M.Geva: Applied Physics Letters, 1995, 66[20], 2685-7 

[446-121/122-079] 

General 

InGaAs/InGaAsP/InP: Diffusion 

Laser structures, with InGaAs quantum wells which were about 1.85µ below the surface, 

were implanted with ions that had energies of up to 8.6MeV. Intermixing of the quantum 

wells, during rapid thermal annealing, was monitored via changes in the energy, line- 

420

General (Ga,In)As Surface 

width and intensity of the photoluminescence peaks from the quantum wells. When 

diffusion occurred mainly through InP, the photoluminescence data correlated well with 

the calculated total number of vacancies that were created in the sample. This suggested 

that defect diffusion was very efficient in InP. 

P.J.Poole, S.Charbonneau, G.C.Aers, T.E.Jackman, M.Buchanan, M.Dion, 

R.D.Goldberg, I.V.Mitchell: Journal of Applied Physics, 1995, 78[4], 2367-71 

[446-123/124-177] 

InGaAs: Diffusion 

A theoretical model was proposed for the calculation of the layer thickness of a III-V 

ternary crystal which was grown by using the source-current controlled method. The 

latter could be used to control the supply of depleted solute elements to a solution during 

growth. That is, solute elements were continuously supplied to the growth solution from a 

source material and were transported to a substrate (through the temperature gradient in 

the growth solution) via diffusion and electromigration. By using the theoretical model, 

analytical calculations could be made of the diffusion-limited and electromigrationlimited 

growth of InGaAs in temperature gradient In-Ga-As ternary solutions. The 

calculations indicated that the thickness depended upon growth parameters such as the 

growth temperature, the cooling rates of the substrate and source, the solution length, the 

temperature difference between substrate and source, the mobility of the migrating solute 

element, and the electric field in the solution. 


[446-72/73-035] 

InGaAs/GaAs/AlGaAs: Diffusion 

Laser structures, with InGaAs quantum wells which were about 1.85µ below the surface, 

were implanted with ions that had energies of up to 8.6MeV. Intermixing of the quantum 

wells, during rapid thermal annealing, was monitored via changes in the energy, linewidth 

and intensity of the photoluminescence peaks from the quantum wells. When the 

defects had to diffuse mainly through Al 0.71 Ga 0.29 As, these quantities were closely 

related, for short annealing times, to the predicted vacancy generation and ion deposition 

at the depth of the quantum well before annealing. This suggested that the defect 

diffusion length in AlGaAs and/or GaAs was quite low. 

P.J.Poole, S.Charbonneau, G.C.Aers, T.E.Jackman, M.Buchanan, M.Dion, 

R.D.Goldberg, I.V.Mitchell: Journal of Applied Physics, 1995, 78[4], 2367-71 

[446-123/124-177] 


Ga 

InGaAs: Ga Surface Diffusion 

Surface diffusion during molecular beam epitaxy was studied. Firstly, the mode transition 

between 2-dimensional nucleation and step flow during molecular beam epitaxial growth 

421

Surface (Ga,In)As Interdiffusion 

on vicinal surfaces was studied theoretically and experimentally. The basis of the theory 

was to assume that the transition occurred when the surface supersaturation on the step 

terrace became identical to the critical supersaturation for 2-dimensional nucleation. This 

permitted the diffusion length of Ga to be calculated at the experimentally determined 

critical temperature for the mode transition. It was found that the diffusion length 

increased, as the temperature decreased, due to an increased residence time. Also, the 

diffusion length on (111)B was longer than that on (001) when the same formation energy 

for 2-dimensional nuclei was assumed for both surfaces. The theory was then used to 

elucidate the dependence of the InGaAs composition upon the growth temperature, the 

substrate orientation, and the degree of misorientation. The theory gave good agreement 

with the experimental data, and it was concluded that surface diffusion was one of the 

most important processes which controlled molecular beam epitaxial growth and impurity 

incorporation. 

T.Nishinaga, T.Shitara, K.Mochizuki, K.I.Cho: Journal of Crystal Growth, 1990, 99, 482- 

90 

[446-76/77-009] 

In 

GaInAs/GaAs: In Surface Diffusion 

The lateral profiles of In in a 1.5µ-thick Ga 1-x In x As layers, where x was approximately 

equal to 0.2, were grown onto GaAs channelled substrates with (411)A side-slopes by the 

use of molecular beam epitaxy and were investigated with the use of energy-dispersive X- 

ray spectroscopy. The observed profiles of the In content suggested that the In atoms 

migrated preferentially in the [12¯2¯] direction on the (411)A plane during molecular beam 

epitaxial growth. This preferential migration of In atoms was confirmed by comparing the 

observed lateral profiles of In in GaInAs layers which had been grown onto GaAs 

channelled substrates with simulated In profiles that had been calculated by taking 

account of an additional one-way flow of In atoms along [12¯2¯]. 

T.Kitada, A.Wakejima, N.Tomita, S.Shimomura, A.Adachi, N.Sano, S.Hiyamizu: Journal 

of Crystal Growth, 1995, 150[1-4], 487-91 

[446-127/128-141] 


GaInAs/AlGaInAs: Interdiffusion 

Investigations were made of the interdiffusion behavior and thermal stability of n-doped 

and undoped multiple quantum-well structures that were lattice-matched to InP by means 

of molecular beam epitaxy. The activation energy for the main interdiffusion process was 

deduced to be 2.5eV for doped structures and 2.9eV for undoped structures. The differing 

interdiffusion processes were monitored by using photoluminescence spectroscopy at 8K, 

after rapid thermal annealing. The effect of doping was studied by comparing the results 

422

Interdiffusion (Ga,In)As Interdiffusion 

for n-doped and undoped structures. Photoluminescence excitation spectroscopic data for 

2K confirmed the differing interdiffusion processes. 

V.Hofsäss, J.Kuhn, H.Schweizer, H.Hillmer, R.Lösch, W.Schlapp: Journal of Applied 

Physics, 1995, 78[5], 3534-6 

[446-123/124-172] 

GaInAs/AlInAs: Interdiffusion 

The interaction of lattice-matched heterostructures, grown by means of molecular beam 

epitaxy, was studied by using electron microscopy. An appreciable amount of 

interdiffusion was observed at temperatures as low as 700C. The use of X-ray microanalysis 

revealed that interdiffusion occurred along a non-linear (non lattice-matched) 

path. 

R.E.Mallard, N.J.Long, G.R.Booker, E.G.Scott, M.Hockly, M.Taylor: Journal of Applied 

Physics, 1991, 70[1], 182-92 

[446-91/92-018] 

1.0E-14 

D (cm 2 /s) 

1.0E-15 

1.0E-16 

1.0E-17 

table 34 

table 37 

table 38 

table 39 

table 40 

table 41 

table 42 

1.0E-18 

1.0E-19 

7 8 9 10 11 

10 4 /T(K) 

Figure 8: Interdiffusivity in InGaAs/GaAs 

423


GaInAs/GaAs: Interdiffusion 

The dependence of impurity-free interdiffusion upon the properties of a dielectric cap 

layer was studied in pseudomorphic multi-quantum well structures that had been grown 

by means of molecular beam epitaxy. Electron-beam evaporated SiO 2 films, chemical 

vapor deposited SiO x N y films, and spun-on SiO 2 films were used as cap layers during 

rapid thermal annealing at temperatures of between 850 and 950C. The 

photoluminescence at 10K was used to monitor interdiffusion-induced band-gap shifts, 

and to calculate the corresponding In-Ga interdiffusion coefficients. The latter were found 

to increase with cap layer thickness (electron-beam SiO 2 ) up to a limit which was 

governed by saturation of the out-diffused Ga concentration in the SiO 2 caps. A 

maximum concentration of between 4 x 10 19 and 7 x 10 19 /cm 3 in the SiO 2 caps was found 

by using secondary ion mass spectroscopic profiling. Larger band-edge shifts were also 

obtained when the O content of SiO x N y cap layers was increased, but the differences were 

insufficient to suggest a laterally selective interdiffusion process that was based upon 

variations in cap layer composition alone. Much larger differences were obtained by 

using various deposition techniques for the cap layers. This indicated that the porosity of 

the cap layer was a much more important factor than was the film composition in 

obtaining a laterally selective interdiffusion process. In the case of Ga 0.8 In 0.2 As/GaAs 

interdiffusion, the activation energies and pre-factors were estimated to range from 3.04 

to 4.74eV and from 5 x 10 -3 to 2 x 10 5 cm 2 /s, respectively; depending upon the cap layer 

deposition technique and the depth of the multi-quantum well below the sample surface. 

S.Bürkner, M.Maier, E.C.Larkins, W.Rothemund, E.P.O’Reilly, J.D.Ralston: Journal of 


[446-125/126-131] 

337 GaInAs/GaAs: Interdiffusion 

The effects of Si and Be, at doping levels of up to 10 19 /cm 3 , upon the interdiffusion of 

quantum wells after annealing were studied by using photoluminescence techniques (table 

37). It was found that, for Be concentrations of up to 2.5 x 10 19 /cm 3 , and for Si 

concentrations of up to 10 18 /cm 3 , no change in the interdiffusion coefficients could be 

measured. At a Si dopant concentration of 6 x 10 18 /cm 3 , there was a dramatic degradation 

of the material quality after annealing (750C, 15s). This caused the luminescence from 

the well to disappear, while a deep-level luminescence that was related to donor-Ga 

vacancy complexes and As antisite defects appeared. On the basis of these results, it was 

suggested that the position of the Fermi level played no role in the intermixing of III-V 

heterostructures. It was also concluded that most of the enhanced intermixing which was 

observed in Si-doped GaAs/AlGaAs structures was related to Si relocation at very high 

doping levels. 

W.P.Gillin, I.V.Bradley, L.K.Howard, R.Gwilliam, K.P.Homewood: Journal of Applied 

Physics, 1993, 73[11], 7715-9 

[446-106/107-082] 

424


Table 37 

Interdiffusion Data for Ga 0.8 In 0.2 As/GaAs 

Dopant Amount (/cm 3 ) Temperature (C) Coefficient (cm 2 /s) 

- - 900 4.70 x 10 -17 

- - 950 1.96 x 10 -16 

- - 1000 5.50 x 10 -16 

- - 1050 1.00 x 10 -15 

Si 10 17 900 3.50 x 10 -17 

Si 10 17 950 1.26 x 10 -16 

Si 10 17 1000 3.90 x 10 -16 

Si 10 17 1050 7.50 x 10 -16 

Si 10 18 900 1.40 x 10 -16 

Si 10 18 950 4.00 x 10 -16 

Si 10 18 1000 1.20 x 10 -15 

Si 10 18 1050 2.80 x 10 -15 

Be 10 17 900 5.40 x 10 -17 

Be 10 17 950 1.50 x 10 -16 

Be 10 17 1000 2.50 x 10 -16 

Be 10 17 1050 1.60 x 10 -15 

Be 10 18 900 1.00 x 10 -17 

Be 10 18 950 1.02 x 10 -16 

Be 10 18 1000 4.20 x 10 -16 

Be 10 18 1050 1.33 x 10 -15 

Be 2.5 x 10 19 900 3.70 x 10 -17 

Be 2.5 x 10 19 950 1.64 x 10 -16 

Be 2.5 x 10 19 1000 5.00 x 10 -16 

Be 2.5 x 10 19 1050 2.40 x 10 -15 


Simple analytical expressions were derived for the approximate estimation of the 

interdiffusion coefficient, of partially disordered quantum-well heterostructures, directly 

from measurements of the photoluminescence phase shift which was associated with layer 

interdiffusion. The phase shift was calculated as a function of the interdiffusion length, 

(Dt) ½ , in the strained-layer system, Ga 0.8 In 0.2 As/GaAs. The calculations were performed 

within the framework of the envelope function approximation and Fick's law. A simple 

relationship was derived for the variation in phase shift as a function of the dimensionless 

425


parameter, (Dt) ½ /L, where L was the quantum-well thickness. This satisfactorily 

accounted for most of the published interdiffusivity values, to within a factor of 2. 

M.T.Furtado, M.S.S.Loural: Superlattices and Microstructures, 1993, 14[1], 21-5 

[446-113/114-029] 


Molecular-beam epitaxially grown highly strained Ga 0.65 In 0.35 As/GaAs multiple quantumwell 

structures were investigated. Interdiffusion was carried out via rapid thermal annealing, 

at temperatures of between 700 and 950C, by using GaAs proximity caps and electron-beam 

evaporated SiO 2 cap layers, respectively. Reduced photoluminescence line-widths and 

increased photoluminescence intensities were observed after diffusion-induced band-gap 

shifts that ranged from 0.006 to 0.220eV. Microscopic photoluminescence methods were 

used to study the onset of strain relaxation due to dislocation generation. Two types of line 

defect were found in samples which had been annealed using proximity caps; depending 

upon the annealing temperature and the number of quantum wells. These were misfit 

dislocations with their lines parallel to directions, and -oriented line defects. 

No dislocations were observed, in samples which had been annealed using a SiO 2 cap, over 

the entire temperature range which was investigated here. Resonant Raman scattering 

measurements of the 1LO/2LO phonon intensity ratio were used to make semi-quantitative 

assessments of the total defect content; including point defects. It was found that, whereas 

increasing point defect densities, and the formation of line defects, were observed in 

proximity-capped samples as the annealing temperature was increased, no deterioration of 

structural quality due to an increased point defect density was observed in samples which 

had been annealed using a SiO 2 cap. 

S.Bürkner, M.Baeumler, J.Wagner, E.C.Larkins, W.Rothemund, J.D.Ralston: Journal of 


[446-134/135-139] 

InGaAs/GaAs/Si: Interdiffusion 

Heterostructures of the form, In x Ga 1-x As(5nm)/GaAs(5nm)/Si, were fabricated via 

metalorganic chemical vapor deposition and were annealed at 750C. A transmission electron 

microscopic study was made of the heterostructure, with particular attention being paid to 

interdiffusion between group-III elements. Plan-view transmission electron microscopy 

revealed that the epitaxial In x Ga 1-x As /GaAs layer consisted basically of small island 

crystals. The morphology of the islands varied as a function of the In content. When x was 

greater than 0.1, the islands coalesced into larger ones; leaving small regions of so-called 

sea. The spacings of 022 moiré fringes varied spatially as a result of variations in In content 

in the epilayers. Cross-sectional transmission electron microscopic observations showed that 

there was no sharp heteroboundary between the In x Ga 1-x As layer and the GaAs layer. The 

contrast of the 002 dark-field image was sensitive to the In content and revealed 

interdiffusion between In and Ga. 

K.Kamei, K.Fujita, Y.Shiba, H.Katahama, Y.Maehara: Defect and Diffusion Forum, 

1993, 95-98, 977-82 

[446-95/98-977] 

426


GaInAs/GaInAsP: Interdiffusion 

Electron microscopy was used to characterize metalorganic chemical vapor deposited 

multiple quantum well structures which underwent a so-called blue shift in luminescence 

during thermal processing. The sample exhibited a shift, towards shorter wavelengths, of 

more than 100nm during annealing at 750C. The structural modifications which led to the 

blue shift included the elimination of atomic ordering in the quaternary barrier layers of 

the material, plus appreciable layer interdiffusion. A method was described by which 

quantitative analyses of the layer composition and lattice parameter could be obtained, at 

a spatial resolution of better than 2nm, by making energy-dispersive X-ray microanalyses 

in a scanning transmission electron microscope. Such analyses showed that interdiffusion 

occurred along a non-linear (non lattice-matched) path, where the group-V diffusivities 

exceeded those of group-III elements. This resulted in the incorporation of excess 

coherency strains, of up to 0.5%, in the quantum-well regions. 

R.E.Mallard, N.J.Long, E.J.Thrush, K.Scarrott, A.G.Norman, G.R.Booker: Journal of 


[446-106/107-091] 

InGaAs/GaAs: Interdiffusion 

The compositional disordering of superlattices which had been grown by means of 

molecular beam epitaxy, using a low-temperature GaAs cap layer, was studied. 

Disordering of the superlattice was verified by photoluminescence and double-crystal X- 

ray rocking curve measurements. The disordering mechanism was found to involve Ga 

vacancy-enhanced interdiffusion, due to the presence of the low-temperature GaAs. The 

diffusion and Schrödinger’s equations were solved numerically in order to obtain the 

compositional profile and the transition energies in the disordered quantum well, 

respectively. The simulated energy shifts for samples under various annealing conditions 

agreed well with experimental data. The calculated effective diffusivity for In-Ga 

interdiffusion involved an activation energy of 1.63eV. This was smaller than the 

activation energy, of 1.93eV, for intrinsic interdiffusion. The diffusivity for the enhanced 

In-Ga interdiffusion, due to the presence of low-temperature GaAs, was some 2 orders of 

magnitude larger than the intrinsic In-Ga diffusivity. 

J.S.Tsang, C.P.Lee, S.H.Lee, K.L.Tsai, C.M.Tsai, J.C.Fan: Journal of Applied Physics, 

1996, 79[2], 664-70 

[446-131/132-179] 


The effect of strain upon cation interdiffusion in quantum wells was described. It was 

found that Fick’s diffusion equation did not correctly describe interdiffusion in a 

heterostructure with strained layers. It was suggested that the strain altered the crystal 

defect concentration, and that the diffusivity was therefore affected by the strain. A 

diffusion equation which included the effects of strain was derived and solved 

numerically. Experimental photoluminescence peak shifts, as a function of annealing 

427


time, were closely fitted by this analysis and useful parameters such as a coefficient 

which described InGaAs/GaAs quantum well interdiffusion were deduced. 

S.W.Ryu, I.Kim, B.D.Choe, W.G.Jeong: Applied Physics Letters, 1995, 67[10], 1417-9 

[446-125/126-141] 

Table 38 

Interdiffusivity in Kr-Implanted InGaAs/GaAs 


1050 7.71 x 10 -15 

1000 3.11 x 10 -15 

950 5.90 x 10 -16 

925 3.60 x 10 -16 

900 1.10 x 10 -16 

875 9.00 x 10 -17 

825 1.50 x 10 -17 

750 8.00 x 10 -19 


Low-pressure metalorganic vapor phase epitaxially grown strained quantum well 

structures were characterized by using photoluminescence and X-ray diffraction 

techniques. It was shown that, beyond the pseudomorphic limit, these structures exhibited 

considerable Ga/In interdiffusion at the interfaces, and partial strain relaxation in the 

quantum well layers. 

A.K.Srivastava, B.M.Arora, S.Banerjee: Journal of Electronic Materials, 1994, 23[2], 

191-4 

[446-113/114-036] 

338,39,40,41 InGaAs/GaAs: Interdiffusion 

Photoluminescence techniques and repeated annealing were used to determine the 

diffusion coefficients for intermixing in quantum wells and to study the subsequent 

effects of ion implantation upon intermixing. It was shown that, after ion implantation, a 

very fast interdiffusion process occurred which was independent of the nature of the 

implanted ion and was thought to be due to the rapid diffusion of interstitials which were 

created during implantation. Following this rapid process, it was found that neither Ga 

nor Kr ions had any effect upon the subsequent interdiffusion coefficient (tables 38 and 

39). After As implantation, in addition to the initial damage-related process, an enhanced 

region of interdiffusion was observed; with a diffusion coefficient which was an order of 

magnitude greater than that of an non-implanted control wafer (tables 40 and 41). This 

enhancement was suggested to be due to the creation of group-III vacancies by As atoms 

which moved into group-V lattice sites. The fast process continued until the structure had 

broadened by about 7.5nm, whereupon the diffusion coefficient returned to the non- 

428


implanted control value. The activation energy for interdiffusion was measured at 

temperatures ranging from 1050 to 750C, and a value of 3.7eV was deduced. This value 

was found to be independent of the nature of the implanted ion. 

I.V.Bradley, W.P.Gillin, K.P.Homewood, R.P.Webb: Journal of Applied Physics, 1993, 

73[4], 1686-92 

[446-109/110-041] 

Table 39 

Interdiffusivity in Ga-Implanted InGaAs/GaAs 


1050 2.68 x 10 -15 

1000 9.10 x 10 -16 

950 2.20 x 10 -16 

925 1.30 x 10 -16 

900 4.10 x 10 -17 

875 3.10 x 10 -17 

825 4.30 x 10 -18 

750 2.00 x 10 -19 

342 InGaAs/GaAs: Interdiffusion 

Interdiffusion in quantum wells was studied by using photoluminescence methods to 

monitor the temporal development of the diffusion process in a single sample (table 42). 

Two distinct regimes were detected: a fast initial diffusion and a second, steady-state, 

diffusion. The steady-state diffusion was found to depend upon the depth of the quantum 

well from the surface, and could be correlated with published data on the diffusion of Ga 

vacancies into GaAs. 

W.P.Gillin, D.J.Dunstan, K.P.Homewood, L.K.Howard, B.J.Sealy: Journal of Applied 

Physics, 1993, 73[8], 3782-6 

[446-109/110-042] 


The interdiffusion of single quantum wells was studied as a function of temperature for 

both p-type (Be) and n-type (Si) doping. The interdiffusion of group-III elements was 

monitored via photoluminescence from the ground states of valence and conduction band 

quantum wells. Intermixing was modelled by using a Green's function method to solve the 

diffusion equation that described the evolution of well shapes during processing. It was 

deduced that the activation energy for interdiffusion was 3.4eV. 

W.P.Gillin, K.P.Homewood, L.K.Howard, M.T.Emeny: Superlattices and 

Microstructures, 1991, 9[1], 39-42 

[446-78/79-042] 

429


Table 40 

Interdiffusivity in Non-Implanted InGaAs/GaAs 


1050 3.42 x 10 -15 

1000 8.45 x 10 -16 

950 1.97 x 10 -16 

925 1.77 x 10 -16 

900 4.25 x 10 -17 

875 2.27 x 10 -17 

825 1.00 x 10 -17 

750 2.00 x 10 -19 


It was shown that the Si-doping of quantum-well materials to 10 19 /cm 3 promoted a timeand 

temperature-dependent diffusion process which was related to group-III vacancy 

formation. The effect of the formation of such vacancies upon subsequent interdiffusion 

was modelled and was shown to reproduce variations, in the diffusion coefficient as a 

function of depth, without invoking a Fermi-level model. Experiments which were 

performed on layers that were doped to between 10 17 and 10 18 /cm 3 did not reveal an 

enhanced interdiffusion; contrary to the predictions of the Fermi-level model. Other 

results suggested that the reason for this was that interdiffusion in III-V materials was not 

governed by thermal equilibrium vacancy concentrations but rather by the vacancy 

concentrations which were grown into the substrate materials. It was demonstrated that 

the position of the Fermi level played no role in III-V intermixing. 

Z.H.Jafri, W.P.Gillin: Journal of Applied Physics, 1997, 81[5], 2179-84 

[446-148/149-182] 

InGaAs/InAlAs: Interdiffusion 

Quantum well structures were grown on InP(Fe) substrates by using metalorganic vapor 

phase epitaxial techniques. Each sample contained 3 InGaAs wells, which were 2.6, 5.9, 

or 17.6nm in thickness and were separated by 24nm-thick InAlAs barrier layers. The 

samples were implanted with Si ions to uniform densities which ranged from 1.8 x 10 17 to 

3.9 x 10 19 /cm 3 over the quantum wells, and were then annealed under various conditions. 

A photoluminescence peak energy for each well was monitored in order to study 

intermixing at the interface. Blue shifts in the photoluminescence peak energy were found 

to occur within the first 15s of thermal annealing when the Si dose exceeded a critical 

value of 2 x 10 18 to 3 x 10 18 /cm 3 . The saturation value of the energy shift was governed 

mainly by the Si density, but hardly depended upon the annealing temperature and time. It 

was concluded that the defects which were formed by Si ion implantation enhanced the 

430


thermal interdiffusion of Ga and Al atoms at the InGaAs/InAlAs interface. This ended 

when the implantation-induced defects annealed out. 

S.Yamamura, R.Saito, S.Yugo, T.Kimura, M.Murata, T.Kamiya: Journal of Applied 

Physics, 1994, 75[5], 2410-4 

[446-117/118-184] 

Table 41 

Interdiffusivity in As-Implanted InGaAs/GaAs 

Temperature (C) Diffusion Type D (cm 2 /s) 

1050 steady-state 6.42 x 10 -15 

1000 steady-state 3.90 x 10 -15 

950 steady-state 6.96 x 10 -16 

950 enhanced 3.50 x 10 -15 

925 steady-state 4.50 x 10 -16 

925 enhanced 1.73 x 10 -15 

900 steady-state 3.05 x 10 -16 

900 enhanced 1.01 x 10 -15 

750 enhanced 3.30 x 10 -18 

InGaAs/InGaAsP: Interdiffusion 

An investigation was made of the effect of interdiffusion upon the photoluminescence and 

Raman spectra of a single quantum well of InGaAs (8nm wide) which was sandwiched 

between 2 large InGaAsP barriers. Firstly, an investigation was made of a blue shift of the 

recombination line (2K) which occurred after annealing at 650 or 750C for times which 

ranged from 0.25 to 2h. A single diffusivity coefficient was assumed for all of the atomic 

species, and the amount of intermixing was deduced by means of model calculations. 

Average diffusivity coefficients of 0.0095 and 0.2Å²/s were found at 650 and 750C, 

respectively. This agreed well with published data on InGaAs/InP, and suggested that the 

activation energy was 2.54eV. 

H.Peyre, F.Alsina, J.Camassel, J.Pascual, R.W.Glew: Journal of Applied Physics, 1993, 

73[8], 3760-8 

[446-106/107-117] 

InGaAs/InGaAsP: Interdiffusion 

Thermal interdiffusion on the group-V sub-lattice in quantum-well structures was studied 

by using samples which were annealed under silicon nitride encapsulation, or under 

various phosphine over-pressures. It was found that the interdiffusion length was 

comparable, under all of these conditions, with only small effects of the phosphine overpressure 

being observed. It was suggested that interdiffusion results which were obtained 

431


for nitride-capped samples could be applied to the interdiffusion which occurred during 

growth. 

W.P.Gillin, S.D.Perrin, K.P.Homewood: Journal of Applied Physics, 1995, 77[4], 1463-5 

[446-121/122-078] 

Table 42 

Interdiffusion in InGaAs/GaAs 


900 4.6 x 10 -17 

950 2.6 x 10 -16 

1000 6.4 x 10 -16 

1050 1.6 x 10 -15 

InGaAs/InP: Interdiffusion 

A formula was derived which described the interdiffusion profiles of quantum wells. It 

was shown that it accurately modelled interdiffusion in quantum wells of lattice-matched 

InGaAs. The formula took account of the differing interdiffusion coefficients between 

layers, and of the interfacial discontinuity of interdiffused species. The formula explained 

how quantum energy shifts due to interdiffusion varied with annealing time and annealing 

temperature in various wide-well layers of both InGaAsP/InP and GaAs/AlGaAs quantum 

wells. The quantitative difference between the interdiffusion profiles in these two 

materials was also demonstrated. 

K.Mukai, M.Sugawara, S.Yamazaki: Physical Review B, 1994, 50[4], 2273-82 

[446-115/116-130] 


An investigation was made of intermixing effects, in In 0.53 Ga 0.47 As single quantum wells, 

which were caused by 30keV Ar + ion beam implantation to doses which ranged from 10 12 

to 10 14 /cm 2 and by subsequent rapid thermal annealing at temperatures of between 600 

and 900C. After implantation and rapid thermal annealing at 600C, a significant increase 

was observed (in the photoluminescence emission energy of about 0.06eV) as compared 

with non-implanted heterostructures. This indicated that the intermixing was produced by 

implantation. However, in the case of rapid thermal annealing at temperatures above 

850C, the energy shifts (of up to 0.2eV) which were observed in implanted samples were 

similar to the shifts that occurred in non-implanted samples. This indicated that a 

prominent contribution arose from thermal interdiffusion. A significant decrease in the Ga 

concentration after interdiffusion was confirmed by Raman data. 

J.Oshinowo, J.Dreybrodt, A.Forchel, N.Mestres, J.M.Calleja, I.Gyuro, P.Speier, 

E.Zielinski: Journal of Applied Physics, 1993, 74[3], 1983-6 

[446-106/107-117] 

432



The effects of differing cation and anion interdiffusion rates upon the disordering of 

In 0.53 Ga 0.47 As/InP single quantum-wells were investigated by using an erf distribution to 

model the concentration profile after interdiffusion. The early stages of disordering 

caused a spatially dependent strain build-up, which could be compressive or tensile. This 

strain profile, and the concentration distribution, gave rise to interesting carrier 

confinement profiles after disordering. A significantly higher cation interdiffusion rate 

produced a red-shift of the ground-state transition energy which, with prolonged 

interdiffusion, saturated and then decreased. A significantly higher anion interdiffusion 

rate caused a blue-shift in the ground-state transition energy, and shifted the light-hole 

ground-state to above the heavy-hole ground-state. 

W.C.Shiu, J.Micallef, I.Ng, E.H.Li: Japanese Journal of Applied Physics, 1995, 34[1- 

4A], 1778-83 

[446-121/122-078] 


The interdiffusion of superlattice structures, during annealing at temperatures of between 

700 and 850C, was studied by means of Raman spectroscopy. The peak intensities and 

peak energies of InAs-, GaAs-, and InP-like longitudinal optical phonon modes varied as 

a function of annealing temperature and time. The depth profiles of group-III and group-V 

atoms were estimated quantitatively by monitoring variations in the peak energies of the 

longitudinal optical phonon modes, and in the ratios of the mode intensities. This revealed 

that the interdiffused superlattice consisted of InGaAsP of uniform composition, and InP 

barrier layers with sharp interfaces. It was found that the resultant InGaAsP well was 

roughly lattice-matched to InP (to within 0.5%). The diffusion coefficient in the well 

region was larger than that in the barrier region, and the interdiffusion could be described 

by: 

D(cm 2 /s) = 8.56 x 10 10 exp[-5.82(eV)/kT] 

Interdiffusion in this superlattice was governed by diffusion in the InP region.' 

S.J.Yu, H.Asashi, S.Emura, S.Gonda, K.Nakashima: Journal of Applied Physics, 1991, 

70[1], 204-8 

[446-93/94-039] 


The thermal stability and interdiffusion of In 0.53 Ga 0.47 As/InP surface quantum wells was 

investigated. Well-defined high-intensity photoluminescence emission spectra were 

obtained. After rapid thermal annealing (500 to 900C, 60s), strong emission energy shifts 

of up to 0.316eV were detected. By assuming a simple model for ion intermixing, the 

interdiffusion coefficient was estimated to be equal to 1.7 x 10 -14 cm 2 /s at 900C; with an 

activation energy of 1.3eV. 

J.Oshinowo, A.Forchel, D.Grützmacher, M.Stollenwerk, M.Heuken, K.Heime: Applied 


[446-88/89-044] 

433



Quantum-well structures were studied using magneto-optical transmission spectroscopy. 

The effects of dopants, overgrowth and annealing were investigated. The blue-shift effect, 

which was often observed in multiple quantum-well structures that were subjected to 

heat-treatment, was here attributed to a dominant group-V interdiffusion which could be 

suppressed by high defect densities in the substrate. The presence of Zn in an overgrown 

layer on top of the multiple quantum-well structures caused a counteractive red-shift 

effect after long annealing times. This was due to group-III diffusion. On the other hand, 

in situ Zn or S produced no observable shift in transition energies due to interdiffusion. 

This was attributed to an enhanced group-III interdiffusion that was induced by Zn 

diffusion into the multiple quantum-wells. It was concluded that very different 

interdiffusion mechanisms operated in the case of group-III and group-V elements; thus 

supporting the suggestion of vacancy-related group-V interdiffusion rather than the 

interstitialcy mechanism of group-III interdiffusion. 

S.L.Wong, R.J.Nicholas, R.W.Martin, J.Thompson, A.Wood, A.Moseley, N.Carr: Journal 


[446-134/135-148] 

434

(Ga,In)(As,P) 

As 

InGaAsP/InP: As Diffusion 

The effect of diffusion upon quantum-well structures was studied by using secondary ion 

mass spectrometry, Auger electron spectroscopy, capacitance-voltage measurements, 

photoluminescence, and X-ray diffraction methods. It was found that the interdiffusion of 

As was negligible. 

G.J.Van Gurp, W.M.Van de Wijgert, G.M.Fontijn, P.J.A.Thijs: Journal of Applied 

Physics, 1990, 67[6], 2919-26 

[446-74-040] 

Be 

InGaAsP/InP: Be Diffusion 

A study was made of dopant redistribution between a chemical beam epitaxially grown 

InGaAsP laser structure, and a metalorganic vapor phase epitaxially over-grown InP 

layer. Secondary ion mass spectroscopic data revealed that Zn and Be atoms interdiffused 

deeply into the adjacent InP layers, at a Zn doping level of 10 18 /cm 3 . A fraction of the Zn 

atoms went through the chemical beam epitaxial InP, and penetrated the laser structure 

guide layer. It was found that Zn out-diffusion was appreciably suppressed by reducing 

the Be dopant concentration from 10 18 to 5 x 10 17 /cm 3 . 

H.Sugiura, S.Kondo, M.Mitsuhara, S.Matsumoto, M.Itoh: Applied Physics Letters, 1997, 

70[21], 2846-8 

[446-152-0435] 

Ga 

InGaAsP/InP: Ga Diffusion 



435

Ga (Ga,In)(As,P) Si 

photoluminescence, and X-ray diffraction methods. It was found that there was marked 

interdiffusion of Ga. In a multiple quantum well, Zn diffusion at 500C caused Ga 

intermixing and the Auger electron spectroscopic profiles became flat. At higher 

temperatures, an ordering was found such that the Ga concentration was greatest in the 

original InP layers. This was attributed to the effect of minimization of the free energy; 

balanced by an increase in the mismatch strain energy. Photoluminescence data revealed 

that interdiffusion began at temperatures above 420C. 


Physics, 1990, 67[6], 2919-26 

[446-74-040] 

In 

InGaAsP/InP: In Diffusion 



photoluminescence, and X-ray diffraction methods. It was found that there was marked 

interdiffusion of In. In a multiple quantum well, Zn diffusion at 500C caused In 

intermixing and the Auger electron spectroscopic profiles became flat. 


Physics, 1990, 67[6], 2919-26 

[446-74-040] 

P 

InGaAsP/InP: P Diffusion 



photoluminescence, and X-ray diffraction methods. It was found that the interdiffusion of 

P was negligible. 


Physics, 1990, 67[6], 2919-26 

[446-74-040] 

Si 

GaInAsP/GaAs: Si Diffusion 

The effect of concurrent Si and Zn diffusion, upon interdiffusion between the cation and 

anion sub-lattices, was studied in Ga 0.95 In 0.05 As 0.95 P 0.05 /GaAs heterostructures which had 

been grown by using liquid-phase epitaxy techniques. The diffusion sources were 

equilibrium ternary tie-triangle compositions. The extent of interdiffusion of both group- 

III and group-V atoms was determined by depth profiling In and P, respectively, using 

436

Si (Ga,In)(As,P) Zn 

secondary ion mass spectrometry. It was found that Si diffusion enhanced both cation and 

anion interdiffusion to the same extent. A mono-vacancy mechanism was used to explain 

the effect of Si. It was concluded that the impurity diffusion mechanism was a major 

factor which affected the degree of enhancement. 

K.H.Lee, H.H.Park, D.A.Stevenson: Journal of Applied Physics, 1989, 65[3], 1048-52 

[446-72/73-038] 

InGaAsP/GaAs: Si Diffusion 





column-V atoms in various Si-diffused heterostructures which were closely latticematched 

to GaAs. An enhancement of lattice atom self-diffusion, due to impurity 







[446-64/65-157] 

Zn 

InGaAsP: Zn Diffusion 

The ampoule diffusion of Zn into liquid-phase epitaxial layers was studied, at 

temperatures of between 425 and 525C, by means of secondary ion mass spectrometry. It 

was found that the incorporation and diffusion of Zn could be described in terms of the 

interstitial-substitutional model. The difference between the acceptor and Zn 

concentrations was attributed to compensation by Zn interstitial donors or by neutral Znvacancy 

complexes. The diffusion depth was slightly smaller than that in InP and was 

much larger than that in GaAs. In the case of n-type InGaAsP, the profiles exhibited a 

cut-off which was like that seen in InP. 

G.J.Van Gurp, D.L.A.Tjaden, G.M.Fontijn, P.R.Boudewijn: Journal of Applied Physics, 

1988, 64[7], 3468-71 

[446-72/73-037] 


Closed-ampoule diffusion led to a net acceptor concentration which was lower than the 

Zn concentration. Upon annealing in an atmosphere without Zn, the Zn and net acceptor 

concentrations became almost identical. This was attributed to a decreased Zn 

concentration and an increased net acceptor concentration. The results were quantitatively 

explained by assuming that the Zn was incorporated as both substitutional acceptors and 

437

Zn (Ga,In)(As,P) Zn 

interstitial donors, and that only the interstitial Zn was driven out by annealing; due to its 

large diffusion coefficient. The profiles which were calculated by using this interstitialsubstitutional 

model could be fitted to experimentally determined profiles by assuming 

that the Zn diffused as singly-ionized interstitial donors. The present model also 

explained published data, on diffusion in n-type InP, in which a profile cut-off was found 

at a depth where the acceptor concentration equalled the background donor concentration. 

G.J.Van Gurp, T.Van Dongen, G.M.Fontijn, J.M.Jacobs, D.L.A.Tjaden: Journal of 


[446-72/73-037] 


Shallow diffusion was improved by using a new spun-on source which was based upon 

Zn-doped alumina. In this case, the thermal expansion coefficients of the diffusion source 

and the semiconductor were better matched than when using Zn-doped silica films. As 

well as excellent mechanical stability of the spun-on films over a wide temperature range, 

the effect of mechanical stresses upon the diffusion process was effectively reduced. 

M.C.Amann, G.Franz: Journal of Applied Physics, 1987, 62[4], 1541-3 

[446-55/56-029] 

GaInAsP/GaAs: Zn Diffusion 

The effect of concurrent Zn diffusion upon interdiffusion in an Ga 0.94 In 0.06 As 0.95 P 0.05 - 

GaAs heterostructure, grown by means of liquid phase epitaxy, was investigated. 

Diffusion annealing (700C, 25h) was performed by using an equilibrium ternary diffusion 

source, and In and P profiles were measured by using secondary ion mass spectrometry. It 

was found that Zn diffusion selectively enhanced cation (In-Ga) interdiffusion. In the case 

of concurrent Zn diffusion, the interdiffusion coefficient for the In-Ga components was 

about 5 x 10 -14 cm 2 /s; as compared with about 6 x 10 -16 cm 2 /s for anions (As-P). A kick-out 

mechanism was suggested to explain the results. 

H.H.Park, K.H.Lee, D.A.Stevenson: Applied Physics Letters, 1988, 53[23], 2299-301 

[446-64/65-174] 

GaInAsP/GaAs: Zn Diffusion 

The effect of concurrent Si and Zn diffusion, upon interdiffusion between the cation and 

anion sub-lattices, was studied in Ga 0.95 In 0.05 As 0.95 P 0.05 /GaAs heterostructures which had 

been grown by using liquid-phase epitaxy techniques. The diffusion sources were 

equilibrium ternary tie-triangle compositions. The extent of interdiffusion of both group- 

III and group-V atoms was determined by depth profiling In and P, respectively, using 

secondary ion mass spectrometry. It was found that Zn diffusion selectively enhanced 

cation (In, Ga) interdiffusion. A kick-out mechanism was used to explain the selective 

enhancement of cation interdiffusion by Zn. It was concluded that the impurity diffusion 

mechanism was a major factor which affected the degree of enhancement. 

K.H.Lee, H.H.Park, D.A.Stevenson: Journal of Applied Physics, 1989, 65[3], 1048-52 

[446-72/73-038] 

438


GaInAsP/InP: Zn Diffusion 

It was shown that Zn diffusion into a Ga x In 1-x As y P 1-y -InP quantum well structure and 

superlattice, with a thickness of 10nm, completely disordered the quantum well and the 

superlattice layers. It was found that the photoluminescence wavelength of the quantum 

well and the superlattice had increased after Zn diffusion. This was attributed to In-Ga 

interdiffusion at the interface, due to an interchange process between interstitials and 

substitutional Zn atoms. 

M.Razeghi, O.Acher, F.Launay: Semiconductor Science and Technology, 1987, 2[12], 

793-6 

[446-61-077] 

InGaAsP/InP: Zn Diffusion 

The effect of Zn diffusion upon quantum-well structures was studied by using secondary 

ion mass spectrometry, Auger electron spectroscopy, capacitance-voltage measurements, 

photoluminescence, and X-ray diffraction methods. The structures were stable to 

annealing, in the absence of Zn. In a multiple quantum well, Zn diffusion at 500C caused 

In and Ga intermixing and the Auger electron spectroscopic profiles became flat. 

Photoluminescence data revealed that interdiffusion began at temperatures above 420C. 

The diffusion of Zn changed the lattice parameter values of InGaAsP and InP so that the 

average value decreased. 


Physics, 1990, 67[6], 2919-26 

[446-74-040] 


The effect of concurrent Zn diffusion upon interdiffusion in a In 0.72 Ga 0.28 As 0.61 P 0.39 /InP 

heterostructure was investigated by using Auger electron spectroscopy and secondary ion 

mass spectrometry. The measured profiles showed that Zn diffusion (600C, 1 to 4h) 

mainly enhanced cation (In, Ga) interdiffusion. This could not be explained in terms of 

the Zn-vacancy complex model. Results which were obtained under conditions of group- 

V element over-pressure suggested that cation interstitials could control both the rate of 

Zn diffusion and the mixing of group-III sub-lattices in the InP-based alloy system. 

H.H.Park, B.K.Kang, E.S.Nam, Y.T.Lee, J.H.Kim, O.Kwon: Applied Physics Letters, 

1989, 55[17], 1768-70 

[446-72/73-029] 


The degradation of lattice-matched In 0.72 Ga 0.28 As 0.61 P 0.39 /InP hetero-interfaces during Zn 

diffusion was investigated by using high-resolution transmission electron microscopy and 

Auger electron spectroscopy. Diffusion-induced intermixing of In and Ga across the 

GaInAsP/InP interface produced tensile stresses in the Ga-mixed InP side and 

compressive stresses in the In-mixed GaInAsP side. The effects of localized interfacial 

stresses upon the nucleation of misfit dislocations and upon their strain accommodation 

439


behavior were clearly revealed throughout the intermixed region, and reached several 

thousand Å on both sides of the interface. The interfacial strain was relaxed by the 

generation of paired dislocations, with anti-parallel Burgers vectors, which arose from the 

intermixed GaInAsP/GaInP interface. The dislocation morphologies revealed striking 

contrasts across the intermixed interface; involving stacking faults in the tensile layer and 

perfect dislocation tangles in the compressive layer. The dislocation lines were 

concentrated at the GaInAsP/GaInP interface and along misfit boundaries in the frontal 

areas of the intermixed region. A model was proposed, in order to explain the strain 

relaxation behavior in the intermixed region, which invoked an homogeneous nucleation 

mechanism and splitting of the paired dislocations from the intermixed interface. 

H.H.Park, K.H.Lee, J.K.Lee, Y.T.Lee, E.H.Lee, J.Y.Lee, S.K.Hong, O.Kwon: Journal of 


[446-106/107-118] 


An open-tube Zn diffusion method that involved the use of solid Zn 3 P 2 or InP was 

developed, and was used to fabricate heterojunction devices. The diffusion profiles were 

determined by means of secondary ion mass spectrometry. It was found that the diffusion 

depth was proportional to the square root of the diffusion time, and that a sharp change in 

the Zn concentration was observed at the diffusion front. The activation energies were 

equal to 1.2eV for undoped InP, 1.7eV for Sn-doped InP, and 1.1eV for InGaAs. 

T.Ohishi, K.Ohtsuka, Y.Abe, H.Sugimoto, T.Matsui: Japanese Journal of Applied 

Physics, 1990, 29[2], L213-6 

[446-76/77-024] 


The diffusion of Zn into heterostructures, from spun-on polymer films which contained 

no Si-organic compounds, was investigated. It was shown that the initial non-equilibrium 

stage of diffusion, and the defect relaxation process, determined the final profile of the Zn 

distribution; as well as its anomalies. The results demonstrated the important role which 

was played by the kick-out mechanism in the formation of a sharp diffusion profile slope 

as well as in Zn gettering at a hetero-interface. 

B.J.Ber, L.A.Busygina, A.T.Gorelenok, A.V.Kamanin, A.V.Merkulov, I.A.Mokina, 

N.M.Shmidt, I.J.Yakimenko, T.A.Yurre: Materials Science Forum, 1994, 143-147, 1415- 

20 

[446-113/114-037] 


A study was made of dopant redistribution between a chemical beam epitaxially grown 

InGaAsP laser structure, and a metalorganic vapor phase epitaxially over-grown InP 

layer. Secondary ion mass spectroscopic data revealed that Zn and Be atoms interdiffused 

deeply into the adjacent InP layers, at a Zn doping level of 10 18 /cm 3 . A fraction of the Zn 

atoms went through the chemical beam epitaxial InP, and penetrated the laser structure 

440

Zn (Ga,In)(As,P) Interdiffusion 

guide layer. It was found that Zn out-diffusion was appreciably suppressed by reducing 

the Be dopant concentration from 10 18 to 5 x 10 17 /cm 3 . 

H.Sugiura, S.Kondo, M.Mitsuhara, S.Matsumoto, M.Itoh: Applied Physics Letters, 1997, 

70[21], 2846-8 

[446-152-0441] 


InGaAsP/InGaAsP: Interdiffusion 

The interdiffusion behavior of In 1-x Ga x As 1-y P y /In 1-x' Ga x' As 1-y' P y' multiple quantum-well 

heterostructures, with various built-in strains and layer thicknesses, was investigated by 

monitoring their photoluminescence. All of the samples which were annealed at 850C 

consistently exhibited a blue-shift of the heavy hole exciton line as a result of 

interdiffusion across the hetero-interfaces. Data on a nearly strain-compensated structure, 

with a constant P/As ratio and In-rich wells, showed that the blue-shift of the excitonic 

line was the result of group-III atom interdiffusion alone, as in the GaAs/GaAlAs system. 

The In-Ga interdiffusion could be described by a coefficient of about 4.72 x 10 -16 cm 2 /s. 

In the case of lattice-matched and compressively strained structures, simultaneous 

interdiffusion on the group-III and group-V sub-lattices yielded an effective interdiffusion 

coefficient which ranged from 3.83 x 10 -16 to 5.51 x 10 -16 cm 2 /s. 

A.Hamoudi, A.Ougazzaden, P.Krauz, K.Rao, M.Juhel, H.Thibierge: Japanese Journal of 


[446-119/120-215] 

InGaAsP/InGaAsP: Interdiffusion 

By using quaternary/quaternary multiple quantum wells, with constant P/As ratios and Inrich 

wells, it was shown that it was possible to produce a blue-shift of the heavy hole 

exciton line by interdiffusing only group-III atoms. The use of this type of structure, 

which was particularly suitable for group-III diffusion studies, yielded a value (for the In- 

Ga interdiffusion coefficient) of about 4.72 x 10 -16 cm 2 /s at 850C. 

A.Hamoudi, A.Ougazzaden, P.Krauz, E.V.K.Rao, M.Juhel, H.Thibierge: Applied Physics 

Letters, 1995, 66[6], 718-20 

[446-121/122-078] 

441

(Ga,In)P 

Zn 

InGaP: Zn Diffusion 

The migration of Zn was studied by using angle-lapping and stain-etching techniques. It 

was found that the Zn diffusion depth was proportional to the square root of the diffusion 

time, and depended upon the composition of In 1-x Ga x P. That is, the effective diffusion 

coefficient at a given diffusion temperature decreased according to: 

D(cm 2 /s) = 3.935 x 10 -8 exp[-6.84x] 

while the activation energy increased according to: 

Q(eV) = 1.28 + 2.38x 

This was attributed to a decreasing contribution from interstitial diffusion with increasing 

Ga content. 

S.T.Kim, D.C.Moon: Japanese Journal of Applied Physics, 1990, 29[4], 627-9 

[446-76/77-025] 

InGaP: Zn Diffusion 

The diffusion of Zn was studied, using photoluminescence techniques, in liquid-phase 

epitaxial In 0.5 Ga 0.5 P layers which had been grown onto semi-insulating GaAs substrates. 

The photoluminescence exhibited a characteristic emission peak at 1.934eV after 

diffusion. It was found that this peak behaved like donor-acceptor pair recombination and 

was associated with the interstitial Zn donor to substitutional Zn acceptor band transition. 

The calculated activation energy of the substitutional acceptor was 0.047eV. 

I.T.Yoon, B.S.Jeong, H.L.Park: Thin Solid Films, 1997, 300[1-2], 284-8 

[446-157/159-442] 

GaInP/GaAs: Zn Diffusion 

Heterojunction Zn-doped bipolar transistors were prepared by means of metal-organic 

chemical vapor deposition. The characteristics of heterojunction bipolar transistors with 

an undoped spacer layer and an n-type GaAs set-back layer (heterostructure-emitter) 

442

Zn (Ga,In)P|(Ga,In)Sb Interdiffusion 

between base and emitter were investigated. The results showed that base dopant outdiffusion 

was effectively prevented in heterostructure-emitter bipolar transistors. 

Y.F.Yang, C.C.Hsu, E.S.Yang: Semiconductor Science and Technology, 1995, 10[3], 

339-43 

[446-121/122-069] 

(Ga,In)Sb 


GaInSb/InAs: Interdiffusion 

Strained superlattices which exhibited a high degree of structural perfection were grown 

onto GaSb substrates. The superlattices exhibited ideal defect-free structures. Crosssectional 

micrographs revealed that the layers were highly planar, regular and coherently 

strained with respect to the substrate. No crystalline defects were observed by using 

transmission electron microscopy, in spite of the existence of a lattice mismatch of almost 

2%. The planarity of the layers was confirmed by the observation of Pendellösung fringes 

in high-resolution X-ray diffraction patterns, while the observation of numerous sharp 

satellite peaks indicated that essentially no interdiffusion occurred within the 

superlattices. 

R.H.Miles, D.H.Chow, W.J.Hamilton: Journal of Applied Physics, 1992, 71[1], 211-4 

[446-86/87-033] 

443

GaN 

D 

GaN: D Diffusion 



Secondary ion mass spectrometry was used to measure the resultant distributions. In the 

case of plasma-treated samples, 2 populations of 2 H were found. At concentrations that 

were greater than 10 20 /cm 3 , there was a near-surface (less than 0.3µ) region that was 

probably due to the formation of platelet defects. At concentrations of about 10 18 /cm 3 , a 

plateau region was present which extended throughout the film thickness of about 1µ. 

This was attributed to the pairing of 2 H with point defects. The D in the former region 

began to out-diffuse at 300C. In the latter region, out-diffusion did not begin until the 

temperature exceeded 800C. In implanted samples, 2 H redistribution occurred in the same 

manner as the bulk population in plasma-treated material. The thermal stability of the D 

profiles in the nitride was much higher than that in GaAs and similar compounds. 


Technology A, 1995, 13[3], 719-23 

[446-140-029] 

H 

GaN: H Diffusion 

An investigation was made of interactions between H and dopant impurities, in this 

material, by using first-principles calculations. The results revealed a fundamental 

difference in the behaviors of H in p-type and n-type samples. In particular, it was 

explained why the H concentrations in n-type material were low, and why H had a 

beneficial effect upon acceptor incorporation in p-type material. Overall, the conclusions 

supported the potential use of H as a means of improving the doping of wide-bandgap 

semiconductors, although the approach was not generally applicable. In order to be able to 

exploit H passivation as a method for enhancing doping, the H had to be the predominant 

compensating defect. That is, its formation energy had to be lower than that of all native 

444

H GaN Surface 

defects, and be comparable to the formation energy of the dopant impurity. Also, the 

activation energy which was required to dissociate the H-impurity complex, and to 

remove or neutralize H, had to be lower than the activation energies for native defect 

formation and lower than the diffusion barrier of the impurity. Finally, the dissociated H 

atom had to be highly mobile. 

J.Neugebauer, C.G.Van de Walle: Applied Physics Letters, 1996, 68[13], 1829-31 

[446-134/135-263] 

GaN: H Diffusion 

A study was made, of the electronic structure, energetics, and migration of H and H- 

complexes, on the basis of first-principles total-energy calculations. The latter revealed a 

number of features which were very different to those which were exhibited by H in more 

traditional semiconductors such as Si or GaAs. These included a very large negative-U 

effect (of about 2.4eV), an instability of the bond-center site, high energies of H 

molecules, and an unusual geometry of the Mg-H complex. All of these features were 

shown to be a result of the particular properties of the present material, such as its 

strongly ionic nature and the high strength of the Ga-N bond. A simple model was 

proposed, for the negative-U behavior, which was expected to be valid for H in any 

semiconductor. 

J.Neugebauer, C.G.Van de Walle: Physical Review Letters, 1995, 75[24], 4452-5 

[446-127/128-227] 


GaN/GaAs: Ga Surface Diffusion 

Metalorganic vapor phase epitaxial growth of cubic GaN on patterned GaAs(100) 

substrates which had (111)A or (111)B facets was studied. It was found that the growth 

features depended strongly upon the configuration of the pattern. It was deduced that 

these features arose from an orientation-dependent growth rate, which varied in the order: 

(111)B > (100) > (111)A, together with Ga adatom diffusion on the surface. By taking 

advantage of the growth on patterned substrates, the diffusion length of Ga adatoms on 

the GaAs(100) surface was estimated to be typically equal to several microns. 

M.Nagahara, S.Miyoshi, H.Yaguchi, K.Onabe, Y.Shiraki, R.Ito: Journal of Crystal 

Growth, 1994, 145[1-4], 197-202 

[446-119/120-306] 


GaN: Surface Diffusion 

The growth kinetics of GaN/(00?1)sapphire hetero-epitaxial films were studied, at 

substrate temperatures ranging from 560 to 640C, by using reflection high-energy 

electron diffraction specular reflection intensity monitoring techniques. An alternatingelement 

exposure growth method was used in which Ga and N atoms were supplied 

445

Surface GaN|GaP Ga 

separately (rather than simultaneously as usual) to the substrate, with the insertion of a 

time delay between successive Ga-flux and N-flux exposures. A time-dependent recovery 

of the reflection high-energy electron diffraction specular reflection intensity, during the 

time delay, was associated with Ga-N surface molecule migration on Ga-terminated 

surfaces. The activation energy for this migration process was deduced to be 1.45eV. 

H.Liu, J.G.Kim, M.H.Ludwig, R.M.Park: Applied Physics Letters, 1997, 71[3], 347-9 

[446-152-0446] 

GaP 

Fe 

GaP: Fe Diffusion 

The structural properties of samples which had been implanted with 150 or 400keV Fe or 

Ti, to doses of between 10 12 and 10 15 /cm 2 , were studied. The depth distributions of the 

implants were compared before and after annealing with, or without, a Si 3 N 4 cap. 

Rutherford back-scattering, X-ray double-crystal diffractometry, and secondary ion mass 

spectroscopy results indicated that Fe was markedly redistributed in all of the materials 

during annealing. On the other hand, Ti did not redistribute at all. The driving force for 

the redistribution of Fe was thought to be not classical diffusion, but reaction with 

implantation-induced defects and stoichiometric imbalances. The defect chemistry of asimplanted 

arsenides was found to be fundamentally different to that of as-implanted 




1992, 72[8], 3514-21 

[446-106/107-034] 

Ga 

GaP: Ga Diffusion 

The self-diffusion of Ga was measured directly in isotopically controlled heterostructures. 

Secondary ion mass spectroscopy was used to monitor the intermixing of 69 Ga and 71 Ga 

446

Ga GaP General 

between isotopically pure GaP epilayers which had been grown, by molecular beam 

epitaxy, onto GaP substrates. It was found that Ga self-diffusivity in undoped GaP could 


D (cm 2 /s) = 2.0 exp[-4.5(eV)/kT] 

at temperatures of between 1000 and 1190C, under P-rich condition. The entropy of selfdiffusion 

was estimated to be about 4k. 

L.Wang, J.A.Wolk, L.Hsu, E.E.Haller, J.W.Erickson, M.Cardona, T.Ruf, J.P.Silveira, 

F.Briones: Applied Physics Letters, 1997, 70[14], 1831-3 

[446-150/151-144] 

H 

GaP: H Diffusion 

The H passivation effect of Zn was studied by monitoring donor-acceptor pair 

luminescence bands. The existence of a critical pair-separation distance was detected, 

within which neutralization by H was almost entirely suppressed by the presence of a 

nearby donor. This phenomenon was reflected by a relative enhancement of the 

intensities of the discrete pair lines, as compared to that of the remote pair band in S-Zn 

pair emission. Further evidence was provided by the fact that the intensity reduction in 

the O-Zn band was much smaller than that predicted by calculations which assumed a 

random neutralization of Zn acceptors. It was suggested that the suppression was due to 

the Coulomb potential of a nearby donor atom, and was consistent with the concept that 

the principal diffusing form of H was positively charged. 

Y.Mochizuki, M.Mizuta: Materials Science Forum, 1992, 83-87, 575-80 

[446-99/100-085] 

Ti 

GaP: Ti Diffusion 







1992, 72[8], 3514-21 

[446-106/107-034] 

General 

GaP/InP: Diffusion 

A theoretical study was made of the defect characteristics of a strained-layer superlattice. 

On the basis of the calculated charge states, some closely-bound interstitial-antisite pairs 

447

General GaP Interdiffusion 

were found to form at a distance of one bond length in a stable GaP/InP (1 x 1) structure. 

The defect-related states that were localized in the energy gap suggested that the 

formation mechanism of the EL2 complex in the GaAs system was inapplicable to 

GaP/InP multilayer structures. A slow self-diffusion process was proposed for both 

group-III and group-V atoms in this superlattice. Kick-out reactions were suggested to be 

favored at higher dopant concentrations. 

E.G.Wang: Physica Status Solidi B, 1992, 174[2], 367-74 

[446-106/107-096] 


GaP: Surface Diffusion 

A scanning tunnelling microscope tip-induced migration of di-vacancies was observed on 

the (110) surface of n-doped material. Such motion occurred only for a negative polarity 

of the tunnelling voltage, and was directed along the [1¯10] direction; regardless of the 

scanning orientation. It was therefore purely non-thermal. It was suggested that a fieldinduced 

reduction of the barrier to migration initiated the motion. A quantitative analysis 

of the migration indicated the existence of long-lived excited states. This was consistent 

with charged trap states in doped semiconductors. 

P.Ebert, M.G.Lagally, K.Urban: Physical Review Letters, 1993, 70[10], 1437-40 

[446-106/107-093] 


GaP/InP: Interdiffusion 

The interdiffusion of lateral composition-modulated (GaP) 2 /(InP) 2 short-period 

superlattices was studied. Lateral compositional modulation was obtained by using the 

strain-induced lateral layer ordering process. A blue-shift of the inter-band transition was 

observed, by means of photoluminescence spectroscopy, in cap-less and SiO 2 - 

encapsulated annealed (800C, 5.5h) short-period superlattices. The intensity and 

wavelength of Si 3 N 4 -encapsulated annealed short-period superlattices were only slightly 

perturbed. Transmission electron microscopy showed that capless-annealed (800C, 5.5h) 

short-period superlattices retained their lateral composition modulation. However, the 

(00½) satellite reflections disappeared. After long (48h) annealing, the inter-band 

transition corresponded to that of an In 0.50 Ga 0.50 P alloy. This suggested that the lateral 

composition modulation disappeared. The observed lateral interdiffusion coefficient 

exceeded the vertical one by a factor of about 30; thus suggesting that short-period 

superlattice interdiffusion was enhanced by native point defects. 

J.I.Malin, A.C.Chen, J.E.Bonkowski, J.E.Baker, K.Y.Cheng, K.C.Hsieh: Journal of 


[446-136/137-120] 

448

GaSb 

H 

GaSb: H Diffusion 

The use of spreading resistance and capacitance-voltage methods showed that atomic H 

passivated shallow acceptors and donors in this material. Deep-level passivation by H 

also occurred, as revealed by deep-level transient spectroscopic measurements of 

Schottky diode structures. Effective H diffusion coefficients were deduced for both n + - 

type and p + -type samples. In the former case, the diffusion was thermally activated and 

could be described by: 

D(cm 2 /s) = 3.4 x 10 -5 exp[-0.55(eV)/kT] 

at temperatures of between 100 and 250C. In the latter case, the diffusivity could be 

described by: 

D(cm 2 /s) = 1.5 x 10 -6 exp[-0.45(eV)/kT] 

Reactivation of passivated shallow and deep levels occurred at temperatures of between 

250 and 300C. 

A.Y.Polyakov, S.J.Pearton, R.G.Wilson, P.Rai-Choudhury, R.J.Hillard, X.J.Bao, 

M.Stam, A.G.Milnes, T.E.Schlesinger, J.Lopata: Applied Physics Letters, 1992, 60[11], 

1318-20 

[446-86/87-036] 

Sb 

GaSb: Sb Diffusion 

Recovery in crystals which had been implanted with Ga ions, and then subjected to rapid 

thermal annealing or furnace annealing, was studied by using Raman scattering 

techniques. The intensity of the LO phonon mode decreased with increasing ion 

implantation fluence. It was found that the threshold fluence for the amorphization of Gaimplanted 

GaSb was 5 x 10 13 /cm 2 . This value was much lower than that for InP 

(10 14 /cm 2 ). In the case of furnace annealing, the recovery processes in Ga-implanted 

material were very different above and below a fluence of 5 x 10 14 /cm 2 . No recovery was 

observed at this critical fluence. At lower fluences, the LO mode intensity increased with 

449

Sb GaSb Zn 

increasing annealing temperatures and times of up to 400C and of up to 0.25h. However, 

the damage recovery was very poor when compared with that of GaAs, InP, or GaP. In 

the case of Si 3 N 4 -capped rapid thermal annealing, recovery was observed even for a 

fluence of 5 x 10 14 /cm 2 . New modes were observed, at about 114 and 150/cm, in 

implanted and annealed samples. These 2 modes were related to the E g and A lg modes of 

Sb-Sb bond vibrations, respectively, and were attributed to the out-diffusion of Sb atoms. 

It was found that furnace annealing enhanced Sb out-diffusion, and that the capped rapid 

thermal annealing process was superior to the furnace annealing process for the healing of 

damaged layers. These anomalous behaviors were thought to be closely related to the 

weak bond-strength of Sb-containing materials. The degree of recovery, as a function of 

annealing temperature, annealing time, and fluence was also investigated. 

S.G.Kim, H.Asahi, M.Seta, J.Takizawa, S.Emura, R.K.Soni, S.Gonda, H.Tanoue: Journal 


[446-106/107-096] 

Te 

GaSb: Te Diffusion 

The diffusion of Te into undoped p-type material was carried out, and the samples were 

studied using cathodoluminescence and photoluminescence techniques. The luminescence 

centers in Te-diffused samples were identified and were compared with those in Tedoped 

bulk material. Essential differences in the radiative levels, between diffused and 

as-grown doped samples, were observed. Evidence of the presence of self-compensating 

acceptor complexes was found in the case of diffused samples. At short and medium 

diffusion times, a compensating acceptor complex, V Ga Ga Sb Te Sb , was observed. At long 

diffusion times, the predominant acceptor center was suggested to be the antisite defect, 

Ga Sb , or a related complex. 

P.S.Dutta, B.Méndez, J.Piqueras, E.Dieguez, H.L.Bhat: Journal of Applied Physics, 

1996, 80[2], 1112-5 

[446-136/137-120] 

Zn 

GaSb: Zn Diffusion 

It was recalled that Zn diffusion was usually performed in an evacuated quartz ampoule, 

with one source for the dopant and with a second antimonide source being used to 

provide an over-pressure over the GaSb during diffusion. Strict control of the system 

partial pressure was required in order to produce reproducible diffusion and damage-free 

surfaces. An alternative technique was open-tube diffusion from doped spun-on films. 

Such dopant emulsions had been successfully applied to some III/V semiconductors. The 

present work was the first report of the reproducible diffusion of Zn into GaSb from a 

spun-on diffusion source in an open-tube system. A Zn/SiO 2 film was used as the 

450

Zn GaSb Zn 

diffusion source at temperatures of 630 or 680C. The concentration in the semiconductor 

could fall below the substrate dopant level due to the depletion effects of the source. 

C.Heinz: Journal of the Electrochemical Society, 1988, 135[1], 250-2 

[446-61-078] 


The diffusion of Zn into Te-doped samples was studied as a function of time, temperature 

and Sb over-pressure. The overall Zn profiles, as well as carrier concentration profiles, 

were determined. The results indicated an inverse dependence of the diffusivity upon the 

Sb over-pressure, and reflected the operation of an interstitial-substitutional vacancy 

mechanism. At high Zn concentrations, the profiles indicated the existence of an 

additional component which was associated with a non-electrically active Zn species that 

had a small, but highly temperature-dependent, diffusion coefficient. 

G.J.Conibeer, A.F.W.Willoughby, C.M.Hardingham, V.K.M.Sharma: Optical Materials, 

1996, 6[1-2], 21-5 

[446-141/142-104] 


The diffusion of Zn into Te-doped samples was studied as a function of time, 

temperature, and Sb over-pressure. The overall Zn profiles, as well as carrier 

concentration profiles, were measured. The results indicated the operation of a 

substitutional-interstitial vacancy (Frank-Turnbull) or kick-out (Gosele-Morehead) 

mechanism; although there was insufficient evidence to decide between them. There was 

also an inverse dependence of the diffusivity upon the Sb over-pressure. This was 

explained in terms of a Zn diffusion that was superposed on Ga vacancy diffusion. It was 

noted that Te doping appeared to have little effect upon diffusion, due to its low level 

when compared to that of Zn. Moreover, at high Zn concentrations, the profiles indicated 

the presence of an additional component that was associated with a non-electrically active 

Zn species which had a small strongly temperature-dependent diffusion coefficient. 

G.J.Conibeer, A.F.W.Willoughby, C.M.Hardingham, V.K.M.Sharma: Journal of 


[446-141/142-105] 


Spin-on diffusion of Zn into samples in an open tube-furnace was investigated by using a 

Zn-doped silica film as a diffusion source. Theoretical calculations were carried out which 

assumed the silica film to be an exhaustible source. Diffusion treatments were performed 

at 630 or 680C, using undiluted or diluted Zn-containing silica films. It was found that the 

junction depth exhibited a linear dependence upon the square root of the diffusion time for 

undiluted films with no source depletion, whereas results which were obtained by diffusion 

from diluted films clearly reflected depletion effects in an exhaustible diffusion 

451

Zn GaSb General 

source. The method was also applicable to Ga 0.935 Al 0.065 Sb, due to its very small Al 

content. 

C.Heinz: Solid-State Electronics, 1993, 36[12], 1685-8 

[446-115/116-134] 


A closed-ampoule technique was used to introduce Zn into Te-doped material. Annealing 

was carried out for various times; with or without an Sb over-pressure. The total Zn 

profiles were measured by using secondary ion mass spectrometry, and carrier profiles 

were deduced from incremental sheet resistance data. It was found that the diffusivity 

varied, with Zn concentration, from 6.3 x 10 -12 cm 2 /s for a concentration of 3 x 10 20 /cm 3 

at 500C. 

G.J.Conibeer, A.F.W.Willoughby, C.M.Hardingham, V.K.M.Sharma: Materials Science 

Forum, 1994, 143-147, 1427-32 

[446-113/114-031] 


The diffusion of Zn in n-type material, at temperatures ranging from 450 to 540C, was 

studied. The diffusion was carried out in a closed system, using a Zn-Ga source. The 

diffusion profiles were measured by using secondary ion mass spectrometry. On the basis 

of the diffusion profiles, concentration-dependent diffusion coefficients were calculated 

by using Boltzmann-Matano analyses. These data were qualitatively interpreted in terms 

of an interstitial-substitutional diffusion model which had originally been proposed for Zn 

diffusion in GaAs. 

V.S.Sundaram, P.E.Gruenbaum: Journal of Applied Physics, 1993, 73[8], 3787-9 

[446-109/110-036] 

General 

GaSb: Diffusion 

An extensive review was presented of advances in the use of GaSb-based systems. It 

detailed all aspects: from bulk crystal growth or epitaxy, post-growth processing to device 

design. Some current areas of research and development were critically assessed, and 

their significance with respect to the understanding of basic physical phenomena and 

practical applicability was addressed. These areas included the role played by defects and 

impurities in affecting the structural and other properties of the material, and the 

techniques which were used for surface- and bulk-defect passivation. It was concluded 

that current knowledge concerning this material was sufficient to explain its basic 

properties and permit its further application. Among the topics which were listed were: 

defects and impurities, extended defects, native defects, isotopic effects, self- and 

impurity diffusion, ion bombardment-induced defects, surface and bulk defect 

452

General GaSb|InAs H 

passivation, H-plasma passivation, and the passivation of amorphous hydrogenated 

material. 

P.S.Dutta, H.L.Bhat, V.Kumar: Journal of Applied Physics, 1997, 81[9], 5821-70 

[446-150/151-145] 

InAs 

Fe 

InAs: Fe Diffusion 













1992, 72[8], 3514-21 

[446-106/107-034] 

H 

InAs: H Diffusion 

Accidentally doped InAs layers were grown, by molecular beam epitaxy, onto semiinsulating 

GaAs substrates, and were exposed to H plasma. It was found that the 

453

H InAs Zn 

diffusivity of H in these layers was high. Hydrogenation led to an order of magnitude 

increase in the free carrier density. At the same time, defect-related peaks disappeared 

from the near-bandedge luminescence spectrum. On the other hand, the properties of bulk 

InAs or those of InAs-on-InAs layers were not altered by exposure to the H plasma. A 

model which was based upon the passivation, by H, of electronic states which were 

induced by dislocations in InAs-on-GaAs layers was proposed in order to explain these 

effects. 

B.Theys, S.Kalem, A.Lusson, J.Chevallier, C.Grattepain, M.Stutzmann: Materials 

Science Forum, 1992, 83-87, 629-34 

[446-99/100-090] 

InAs/GaAs: H Diffusion 

Atomic H was introduced, into molecular beam epitaxial InAs layers on GaAs substrates, 

from a plasma source. It was found that the H diffused very rapidly into the material. Its 

presence modified the electronic transport properties, the near-bandedge luminescence 

spectra, and the far-infrared reflectivity. The free-carrier density had increased by an 

order of magnitude after hydrogenation. These effects could be removed by thermal 


B.Theys, A.Lusson, J.Chevallier, C.Grattepain, S.Kalem, M.Stutzmann: Journal of 


[446-91/92-023] 

Ti 

InAs: Ti Diffusion 







1992, 72[8], 3514-21 

[446-106/107-034] 

Zn 

InAs: Zn Diffusion 

Elemental Zn was diffused into (100)-oriented samples by using the closed-tube method. 

The diffusion temperatures ranged from 350 to 500C. It was found that the results could 


D(cm 2 /s) = 0.00016 exp[-1.07(eV)/kT] 

454

Zn InAs Surface 

H.Khald, H.Mani, A.Joullie: Journal of Applied Physics, 1988, 64[9], 4768-70 

[446-72/73-034] 


In 

InAs: In Surface Diffusion 

The inter-surface diffusion of In adatoms between a (111)A and a (001) surface on an 

InAs (111)A-(001) non-planar substrate was investigated for the first time by using 

microprobe reflection high-energy electron diffraction techniques. It was found that the 

surface diffusion of In adatoms depended strongly upon the growth temperature and the 

As pressure, but was independent of the In flux. It was also observed that the migration 

direction of the In lateral flux between the (111)A and (001) InAs surfaces changed; 

depending upon the growth conditions. 

X.Q.Shen, T.Nishinaga: Journal of Crystal Growth, 1995, 146[1-4], 374-8 

[446-127/128-148] 


InAs: Surface Diffusion 

The sensitivity, of the 2- to 3-dimensional growth transition of InAs self-assembled 

islands, to the InAs coverage was used to demonstrate the growth of self-aligned InAs 

islands on etched GaAs ridges by molecular beam epitaxy. The differing migration 

behavior of In adatoms upon the various crystal planes of etched ridges was used to 

modulate spatially the supply of In adatoms. The ridges were oriented along the [011] and 

[01¯1] directions on (100) substrates with grating spacings of 0.28, 1 or 5µ. Atomic force 

microscopy revealed that the InAs islands were self-aligned along the ridges and were 

typically 40nm in diameter and 12nm in height. In samples with [011]-oriented ridges, the 

islands were located on the side-walls. On the other hand, in the case of [01¯1]-oriented 

ridges, the islands were on the (100) planes on, and at the foot of, the mesa. On samples 

with a grating pitch of 0.28µ, all of the islands were located either on the side-walls or at 

the bottom of the so-called V-groove: for both grating orientations. 

D.S.L.Mui, D.Leonard, L.A.Coldren, P.M.Petroff: Applied Physics Letters, 1995, 66[13], 

1620-2 

[446-121/122-074] 

InAs: Surface Diffusion 

The surface diffusion of group-III atom incorporation during molecular beam epitaxial 

growth was considered. Firstly, the diffusion length for incorporation on the (001) top 

surface, with (111)A or (411)A side surfaces on V grooves, was studied. It was shown 

that the diffusion length took the same value for both cases and was inversely 

455

Surface InAs Interdiffusion 

proportional to the As pressure. The same relationship was also found for the diffusion of 

In in InAs during molecular beam epitaxy. However, the diffusion length of Ga on (111)B 

exhibited an inverse parabolic dependence of the As pressure. It was suggested that, on 

the (001) surface, two As 4 molecules met to furnish active As atoms for growth. On the 

other hand, the behavior of the As 4 molecule on the (111)B surface remained unclear. The 

ratio of the surface diffusion coefficients on (111)B and (001) was calculated. It was 

found that the ratio took a value of about 140. Using this ratio, the incorporation lifetimes 

on (111)B and (001) surfaces were calculated as functions of the As pressure. It was 

found that the curves of incorporation lifetime intersected at the As pressure where flow 

inversion occurred. 

T.Nishinaga, X.Q.Shen, D.Kishimoto: Journal of Crystal Growth, 1996, 163[1], 60-6 

[446-138/139-078] 


InAs/GaAs: Interdiffusion 

The existence of interdiffusion, between self-assembled InAs quantum dots and a GaAs 

substrate, was investigated by using Rutherford back-scattering techniques. These were 

also useful for determining the value of the average InAs layer thickness. As a result, data 

on the diffusion of Ga atoms into the dot were obtained. 

T.Haga, M.Kataoka, N.Matsumura, S.Muto, Y.Nakata, N.Yokoyama: Japanese Journal of 

Applied Physics, 1997, 36[2-8B], L1113-5 

[446-157/159-456] 

InAs/InP: Interdiffusion 

The intermixing of ultra-thin strained quantum-well structures, during annealing at 

temperatures of 730 to 830C, was investigated by means of photoluminescence 

measurements. Upon analyzing the results, using a microscopic model, the interdiffusion 

process was found to be characterized by an activation energy of about 3.8eV. The 

interdiffusion coefficient was close to 7 x 10 -7 cm 2 /s at 830C. 

J.M.Sallese, S.Taylor, H.J.Bühlmann, J.F.Carlin, A.Rudra, R.Houdré, M.Ilegems: 

Applied Physics Letters, 1994, 65[3], 341-3 

[446-125/126-140] 

456

In(As,P) 


InAsP/InP: Interdiffusion 

Evidence was presented for the occurrence of anomalously high strain-dependent 

interdiffusion in InAsP layers which had been grown onto InP(001) substrates by means 

of organometallic vapor phase epitaxy at 620C. In particular, there were clear indications 

of the existence of a critical strain. If the strain was equal to about 1.9% or more, marked 

P-As mixing occurred. At smaller strains, the degree of mixing was greatly reduced. The 

incidence of interdiffusion was also highly sensitive to the temperature. A set of samples 

which was prepared at 580C exhibited an approximately 2-fold decrease in P-As mixing, 

as compared with samples that were prepared at 620C. 

D.J.Tweet, H.Matsuhata, P.Fons, H.Oyanagi, H.Kamei: Applied Physics Letters, 1997, 

70[25], 3410-2 

[446-152-0457] 

In(As,Sb) 

Zn 

InAsSb: Zn Diffusion 

Elemental Zn was diffused into (100)-oriented samples of InAs 1-x Sb x , where x ranged 

from 0.10 to 0.12, by using the closed-tube method. The diffusion temperatures ranged 

from 350 to 500C. It was found that the results could be described by: 

D(cm 2 /s) = 0.00016 exp[-1.07(eV)/kT] 

H.Khald, H.Mani, A.Joullie: Journal of Applied Physics, 1988, 64[9], 4768-70 

[446-72/73-034] 

457

InN 

D 

InN: D Diffusion 



Secondary ion mass spectrometry was used to measure the resultant distributions. At 

concentrations that were greater than 10 20 /cm 3 , there was a near-surface (less than 0.3µ) 

region that was probably due to the formation of platelet defects. At concentrations of 

about 10 18 /cm 3 , a plateau region was present which extended throughout the film 

thickness of about 1µ. This was attributed to the pairing of 2 H with point defects. In 

implanted samples, 2 H redistribution occurred in the same manner as the bulk population 

in plasma-treated material. The thermal stability of the D profiles in the nitride was much 

higher than that in GaAs and similar compounds. 


Technology A, 1995, 13[3], 719-23 

[446-140-029] 

InP 

Au 

InP: Au Diffusion 

Migration of Au at temperatures ranging from 400 to 700C was studied by using 

secondary ion mass spectrometry. Low values were found for the diffusion coefficient; 

458

Au InP Cd 

which was equal to 2 x 10 -12 cm 2 /s at 550C. By using deep-level transient spectroscopy, 

Au was found to behave as a shallow donor with a level that was situated at 0.55eV from 

the conduction band. It was concluded that Au thermal migration from contacts was not 

the mechanism which was responsible for device degradation. 

V.Parguel, P.N.Favennec, M.Gauneau, Y.Rihet, R.Chaplain, H.L'Haridon, C.Vaudry: 


[446-55/56-029] 

Be 

InP: Be Diffusion 

Implantation (1.5 x 10 14 /cm 2 ) of 30keV Be into (x11)A-oriented semi-insulating InP 


(100)-oriented substrates were also implanted. The in-diffusion of Be in (311)A-oriented 

substrates was lower than that in (100) material. 



[446-117/118-166] 

InP: Be Diffusion 

Models were presented for the distribution profiles of Be in ion-doped layers after 

implantation and annealing. The possibility of predicting the mean free path of Be in III- 

V compounds was considered. The effect of defect-impurity interactions upon Be 

diffusion was also examined. It was found that a flux of impurities towards the surface 

occurred which was not diffusive in nature. 

G.I.Koltsov, V.V.Makarov, S.J.Yurchuk: Fizika i Tekhnika Poluprovodnikov, 1996, 

30[10], 1907-16 (Semiconductors, 1996, 30[10], 996-1000) 

[446-148/149-171] 

Cd 

InP: Cd Diffusion 

The so-called leaky tube method was used to diffuse elemental Cd into InP at 500C, and 

the junction depths were determined after various times. It was found that the diffusion 

coefficient was concentration-dependent, and ranged from about 10 -14 to 10 -10 cm 2 /s. 

C.B.Wheeler, R.J.Roedel, R.W.Nelson, S.N.Schauer, P.Williams: Journal of Applied 

Physics, 1990, 68[3], 969-72 

[446-86/87-042] 


The Cd was diffused into InP by using Cd 3 P 2 plus P or Cd 3 P 2 plus Cd 3 As 2 as diffusion 

sources. Two diffusion fronts were observed. The diffusion characteristics of Cd 3 P 2 plus P 

459

Cd InP D 

sources were explained in terms of the interstitial-substitutional model or the vacancy 

complex model. The charge state of the diffusing interstitial Cd atom was a singly ionized 

donor. The chemical species of P which reacted with InP was P 2 , and gaseous Cd 

originated from solid-phase CdP 2 . In the case of Cd 3 P 2 plus Cd 3 As 2 diffusion sources, the 

effective diffusion coefficient and the surface acceptor concentration decreased with 

increasing weight fraction of Cd 3 As 2 . The relative depth of the deeper diffusion front 

increased when the supply of vacancies was suppressed. 

K.I.Ohtsuka, T.Matsui, H.Ogata: Japanese Journal of Applied Physics, 1988, 27[2], 253-9 

[446-60-009] 


Phase diagrams between Cd and III-V compounds were investigated in order to develop 

new acceptor diffusion sources. On the basis of thermodynamic data and experimental 

studies of the Cd-Ga-As and Cd-In-P phase diagrams, it was found that Zn 3 Cd 3 B 2 and 

ZnCdB 2 were suitable acceptor sources in that they did not erode the semiconductor 

surface. 

S.F.Marenkin, O.N.Pashkova, V.N.Ravich, I.Z.Babievskaya, N.N.Kazarina, 

J.A.Poroikov: Izvestiya Akademii Nauk SSSR - Neorganicheskie Materialy, 1990, 26[9], 

1814-8. (Inorganic Materials, 1991, 26[9], 1552-5) 

[446-84/85-054] 

Cu 

InP: Cu Diffusion 

A study of Cu diffusion in both p-type and n type samples showed that this material 

exhibited a transition to semi-insulating behavior at relatively low Cu diffusion 

temperatures. It was found that all, or most, of the Cu precipitates formed a Cu-In 

compound, that both originally n-type and p-type material became semi-insulating, and 

that there was a negligibly low concentration of deep-level defects. It was observed that 

there was an abnormal reduction in both electron and hole mobilities, which resulted from 

the introduction of Cu, and that there were isolated pockets of highly conductive material 

in otherwise semi-insulating material. It was concluded that all of these experimental 

observations could be best explained by the buried Schottky barrier model instead of 

compensation by deep levels. 

R.P.Leon, M.Kaminska, K.M.Yu, Z.Liliental-Weber, E.R.Weber: Materials Science 

Forum, 1992, 83-87, 723-8 

[446-99/100-093] 

D 

InP: D Diffusion 

Data on the effusion of D from various samples were presented. It was shown that the 

effusion was limited by surface phenomena, and that decomposition of the sample surface 

played a major role. 

B.Theys, J.Chevallier, M.Miloche, B.Rose: Materials Science Forum, 1994, 143-147, 

945-50 

[446-113/114-037] 

460

D InP Fe 

InP: D Diffusion 

The problem of hydrogenating this material without causing surface degradation was 

solved by exposing the surface to a plasma via a thin SiN x (H) cap. This layer was 

permeable at the hydrogenation pressure of 250C, but was impermeable to P or PH 3 . It 

was found that shallow acceptors were heavily passivated, whereas shallow donors were 

only weakly affected. The presence of acceptors impeded D in-diffusion. Thus, D 

diffusion under the same conditions occurred to a depth of 0.018mm in p-type (2 x 10 16 

Zn/cm 3 ) material, but to a depth of 0.035mm in n-type (S, Sn) material. 

W.C.Dautremont-Smith, J.Lopata, S.J.Pearton, L.A.Koszi, M.Stavola, V.Swaminathan: 


[446-74-040] 

Fe 

InP: Fe Diffusion 

The diffusion of Fe was found to occur via the kick-out mechanism. A published Fe 

diffusion profile in InP was simulated by using the complete set of 3 partial differential 

equations for the kick-out mechanism. A value for the contribution which In selfinterstitials 

made to the self-diffusion coefficient of InP was deduced and was found to be 

much smaller than was suggested by the known self-diffusion coefficients which had 

been determined from In tracer diffusion measurements. 

H.Zimmermann, U.Gösele, T.Y.Tan: Applied Physics Letters, 1993, 62[1], 75-7 

[446-106/107-120] 


The doping and diffusion characteristics of Fe in semi-insulating layers were assessed by 

using secondary ion mass spectrometry. Fairly flat Fe depth profiles, and a linear doping 

curve, were obtained at concentrations of up to 10 17 /cm 3 . Accumulation of Fe at the 

substrate/layer interface was found in some samples; thus revealing a gettering effect of 

the substrate. Very little, and probably negligible, diffusion was observed on alternately 

Fe-doped and undoped structures; even after high-temperature heat treatment (provided 

that the Fe content was about 10 16 /cm 3 or less). 

D.Franke, P.Harde, P.Wolfram, N.Grote: Journal of Crystal Growth, 1990, 100[3], 309- 

12 

[446-76/77-026] 








461

Fe InP Fe 







1992, 72[8], 3514-21 

[446-106/107-034] 


It was pointed out that high-resistivity material could be grown by metalorganic vapor 

phase epitaxy, using ferrocene as a dopant source. Adjacent Zn-doped layers removed the 

resistivity of Fe-doped material. The presence of Zn markedly enhanced the out-diffusion 

of Fe from Fe-doped layers, and into Zn-doped material. It was suggested that interstitial 

Zn took over the lattice sites of substitutional Fe. The Fe then became interstitial and 

mobile. 

E.W.A.Young, G.M.Fontijn: Applied Physics Letters, 1990, 56[2], 146-7 

[446-74-040] 


A new technique, which involved a Zn 3 P 2 diffusion source and rapid thermal annealing, 

was studied. A p + -type layer could be obtained only at temperatures of between 500 and 

550C by using a 15s diffusion time. The diffusivity was calculated and was compared 

with the results of furnace annealing. The present diffusivity was deduced to be equal to 

2.6 x 10 -12 cm 2 /s. In order to form a shallow layer, it was necessary to avoid any 

treatments which might redistribute Fe or dopants. Annealing (850C, 15s) before 

diffusion shifted the carrier profile from 300 to 600nm in depth. The second diffusion 

front extended to 0.0024mm in a semi-insulating substrate. Diffusion treatments which 

were performed on samples without pre-annealing resulted in 2 diffusion fronts. A 

0.0028mm-deep second diffusion front was found when the treatment was applied to a 

pre-annealed epi-wafer. The diffusivity under these conditions was deduced to be equal to 

1.4 x 10 -11 cm 2 /s. 

K.W.Wang, S.M.Parker, C.L.Cheng, J.Long: Journal of Applied Physics, 1988, 63[6], 

2104-9 

[446-72/73-038] 


The diffusivity was measured by using secondary ion mass spectrometry. Deliberately 

doped metalorganic vapor phase epitaxial layers, as well as ion-implanted samples, were 

investigated. In addition, resistivity measurements were performed on Fe-doped layers. It 

was found that the diffusion behavior of Fe was strongly affected by the presence of Zn, 

and vice versa. In adjacent regions of Fe-doped and Zn-doped layers, there was a marked 

interdiffusion of the dopants. The interdiffusion process could be described in terms of a 

kick-out mechanism in which Fe interstitials kicked out substitutional Zn. The diffusion 

of Fe interstitials was an extremely fast transport process, but the concentration of Fe 

462

Fe InP H 

interstitials remained below 5 x 10 14 /cm 3 . Due to the rapid transport, interdiffusion 

proceeded even through barrier layers of (undoped) InP. In the barrier layer itself, the Fe 

concentration remained below the secondary ion mass spectrometric detection limit of 5 x 

10 14 /cm 3 . It was found that a S-doped n-type InP layer prevented the diffusion of Fe. The 

semi-insulating properties of Fe-doped InP were affected by the interdiffusion of Fe and 

Zn. Since S-doped InP inhibited interdiffusion, such a layer could be used as a barrier in 

order to separate Zn-doped and Fe-doped regions, and thus preserve the semi-insulating 

character of the Fe-doped InP. 

E.W.A.Young, G.M.Fontijn, C.J.Vriezema, P.C.Zalm: Journal of Applied Physics, 1991, 

70[7], 3593-9 

[446-93/94-040] 

Ge 

InP: Ge Diffusion 

Implantation of Ge into (x11)A-oriented semi-insulating InP substrates (where x took 


were also implanted. Following 200keV implantation (3 x 10 13 /cm 2 ), after annealing 

(850C, 7s), it was found that the InP was always n-type and had a similar sheet resistance 

regardless of the substrate orientation. No in-diffusion of Ge was observed after annealing 

substrates of any orientation. 



[446-117/118-166] 

H 

InP: H Diffusion 

The H passivation of Zn acceptors, and Zn-H dissociation kinetics, were compared for the 

cases of homo-epitaxial and lattice-mismatched hetero-epitaxial n + p structures. Doping 

profile measurements revealed a marked increase in the depth and degree of passivation 

in the p-type region of hetero-epitaxial samples. This indicated an enhanced diffusion of 

H along dislocations, followed by additional Zn deactivation. Also, the strong affinity 

between H and extended defects was found to promote the subsequent dissociation of Zn- 

H complexes. This was revealed by reverse bias annealing studies which showed that the 

Zn-H dissociation energy decreased, from 1.19eV in homo-epitaxial samples, to 1.12eV 

in hetero-epitaxial samples. Another indicator was the enhanced passivation of extended 

defect-related traps by H that was liberated from Zn acceptors during the reverse bias 

annealing process; as determined by means of deep level transient spectroscopy. 

B.Chatterjee, S.A.Ringel: Applied Physics Letters, 1996, 69[6], 839-41 

[446-138/139-097] 

463

H InP In 


The migration of H was much higher in p-type than in n-type samples. It was concluded 

that H was a deep donor in this material. 

E.M.Omeljanovsky, A.V.Pakhomov, A.Y.Polyakov, O.M.Borodina, E.A.Kozhukhova, 

A.Y.Nashelskii, S.V.Yakobson, V.V.Novikova: Solid State Communications, 1989, 

72[5], 409-11 

[446-72/73-035] 


The problem of hydrogenating this material without causing surface degradation was 

solved by exposing the surface to a H plasma via a thin SiN x (H) cap. This layer was 

permeable to H at the hydrogenation pressure of 250C, but was impermeable to P or PH 3 . 

It was found that shallow acceptors were heavily passivated, whereas shallow donors 

were only weakly affected. The presence of acceptors impeded H in-diffusion. 

W.C.Dautremont-Smith, J.Lopata, S.J.Pearton, L.A.Koszi, M.Stavola, V.Swaminathan: 


[446-74-040] 

In 

InP: In Diffusion 

Single crystals with a [123]-type orientation were deformed by using constant strain rates 

at temperatures of between 540 and 780C. The resultant stress-strain curves were rather 

similar to those which were observed for Ge, Si and InSb. Two stages of dynamic 

recovery could be clearly identified. From the strain-rate and temperature dependences of 

the stress at the beginning of the first recovery stage, an activation energy of 2.3eV was 

deduced. This was regarded as being a lower bound on the activation energy for selfdiffusion 

of the slowest-moving species. A value of between 0.001 and 0.01cm 2 /s was 

estimated for the pre-exponential factor. 

H.Siethoff, K.Ahlborn, H.G.Brion, J.Völkl: Philosophical Magazine A, 1988, 57[2], 235- 

44 

[446-60-009] 

InP: In Diffusion 

The slow evaporation of In from wafers was investigated, and was identified as being a 

major cause of surface roughening during thermal annealing and mass transport under 

adequate P-vapor protection. Analysis showed that In evaporation could be prevented by 

covering the wafer during annealing. However, the usual graphite cover alone was an 

inadequate protection because the In vapor was able to permeate the graphite. On the 

other hand, the use of an InP cover alone resulted in problems which were caused by the 

mass transport that occurred between the InP wafer and the cover. A scheme was 

464

In InP Mg 

developed that used InP covers and a quartz enclosure in addition to a graphite cover. 

This arrangement permitted smooth wafer surfaces to be reproducibly obtained. 

Z.L.Liau: Applied Physics Letters, 1991, 58[17], 1869-71 

[446-81/82-040] 

Mg 

InP: Mg Diffusion 

The diffusion mechanism of Mg was studied during low-pressure metalorganic vaporphase 

epitaxial growth. The Mg dopant profiles were measured by means of secondary 

ion mass spectroscopy. The analysis revealed that abrupt Mg dopant profiles were 

possible. However, the Mg diffusivity depended markedly upon the Mg concentration in 

the crystal lattice. Simultaneous doping with Si led to a distinct decrease in Mg diffusion. 

This behavior was consistent with a model which assumed that the Mg diffused as a 

complex which involved a deep donor. 

E.Veuhoff, H.Baumeister, R.Treichler, O.Brandt: Applied Physics Letters, 1989, 55[10], 

1017-9 

[446-72/73-038] 


Samples of Fe-doped material were implanted with 80keV Mg, Mg and P, or Mg and Ar 

in order to produce shallow p + layers. After rapid thermal annealing (850 or 875C, 5 or 

10s), activations of between 10 and 50% and mobilities of up to 110cm 2 /Vs were 

obtained. Secondary ion mass spectrometry profiles revealed a pile-up of Mg at the 

surface, and in-diffusion tails which were deeper than 2µ. The use of P or Ar coimplantation 

reduced Mg in-diffusion and increased the activation, but not as much as in 

the case of Be implantation. Photoluminescence measurements revealed good crystalline 

quality after annealing. Narrow emissions close to the gap wavelength, and 2 broad bands 

which were centered at about 1.3 and 0.87eV, were found in the photoluminescence 

spectra. The bands were the predominant emission in the photoluminescence spectra of 

layers with higher implanted doses. This band was tentatively attributed to complexes that 

involved Mg and a defect. 

J.M.Martin, S.García, F.Calle, I.Mártil, G.González-Díaz: Journal of Electronic 

Materials, 1995, 24[1], 59-67 

[446-119/120-217] 


The annealing behavior of implanted Mg was studied. It was found that the activated 

fraction of dopants depended markedly upon the implant dose, and upon the substrate 

temperature during implantation. Low activation of high-dose (10 15 /cm 2 ) implants was 

465

Mg InP S 

attributed to the effect of pronounced (80%) out-diffusion. A large variation was found, 

in the apparent activation energy, for implantation temperatures between ambient and 

300C. 

W.H.Van Berlo, M.Ghaffari, G.Landgren: Journal of Electronic Materials, 1992, 21[4], 

431-6 

[446-93/94-040] 

P 

InP: P Diffusion 

Single crystals with a [123]-type orientation were deformed by using constant strain rates 

at temperatures of between 540 and 780C. The resultant stress-strain curves were rather 

similar to those which were observed for Ge, Si and InSb. Two stages of dynamic 

recovery could be clearly identified. From the strain-rate and temperature dependences of 

the stress at the beginning of the first recovery stage, an activation energy of 2.3eV was 

deduced. This was regarded as being a lower bound on the activation energy for selfdiffusion 

of the slowest-moving species. A value of between 0.001 and 0.01cm 2 /s was 

estimated for the pre-exponential factor. 

H.Siethoff, K.Ahlborn, H.G.Brion, J.Völkl: Philosophical Magazine A, 1988, 57[2], 235- 

44 

[446-60-009] 

Rh 

InP: Rh Diffusion 

A secondary-ion mass spectroscopic investigation was made of the thermally induced 

redistribution of Rh in low-pressure metalorganic chemical vapor-deposited InP 

structures. Control measurements were performed on Fe-doped structures. In the case of 

alternately Rh-doped InP and undoped InP structures, an upper limit on the Rh diffusion 

coefficient of about 10 -14 cm 2 /s (at 800C) was established. This was much smaller than the 

Fe diffusivity of about 10 -11 cm 2 /s at 750C. No exchange reactions were observed at the 

interfaces of p-InP and Rh-doped InP structures. Only Rh which was implanted into InP 

exhibited defect-induced redistribution into amorphous areas. 

A.Näser, A.Dadgar, M.Kuttler, R.Heitz, D.Bimberg, J.Y.Hyeon, H.Schumann: Applied 


[446-123/124-177] 

S 

InP[l]: S Diffusion 

Macro-segregation and micro-segregation of S in crystals which had been grown from In 

solutions by using the travelling heater method (under micro-gravity or normal gravity 

conditions) were analyzed by using spatially resolved photoluminescence methods. It was 

466

S InP Si 

found that, whereas macro-segregation in both terrestrially-grown and space-grown 

crystals could be explained by conventional steady-state models which were based upon 

the Burton-Prim-Slichter theory, micro-segregation could be explained only in terms of 

the non steady-state step-exchange model. 

A.N.Danilewsky, Y.Okamoto, K.W.Benz, T.Nishinaga: Japanese Journal of Applied 

Physics, 1992, 31[1-7], 2195-201 

[446-93/94-231] 

Si 

InP: Si Diffusion 

Implantation of Si into (x11)A-oriented semi-insulating InP substrates (where x took 


were also implanted. Following 200keV implantation (5 x 10 13 /cm 2 ), after annealing 

(850C, 7s), it was found that the InP was always n-type and had a similar sheet resistance 

regardless of the substrate orientation. No in-diffusion of Si was observed after annealing 

substrates of any orientation. A similar behavior was observed for Si/B co-implants in 

InP. 



[446-117/118-166] 


Semi-insulating material was implanted with 200keV Xe + ions to 10 14 /cm 2 at room 

temperature, and with 100keV Hg + ions to 10 14 /cm 2 at 200C or room temperature. 

Implanted and non-implanted substrates were encapsulated in AlN/Si 3 N 4 and were 

subjected to rapid (60s) thermal annealing cycles at 650 to 900C. Electrical measurements 

and secondary ion mass spectrometry were used to correlate an observed n-type behavior 

with the presence of Si in the near-surface region. An enhanced near-surface Si 

concentration was found after the rapid thermal annealing of Xe + -implanted samples, and 

n-type surface layers (approximately 80nm thick) were formed. Non-implanted samples 

exhibited no measurable electrical behavior, and there was no enhanced Si concentration 

after rapid thermal annealing. Hot Hg + implantation led to p-type behavior and to low Si 

concentrations. Room-temperature Hg implantation led to semi-insulating behavior and 

high Si levels. It was concluded that the presence of implantation damage enhanced Si indiffusion 

from the Si-based encapsulants which were used to protect the InP surface 

during post-implantation annealing. 

J.H.Wilkie, B.J.Sealy: Thin Solid Films, 1988, 162, 49-57 

[446-70/71-119] 


Migration was investigated by using two configurations. These were diffusion from an 

external source into uniformly n-doped substrates, and diffusion between the layers of n- 

467

Si InP Sn 

p-n-p-n structures which had been grown via metalorganic chemical vapor deposition. 

Alternating layers of n-type material (0.0005mm, [Si] = 10 16 to 3 x 10 19 /cm 3 ) were grown 

by using low-pressure metalorganic chemical vapor deposition at 625C. The distributions 

of Si were determined by means of secondary ion mass spectrometry. No diffusion of Si 

across the grown dopant interface was detected. Electrochemical capacitance-voltage 

profiling indicated that the Si was electrically active. 

C.Blaauw, F.R.Shepherd, D.Eger: Journal of Applied Physics, 1989, 66[2], 605-10 

[446-74-041] 


Material was deposited by means of metalorganic chemical vapor deposition, and was 

simultaneously doped with Si (donor) and Zn (acceptor) species during growth. It was 

found that the incorporation of Si was not affected by the presence of Zn, whereas Zn 

incorporation was markedly enhanced by the presence of Si. The results were consistent 

with the formation of donor-acceptor pairs. 

C.Blaauw, L.Hobbs: Applied Physics Letters, 1991, 59[6], 674-6 

[446-84/85-054] 

InP/Si: Si Diffusion 

The spatial distribution of the charge concentration of InP layers which had been grown 

onto Si substrates by metalorganic vapor-phase epitaxy was investigated. The 

concentrations near to the surface, and within the bulk of the layer, were found to be 

governed by Si doping from the ambient gas. The diffusion of Si across the heterointerface, 

which could be partially assisted by dislocations, predominated in a region near 

to the InP/Si interface. In the vicinity of the hetero-interface, the charge concentration in 

the InP layer was determined by a strong compensation which was attributed to defects 

that were caused by a mismatch between the lattice parameters and thermal expansion 

coefficients of the InP and Si. 

A.Bartels, E.Peiner, R.Klockenbrink, A.Schlachetzki: Journal of Applied Physics, 1995, 

78[1], 224-8 

[446-123/124-179] 

Sn 

InP: Sn Diffusion 

The lattice location and electrical activity of ion implanted Sn, after rapid thermal 

annealing, were determined by means of Mössbauer spectroscopic (using 119m Sn) and 

differential Hall resistivity methods, respectively. It was found that the Sn was 

preferentially located on the In sub-lattice at concentrations below 2 x 10 19 /cm 3 ; resulting 

in high electrical activation and mobility. At Sn concentrations which were above this 

468

Sn InP Zn 

concentration, various electrically inactive Sn complexes were also observed. No sign 

was found of Sn on P sub-lattice sites. 

P.Kringhøj, G.Weyer: Applied Physics Letters, 1993, 62[16], 1973-5 

Ti 

[446-99/100-093] 

InP: Ti Diffusion 







1992, 72[8], 3514-21 

[446-106/107-034] 

InP: Ti Diffusion 

Samples of n-type InP were implanted with Co at 200C. During high-temperature 

annealing, out-diffusion of the implant was as severe as that for room-temperature 

implants. In-diffusion of the implant also occurred, but it was not as severe as the outdiffusion. 

High-temperature annealing of Ti-implanted material resulted in slight Ti indiffusion, 

with minimal redistribution or out-diffusion. In the case of high-temperature 

implants, the lattice quality of the annealed material was close to that of virgin material. 

Regardless of the ion type, resistivities that were close to the intrinsic limit were 

measured in implanted and annealed materials. 

M.V.Rao, S.M.Gulwadi, S.Mulpuri, D.S.Simons, P.H.Chi, C.Caneau, W.P.Hong, 

O.W.Holland, H.B.Dietrich: Journal of Electronic Materials, 1992, 21[9], 923-8 

[446-93/94-038] 

Zn 

InP: Zn Diffusion 

Spun-on SiO 2 films which contained 1, 2, 6, 22 or 36at%Zn were prepared and were used 

for p-diffusion into undoped n-type material. It was found that an increasing Zn 

concentration of the spun-on film led to a higher atomic Zn concentration and diffusion 

depth. On the basis of the experimental data, the InP/film distribution coefficient was 

deduced to be 0.012. A higher Zn concentration in the InP resulted in a lower acceptor 

concentration. The electrical activity of Zn could be significantly increased by means of 

additional heat treatment. 

C.Lauterbach: Semiconductor Science and Technology, 1995, 10[4], 500-3 

[446-121/122-080] 

469

Zn InP Zn 


The formation of defects during Zn diffusion into undoped or semi-insulating Fe-doped 

single crystals at 700C was observed by means of transmission electron microscopy under 

various diffusion conditions. Agglomerates of predominantly perfect interstitial-type 

dislocation loops, dislocations, and small In precipitates within voids were observed in 

the Zn-diffused region. Also, large planar arrays of precipitates were formed by climbing 

dislocations. From these observations, it was deduced that the incorporation of Zn at In 

sub-lattice sites created a supersaturation of In self-interstitials which was removed by 

dislocation loop formation that led to a supersaturation of P vacancies and to void 

formation. 

D.Wittorf, A.Rucki, W.Jäger, R.H.Dixon, K.Urban, H.G.Hettwer, N.A.Stolwijk, 

H.Mehrer: Journal of Applied Physics, 1995, 77[6], 2843-5 

[446-121/122-080] 


The diffusivity was measured by using secondary ion mass spectrometry. Deliberately 

doped metalorganic vapor phase epitaxial layers, as well as ion-implanted samples, were 

investigated. In addition, resistivity measurements were performed on Fe-doped layers. It 

was found that the diffusion behavior of Fe was strongly affected by the presence of Zn, 

and vice versa. In adjacent regions of Fe-doped and Zn-doped layers, there was a marked 

interdiffusion of the dopants. The interdiffusion process could be described in terms of a 

kick-out mechanism in which Fe interstitials kicked out substitutional Zn. 

E.W.A.Young, G.M.Fontijn, C.J.Vriezema, P.C.Zalm: Journal of Applied Physics, 1991, 

70[7], 3593-9 

[446-93/94-040] 


A photoluminescence study was made of Zn-diffused and annealed material. A new peak 

was found near to 1.33eV. After annealing, the peak energy of the luminescence shifted 

towards higher energies. This reflected the out-diffusion of Zn. By depth profiling of this 

luminescence it was deduced that recombination, due to the interstitial Zn donor, 

predominated near to the surface. 

J.S.Choi, H.J.Lim, J.I.Lee, S.K.Chang, H.L.Park: Physica Status Solidi B, 1991, 164[2], 

K69-72 

[446-93/94-041] 


The diffusivity was studied by using dimethylzinc as the source in a PH 3 /H 2 /N 2 

atmosphere at 500 to 600C. The results were compared with the Zn doping of material 

which had been grown via metalorganic vapor-phase epitaxial growth using 

triethylindium, PH 3 and dimethylzinc. It was found that the carrier concentration in Zndiffused 

material was lower, by about 2 orders of magnitude, than that in Zn-doped 

material which had been grown at the same mole fraction. The activation energy for Zn 

470

Zn InP Zn 

incorporation was deduced to be the same for both diffusion and doping. The observed 

dependence of the surface carrier concentration upon the diffusion temperature and the 

mole fraction of dimethylzinc was qualitatively explained by considering an InP surface 

condensation limit for Zn adsorption. Slightly increased carrier concentrations, after 

subsequent annealing (500C, PH 3 /H 2 /N 2 atmosphere) indicated the existence of a low 

concentration of Zn interstitials in as-diffused InP under a high PH 3 flow-rate. The lower 

carrier concentration of shallow Zn-diffused material, after such annealing, suggested the 

occurrence of a large out-diffusion of interstitials near to the surface region during 


M.Wada, K.Sakakibara, M.Higuchi, Y.Sekiguchi: Journal of Crystal Growth, 1991, 

114[3], 321-6 

[446-91/92-025] 


A wafer of InP, with an evaporated thin layer of metallic Zn, was used as a diffusion 

source and was placed 0.15mm away from the sample. In this way, it was possible to 

obtain good diffusion-front planarity, perfect surface quality, and a free-hole 

concentration of about 8 x 10 18 /cm 3 . It was suggested that the first stage of the diffusion 

process was characterized by the incorporation of a large quantity of Zn, such that 

interstitial Zn atoms compensated the acceptors. The Zn concentration then levelled out, 

but most of the interstitials left the crystal. Due to the small free volume which existed 

between sample and source, P loss was decreased and the formation of P vacancies was 

diminished. 

T.Krieghoff, E.Nowak, G.Kühn, B.Schumann, A.Höpner: Crystal Research and 

Technology, 1992, 27[1], 49-57 

[446-88/89-044] 


The migration of Zn during the growth of InP epitaxial layers was investigated in 

structures which consisted of Zn-InP epilayers that had been grown onto S-InP and Fe- 

InP substrates, and onto undoped InP epilayers. The layers were grown by means of 

metalorganic chemical vapor deposition at 625C, under a pressure of 75torr. The dopant 

diffusion profiles were measured by using secondary ion mass spectrometry. At high Zn 

dopant levels (greater than 8 x 10 17 /cm 3 ), diffusion into S-InP substrates took place, with 

Zn accumulation in the substrate at a concentration which was similar to [S]. Diffusion 

into undoped InP epilayers produced a diffusion tail at low [Zn] levels. This was 

suggested to be associated with interstitial Zn diffusion. In the case of diffusion into Fe- 

InP, this low-level diffusion produced a region of constant Zn concentration when the 

dopant concentration was equal to 3 x 10 16 /cm 3 . This was attributed to kicking-out of the 

original Fe species from substitutional sites. Diffusion out of (Zn,Si) co-doped InP 

epilayers which had been grown onto Fe-InP substrates was also investigated. The 

secondary ion mass spectrometry profiles were characterized by a sharp decrease in [Zn] 

at the epilayer/substrate interface. The magnitude of this decrease corresponded to that of 

the Si donor level in the epilayer. When [Si] was greater than [Zn] in the epilayer, no Zn 

471

Zn InP Zn 

diffusion was observed. Hall measurements indicated that the donor and acceptor species 

in those samples were electrically active. All of the results were consistent with the 

presence of donor-acceptor interactions; resulting in the formation of ionized donoracceptor 

pairs which were immobile and did not contribute to the diffusion process. 

C.Blaauw, B.Emmerstorfer, D.Kreller, L.Hobbs, A.J.Springthorpe: Journal of Electronic 

Materials, 1992, 21[2], 173-9 

[446-88/89-045] 


A low-pressure open-tube system, involving diethylzinc and PH 3 , was used to study the 

diffusion of Zn. The Zn and hole concentrations were measured by using secondary ion 

mass spectrometry and capacitance-voltage etch profiling. It was found that annealing of 

the samples increased the hole concentration, due to the out-diffusion of interstitial Zn 

donors. 

J.Wisser, M.Glade, H.J.Schmidt, K.Heine: Journal of Applied Physics, 1992, 71[7], 

3234-7 

[446-86/87-042] 


Phase diagrams between Zn and III-V compounds were investigated in order to develop 

new acceptor diffusion sources. On the basis of thermodynamic data and experimental 

studies of the Cd-Ga-As and Cd-In-P phase diagrams, it was found that Zn 3 Cd 3 B 2 and 

ZnCdB 2 were suitable acceptor sources in that they did not erode the semiconductor 

surface. 

S.F.Marenkin, O.N.Pashkova, V.N.Ravich, I.Z.Babievskaya, N.N.Kazarina, 

J.A.Poroikov: Izvestiya Akademii Nauk SSSR - Neorganicheskie Materialy, 1990, 26[9], 

1814-8. (Inorganic Materials, 1991, 26[9], 1552-5) 

[446-84/85-054] 


Material was deposited by means of metalorganic chemical vapor deposition, and was 

simultaneously doped with Si (donor) and Zn (acceptor) species during growth. It was 

found that the incorporation of Si was not affected by the presence of Zn, whereas Zn 

incorporation was markedly enhanced by the presence of Si. The results were consistent 

with the formation of donor-acceptor pairs; a concept which had previously been used to 

explain Zn diffusion profiles in Si-doped InP. 

C.Blaauw, L.Hobbs: Applied Physics Letters, 1991, 59[6], 674-6 

[446-84/85-054] 


Diethylzinc was used as a p-type dopant source during the chemical beam epitaxial 

growth of InP. Secondary ion mass spectrometry measurements indicated that very 

472

Zn InP Zn 

marked Zn diffusion occurred when the Zn concentration appeared to reduce the pyrolytic 

efficiency of trimethylindium. 

W.T.Tsang, F.S.Choa, N.T.Ha: Journal of Electronic Materials, 1991, 20[7], 541-4 

[446-84/85-054] 


The non-intentional diffusion of Zn acceptors was investigated during low-pressure 

metalorganic vapor-phase epitaxial growth at 550 or 640C. It was found that the diffusion 

of Zn during deposition could be described by an interstitial-substitutional model. The 

diffusivity in the non-intentionally doped material was lower than that in intentionally 

doped material. This was attributed to the low concentration of interstitial Zn atoms in 

samples which were doped during growth. Some deposition parameters, such as a high 

temperature and a high V/III ratio, minimized diffusion. In this way, a normalized 

diffusivity which could be as low as 6.5 x 10 -14 cm 2 /s could be obtained at a dopant level 

of 10 18 /cm 3 . 

M.Glade, J.Hergeth, D.Grützmacher, K.Masseli, P.Balk: Journal of Crystal Growth, 

1991, 108[3-4], 449-54 

[446-84/85-054] 


The saturation behavior of the free carrier concentrations in p-type InP monocrystals 

which had been doped by Zn diffusion was investigated. The maximum free-hole 

concentration appeared at about 5 x 10 18 /cm 3 . The difference in saturation hole 

concentrations of materials was investigated by studying the incorporation and lattice 

location of Zn. The latter was an acceptor when located on a group-III atom site. Zinc 

was diffused into III-V wafers in a sealed quartz ampoule. Particle-induced X-ray 

emission and ion-channelling techniques were then used to determine the exact lattice 

location of Zn atoms. In InP, the substitutional state of Zn depended upon the cooling rate 

of the sample after high-temperature diffusion. In slowly cooled samples, a large fraction 

(about 90%) of the Zn atoms formed random precipitates of Zn 3 P 2 and elemental Zn. 

However, after rapid cooling, only 60% of the Zn atoms formed such precipitates while 

the remainder occupied specific sites. The results were analyzed in terms of the 

amphoteric native defect model. It was shown that differences in the electrical activities 

of Zn atoms were a consequence of differing locations of the Fermi-level stabilization 

energy. 

L.Y.Chan, K.M.Yu, M.Ben-Tzur, E.E.Haller, J.M.Jaklevic, W.Walukiewicz, 

C.M.Hanson: Journal of Applied Physics, 1991, 69[5], 2998-3006 

[446-78/79-015] 


The characteristics of Fe-doped semi-insulating layers, with overgrown Zn-doped p-type 

layers, were investigated by means of scanning electron microscopy, secondary ion mass 

spectrometry, capacitance-voltage, and current-voltage measurements. The resistivity 

which was deduced from the current-voltage characteristics was found to be strongly 

dependent upon the Zn dopant concentration. The secondary ion mass spectrometry depth 

473

Zn InP Zn 

profiles revealed the occurrence of Zn accumulation at the semi-insulating/p-type 

interface, and the peak concentration of Zn accumulation increased with doping level and 

overgrowth time of the p-type layers. The accumulation of Zn at the semi-insulating/ptype 

interface was related to a reduction in semi-insulating layer resistivity. The 

accumulation of Zn at the interface could be minimized by using a short growth time, 

together with low or medium doping of the p-type layers. Such growth conditions led to 

higher semi-insulating layer resistivity. 

W.H.Cheng, H.Kuwamoto, A.Appelbaum, D.Renner, S.W.Zehr: Journal of Applied 

Physics, 1991, 69[4], 1862-5 

[446-78/79-045] 


The migration of Zn was investigated by using two configurations. These were diffusion 

from an external source into uniformly n-doped substrates, and diffusion between the 

layers of n-p-n-p-n structures which had been grown via metalorganic chemical vapor 

deposition. Alternating layers of p-type material (0.0005mm, [Zn] = 4 x 10 17 to 2 x 10 18 

/cm 3 ) and n-type material (0.0005mm, [Si] = 10 16 to 3 x 10 19 /cm 3 ) were grown by using 

low-pressure metalorganic chemical vapor deposition at 625C. The distributions of Zn 

were determined by means of secondary ion mass spectrometry. In the case of un-doped 

spacer layers (with n approximately equal to 10 16 /cm 3 ), the diffusion profiles depended 

markedly upon the Zn dopant level. Little Zn out-diffusion was observed when [Zn] was 

equal to 4 x 10 17 /cm 3 . When [Zn] was greater than 10 18 /cm 3 , the Zn diffused completely 

across the spacer layers during growth times of 1 to 2h. In the case of doped spacer 

layers, the doping level of Si had a marked effect upon the Zn diffusion profiles. The total 

Zn diffusion across the grown dopant interface was not substantially affected, but 

accumulation of Zn occurred in the Si-doped layers; with the formation of Zn spikes for 

which the increase in Zn level - as compared to that (about 10 18 /cm 3 ) of the Zn-doped 

layer - was similar to [Si]. Electrochemical capacitance-voltage profiling indicated that 

the Zn was electrically active. The results were explained in terms of a model in which 

the mobile Zn species that diffused into the Si-doped layers were immobilized by the 

formation of Zn-donor pairs. This model was shown to be consistent with the profiles 

which were obtained for Zn diffusion into n-type material from an external ZnGaCdIn 

source. 

C.Blaauw, F.R.Shepherd, D.Eger: Journal of Applied Physics, 1989, 66[2], 605-10 

[446-74-041] 


The ampoule diffusion of Zn gave rise to donor-acceptor photoluminescence transitions 

with peak positions that depended upon the cooling rate after diffusion. Subsequent 

annealing in a Zn-free ambient caused a shift in the peak position. Luminescence peaks 

were found between 1.30 and 1.38eV. These peaks were attributed to transitions between 

various Zn interstitial donor levels and the Zn substitutional acceptor level. The 

luminescence data were correlated with secondary ion mass spectrometry and Schottky 

barrier capacitance-voltage measurements, and were found to be consistent with an earlier 

474

Zn InP Zn 

model in which Zn was supposed to diffuse as both an interstitial donor and a 

substitutional acceptor. 

E.A.Montie, G.J.Van Gurp: Journal of Applied Physics, 1989, 66[11], 5549-53 

[446-74-041] 


The diffusion of Zn was studied by using boat diffusion, diffusion from As- or P-doped 

spun-on films, or diffusion from In-doped spun-on films. The depth profiles were 

deduced from junction positions. It was found that the p + /p - junction position depended 

upon the diffusion method which was used, but not upon the sample growth technique. 

The p - /n junction position depended upon both factors. Because the amounts of In and P 

(or of the respective vacancies) differed, it was possible to identify diffusion mechanisms. 

It was proposed that interstitially diffusing Zn was independently trapped by 2 immobile 

vacancy centers. These consisted of Zn on V In in the p + region, and Zn on V P ZnV P in the 

p - region. 

U.König, H.Haspeklo, P.Marschall, M.Kuisl: Journal of Applied Physics, 1989, 65[2], 

548-52 

[446-72/73-035] 


Closed-ampoule diffusion led to a net acceptor concentration which was lower than the 

Zn concentration. Upon annealing in an atmosphere without Zn, the Zn and net acceptor 

concentrations became almost identical. This was attributed to a decreased Zn 

concentration and an increased net acceptor concentration. The results were quantitatively 

explained by assuming that the Zn was incorporated as both substitutional acceptors and 

interstitial donors, and that only the interstitial Zn was driven out by annealing; due to its 

large diffusion coefficient. The profiles which were calculated by using this interstitialsubstitutional 

model could be fitted to experimentally determined profiles by assuming 

that the Zn diffused as singly-ionized interstitial donors. The present model also 

explained published data, on diffusion in n-type InP, in which a profile cut-off was found 

at a depth where the acceptor concentration equalled the background donor concentration. 

G.J.Van Gurp, T.Van Dongen, G.M.Fontijn, J.M.Jacobs, D.L.A.Tjaden: Journal of 


[446-72/73-037] 


The Zn was diffused, from a dimethylzinc source, at temperatures ranging from 400 to 

570C using surface concentrations ranging from 10 17 to 10 18 /cm 3 . The resultant diffusion 

profiles were determined by using secondary ion mass spectrometry, electrolytic etching, 

and capacitance-voltage measurements. The results indicated that an interstitialsubstitutional 

mechanism operated at the above concentrations. 

M.Wada, M.Seko, K.Sakakibara, Y.Sekiguchi: Japanese Journal of Applied Physics, 

1989, 28[10], L1700-3 

[446-72/73-039] 

475

Zn InP Zn 


It was pointed out that the use of models, which did not involve dopant effects upon the 

interstitial-substitutional interchange, could lead to the identification of an apparently 

larger charge state for the interstitial. This was expected to increase the apparent 

activation energy for diffusion. This increase was approximately equivalent to the Zn 

solubility activation energy for substrate doping near to the substitutional Zn 

concentration at the surface. 

C.Kazmierski: Journal of Applied Physics, 1988, 64[11], 6573-5 

[446-72/73-039] 


The Zn was diffused into an unpassivated surface, from an open gas flow system, at 

temperatures of between 733 and 773K. In the region where the carrier concentration 

profile could be described by an erfc function, the diffusivity was given by: 

D(m 2 /s) = 3 x 10 -7 exp[-120(kJ/mol)/RT] 

It was shown that thermal processes caused changes in the charge state of Zn in InP. 

These resulted in a variation of the carrier profile. 

T.O.Budko, E.V.Gushchinskaya, J.S.Emelyanenko, S.A.Malyshev: Physica Status Solidi 

A, 1989, 111[2], 451-6 

[446-64/65-174] 


The profiles of Zn in n-type [100] wafers after ampoule diffusion were measured using 

secondary-ion mass spectrometry, Auger electron spectrometry, differential Hall-effect 

measurements, capacitance measurements, and scanning electron microscopy. The results 

could be explained in terms of an interstitial-substitutional mechanism in which the Zn 

diffused as a singly ionized interstitial and was incorporated into the In sub-lattice as an 

electrically active substitutional acceptor or as an electrically inactive complex. At Zn 

concentrations which were lower than the background donor concentration, the profile 

was cut off as interstitial diffusion broke down. The activation energies for diffusion and 

solubility were found to be 1.40 and 1.0eV, respectively. 

G.J.Van Gurp, P.R.Boudewijn, M.N.C.Kempeners, D.L.A.Tjaden: Journal of Applied 

Physics, 1987, 61[5], 1846-55 

[446-60-009] 


The results of open-tube Zn diffusion into undoped or S-doped n-type material at 

temperatures ranging from 550 to 675C were presented. The results were consistent with 

interstitial-substitutional diffusion. In the case of undoped samples, the results were 

described by: 

D (cm 2 /s) = 0.049 exp[-1.52(eV)/kT] 

In the case of heavily S-doped samples, the results were described by: 

D (cm 2 /s) = 1400 exp[-2.34(eV)/kT] 

476

Zn InP Zn 

The difference in the activation energies was comparable to the Fermi level difference for 

the two substrate types, and was consistent with the differing diffusion mechanisms 

which occurred in these two types of InP. 

H.S.Marek, H.B.Serreze: Applied Physics Letters, 1987, 51[24], 2031-3 

[446-60-010] 


The substitutional fraction of Zn atoms which was diffused into single crystals was 

measured by using the proton-induced X-ray excitation technique. The diffusion times 

ranged from 0.25 to 1h at 425 to 650C. For several samples with diffusion depths ranging 

from 0.00075 to 0.0037mm (as determined using secondary ion mass spectrometry), it 

was found that the Zn impurity atoms resided almost entirely on lattice sites and that the 

substitutional fraction was equal to 0.9. There was no evidence of precipitation in the 

diffused layers. Only 1 to 10% of the Zn was electrically active. This was consistent with 

the existence of neutral V P Zn In V P complexes. 

W.N.Lennard, M.L.Swanson, D.Eger, A.J.Springthorpe, F.R.Shepherd: Journal of 


[446-60-010] 


It was recalled that, when p-n junctions were formed by doping with an element that 

diffused via a dissociative mechanism, dopant diffusion was suppressed and dopants 

could pile up near to the junction; at well above their original concentration. Calculations 

confirmed this behavior, if no local neutrality was assumed. The results agreed well with 

published experimental data on Zn diffusion in the present material. It was noted that the 

increased built-in electric field, due to this pile-up, was expelled almost entirely to the 

side of the junction without the pile-up. It was suggested that this effect had important 

implications for devices, which contained thin and/or small regions that were doped with 

such elements, because such regions might become completely depleted. 

I.Lyubomirsky, V.Lyahovitskaya, D.Cahen: Applied Physics Letters, 1997, 70[5], 613-5 

[446-150/151-148] 


The initial stages of Zn diffusion into InP from a polymer spin-on film were investigated. 

It was found that there were high concentrations of micro-defects and extended defects in 

the near-surface region (down to 1µ), an anomalously deep penetration of Zn atoms (with 

a diffusivity that was almost independent of temperature) and a low degree of activation 

of the diffused Zn. A kick-out mechanism was thought to predominate in the initial 

stages. 

A.V.Kamanin, I.A.Mokina, N.M.Shmidt: Solid-State Electronics, 1996, 39[10], 1441-4 

[446-141/142-111] 

477

Zn InP Zn 


It was noted that substitutional Zn was linearly incorporated into device-quality material 

under low Zn-source flow-rates during atmospheric-pressure metalorganic vapor-phase 

epitaxy at 625C. It saturated at about 4 x 10 18 /cm 3 under high Zn-source flow-rates. An 

increase, in the Zn-source flow-rate, to beyond saturation significantly increased the 

amount of interstitial incorporation. The excess interstitials diffused into the undoped 

region via an interstitial-substitutional diffusion mechanism, and revealed themselves via 

an enhanced diffusivity. It was recalled that a model had previously been proposed, for 

surface adsorption-desorption trapping during substitutional Zn incorporation, in which 

the saturation level was assumed to be governed by surface incorporation sites for 

substitutional Zn. This model was applied here, to interstitial Zn incorporation at Zn 

source flow rates which were above the saturation level for substitutional Zn, in order to 

explain the enhanced Zn diffusion. The analysis was extended so as to include the 

incorporation of neutral Zn in the presence of excess P vacancies. It was concluded that 

this model could be used for the simultaneous incorporation of Zn of all 3 types during 

epitaxy; provided that the incorporation processes were independent. 

S.N.G.Chu, R.A.Logan, M.Geva, N.T.Ha, R.F.Karlicek: Journal of Applied Physics, 

1996, 80[6], 3221-7 

[446-138/139-098] 


The diffusion and incorporation characteristics of Zn dopants in organometallic vaporphase 

epitaxially grown material were studied. The Zn diffusion coefficient depended 

strongly upon the concentration, and increased by 4 orders of magnitude for Zn 

concentrations of between 2 x 10 18 and 8 x 10 18 /cm 3 . This marked concentration 

dependence of the Zn diffusion coefficient was shown to govern Zn incorporation during 

organometallic vapor-phase epitaxial growth. The spread of Zn dopants into intentionally 

undoped regions could result in high Zn dopant concentrations. 

E.F.Schubert, C.J.Pinzone, M.Geva: Applied Physics Letters, 1995, 67[5], 700-2 

[446-123/124-178] 


The concentration-dependent diffusion of Zn during metalorganic vapor-phase epitaxy 

from a Zn-doped InP layer, and into the adjacent undoped InP buffer layer, was studied 

by means of secondary ion mass spectroscopy and carrier concentration profiling. If the 

growth rate of the Zn-doped film was faster than the interdiffusion of Zn into the 

underlying undoped buffer layer, the diffusion problem could be treated as a 1- 

dimensional couple between 2 semi-infinite media. Also, Zn diffusion under optimum 

growth conditions completely eliminated the thermal decomposition problem which was 

encountered when using sealed-ampoule or open-tube methods, and also retained all of 

the intrinsic point defects in their thermodynamic equilibrium concentrations. When using 

an optimum growth temperature of 625C, and a maximum Zn flow that was below the 

incorporation limit for substitutional Zn (in order to ensure that the Zn was incorporated 

478

Zn InP Zn 

substitutionally), the diffusion profiles of Zn across the interface could be simulated by 

assuming a concentration-dependent diffusivity. A third-power concentration dependence 

of the effective diffusion coefficient was found. This applied to both Frank-Turnbull and 

kick-out equilibrium mechanisms for an interstitial-substitutional diffusion model. This 

indicated a 2+ charge state for the fast-diffusing Zn interstitials. Extrapolations into the 

high-concentration regime of sealed-ampoule experiments generally agreed with 

published data, although the predominant Zn atoms which were found in the highconcentration 

regime formed complexes with P vacancies in a neutral state. 

S.N.G.Chu, R.A.Logan, M.Geva, N.T.Ha: Journal of Applied Physics, 1995, 78[5], 3001- 

7 

[446-123/124-178] 


A new spin-on solution was proposed for Zn diffusion into this material. The solution 

consisted of a Zn-SiO 2 sol with a very long shelf-life. After spinning-on the sol, a 

ZnO/SiO 2 film was produced on the InP wafer. It was found that, for a thickness of 

150nm, the film acted as an inexhaustible source for short-time (30s, 650C) diffusions. 

U.Schade, B.Unger: Semiconductor Science and Technology, 1993, 8[12], 2048-52 

[446-115/116-145] 


The defects which were introduced by Zn diffusion were studied by measuring the 

photoluminescence and photo-emission spectra of Zn-diffused samples which had been 

fabricated by using a new diffusion technique. The results indicated that Zn diffusion 

generated broad emission bands, with energies ranging from 0.7 to 1eV, only in a surface 

layer with a thickness of less than about 100nm. It also left a P-rich layer with a very high 

Zn concentration and a thickness of less than about 20nm. It was suggested that Zn 

diffusion from a high-Zn concentration source, under P-rich conditions, occurred near to 

the surface and introduced deep centers which were responsible for the bands. 

M.Wada, K.Sakakibara: Japanese Journal of Applied Physics, 1993, 32[2-4A], L469-72 

[446-109/110-042] 


The electrical activity and lattice-site locations of Zn atoms which had been diffused into 

InP were studied by using various characterization techniques. Particle-induced X-ray 

emission channelling showed that, in InP, most of the Zn atoms were situated in 

interstitial sites or formed random Zn precipitates which were electrically inactive. The 

distribution of Zn was shown to depend upon the cooling rate after high-temperature 

diffusion. The difference between the behaviors of Zn in GaAs and InP could be 

understood in terms of the amphoteric native defect model. It was also shown that the 

479

Zn InP Zn 

Fermi level stabilization energy provided a convenient energy reference for the treatment 

of dopant diffusion at semiconductor hetero-interfaces. 

W.Walukiewicz, K.M.Yu, L.Y.Chan, J.Jaklevic, E.E.Haller: Materials Science Forum, 

1992, 83-87, 941-6 

[446-99/100-065] 


A stripe heater was used to diffuse Zn into semi-insulating, or n-type, (001) samples from 

a thin spun-on silica film. The diffusion profiles were determined by means of secondary 

ion mass spectrometry and capacitance-voltage measurements. The diffusion depth and 

activation energy of the effective diffusion coefficient were compared with published data 

on Zn ampoule diffusion. An activation energy of 1.35eV was found for the diffusion of 

Zn in semi-insulating InP. An expression for the effective diffusion coefficient of Zn in 

InP was derived and was compared with the results of a Boltzmann-Matano analysis of 

diffusion profiles. It was concluded that the Zn diffusion could be explained by an 

interstitial-substitutional model, in which Zn diffused as a singly positively charged 

interstitial and acted as an acceptor by filling an In vacancy. 

U.Schade, P.Enders: Semiconductor Science and Technology, 1992, 7[6], 752-7 

[446-99/100-093] 

InP/InGaAs: Zn Diffusion 

It was shown that the growth of emitter layers of InP/InGaAs/InP double heterojunction 

bipolar transistors could result in significant Zn diffusion from the base and into the 

collector. The extent of the diffusion depended upon the n-doping level of the emitter. 

This behavior was explained in terms of non-equilibrium point defects which were 

introduced by a combination of surface pinning of the Fermi level, and n-doping. It was 

also shown that Zn diffusion could be greatly reduced by using AlInAs, instead of InP, as 

the emitter layer. The difference in behavior was shown to be at least partly due to the 

lower diffusivity of group-III interstitials in AlInAs. Moreover, it was shown that the 

introduction of only 50nm of AlInAs between emitter and base resulted in a significant 

reduction of Zn diffusion into the collector. 

R.Bhat, M.A.Koza, J.I.Song, S.A.Schwarz, C.Caneau, W.P.Hong: Applied Physics 

Letters, 1994, 65[3], 338-40 

[446-119/120-218] 


It was recalled that high n + -doping, of the cap layers of heterojunction structures, 

produced anomalous Zn diffusion in the base region during metalorganic vapor phase 

epitaxial growth. This was attributed to non-equilibrium group-III interstitials that were 

generated in the cap layer, and created highly diffusive Zn interstitials via the kick-out 

mechanism. It was shown here that low-temperature (550C) growth was effective in 

reducing the effect of the n + cap layer. Due to a large time constant for the recovery of 

thermal point defect equilibrium, the last-to-grow n + cap layer could not inject excessive 

numbers of group-III interstitials into the base region during growth. However, during 

480

Zn InP Zn 

low-temperature growth the first-to-grow n + sub-collector produced group-III interstitials 

and thus caused anomalous Zn diffusion. In order to prevent this effect, it was suggested 

that growth should be interrupted for 0.5h before growing the base layer, and that growth 

of the n + sub-collector should be carried out at 600C. These changes were effective in 

removing undesirable group-III interstitials. 

K.Kurishima, T.Kobayashi, H.Ito, U.Gösele: Journal of Applied Physics, 1996, 79[8], 

4017-23 

[446-134/135-152] 


It was recalled that highly n + -doped sub-collector layers in such heterojunction bipolar 

transistor structures led to markedly enhanced Zn diffusion in the subsequently grown 

base layer. It was shown that this abnormal Zn diffusion could be suppressed by 

interrupting growth before the Zn-doped layer was grown. It was speculated that this 

interruption of growth permitted excess non-equilibrium group-III self-interstitials, 

coming from the n + -doped sub-collector layer, to disappear before they could enhance Zn 

diffusion in the base layer. 

T.Kobayashi, K.Kurishima, U.Gösele: Applied Physics Letters, 1993, 62[3], 284-5 

[446-106/107-127] 


An investigation was made of the behavior of Zn impurities in heterojunction bipolar 

transistor structures that had been grown by using a low-pressure metalorganic chemical 

vapor deposition technique. In this technique, Zn was anomalously diffused into 

InP/InGaAs heterojunction bipolar transistors with a heavily Si-doped (2 x 10 19 /cm 3 ) subcollector 

layer when the growth temperature before base-layer growth was higher or 

lower than about 600C. On the other hand, an abrupt Zn profile in the same 

heterojunction bipolar transistor structure was obtained when the growth temperature of 

the sub-collector layer was 600C. Two types of point defect reaction, which depended 

strongly upon the growth temperature, were proposed. In one type, excess group-III 

interstitials which were produced in a heavily Si-doped sub-collector layer were easily 

removed. In the other type, the equilibrium concentration of charged group-III vacancies 

decreased as the growth temperature was increased. According to this model, complete 

suppression of unwanted Zn diffusion could be achieved by using a growth interruption 

technique to obtain point defect equilibrium before base-layer growth at 550C when the 

growth temperature of the sub-collector layer was either lower or higher than 600C. It 

was concluded that interruption of the growth for a suitable period of time before baselayer 

growth, or the use of a suitable growth temperature for the sub-collector, was 

essential in order to obtain an abrupt Zn profile in a heterojunction bipolar transistor 

structure with a heavily doped sub-collector layer. 

T.Kobayashi, K.Kurishima, U.Gösele: Journal of Crystal Growth, 1995, 146[1-4], 533-7 

[446-127/128-152] 

481

Zn InP Interdiffusion 

InP/InGaAsP: Zn Diffusion 

In order to prevent carriers in multi-layer heterostructures from being redistributed, a new 

method for open-tube Zn diffusion, using an In-Zn alloy as the source and a 

polycrystalline InP cover to limit surface thermo-damage at temperature as low as 500C, 

was developed. It was found that the diffusion rate was proportional to the square of the P 

content. When using proper masking, the ratio of the lateral width to the diffusion depth 

was about 0.6. 

W.Li, H.Pan: Journal of the Electrochemical Society, 1987, 134[9], 2329-32 

[446-55/56-030] 

General 

InP: Diffusion 

Facet growth near to SiO 2 mask edges, during metalorganic molecular beam epitaxy, was 

studied for various V/III ratios on (100) substrates with a 2°misorientation towards (110). 

It was found that, whereas ideal vertical layer growth occurred at high V/III ratios (even 

after 2µ of growth), oblique (111) planes were kinetically favored near to mask edges at 

lower V/III ratios. The V/III ratio was a key parameter since it determined the facets with 

the lowest kinetically limited growth rate at the border of the growing layer. Also, the 

diffusion length of mobile adsorbed species, which explained the presence of additional 

features near to mask edges and corners, decreased with the V/III ratio. As well as interfacet 

diffusion which was driven by concentration gradients between facets with differing 

growth rates, there was also evidence for the occurrence of anisotropic diffusion along 

[0¯11] on (100) InP. It was suggested that this was the cause of the fine surface ripples 

which were observed on one side, near to the SiO 2 masks. 

R.Matz, H.Heinecke, B.Baur, R.Primig, C.Cremer: Journal of Crystal Growth, 1993, 

127[1-4], 230-6 

[446-106/107-120] 


InP/GaInAsP: Interdiffusion 

Interdiffusion experiments and results for InP/GaInAs(P) heterostructures were 

considered in terms of a thermodynamic model. Important factors which affected 

interdiffusion in the GaInAsP system were shown to include a miscibility gap, differing 

diffusivities on each of the sub-lattices of the 2 materials, Fermi level or impurity-induced 

changes in diffusivity or diffusion mechanism, and the type of experiment. When a 

miscibility gap was present, the activity coefficients and solubilities of all of the species 

varied near to a heterojunction and caused the interdiffusion to become strongly 

composition-dependent. At the usual growth and annealing temperatures, many 

superlattices were expected to equilibrate as 2 quaternary superlattices rather than as an 

homogeneous alloy. Differing diffusivities on the sub-lattices of a superlattice could lead 

482

Interdiffusion InP Interdiffusion 

to widening or narrowing of quantum wells. When this occurred, optical measurements of 

the band-gap energy were likely to be misleading, because of quantum size effects. The 

diffusivity on each sub-lattice could be altered by the presence of group-II, -IV, or -VI 

dopants. Diffusion on the group-III sub-lattice in p-type GaInAsP was found to be 

consistent with an interstitialcy mechanism. The mechanism remained unknown for n- 

type doping and for the group-V sub-lattice. Poorly designed and controlled experiments 

were found to be associated with large discrepancies in the observed diffusivities, with 

unreliable concentration profiles, and with the appearance of new condensed phases in 

the solid. Experiments indicated that the ordered Cu-Pt structure which was often found 

in GaIn(As)P epilayers was unstable, and was not strain-stabilized relative to the 

disordered structure at normally used growth and annealing temperatures. 


[446-106/107-127] 

343 InP/GaP: Interdiffusion 

The formation of a solid solution was monitored by means of X-ray studies of annealed 

powder mixtures of the components. The results (table 43) indicated that the overall 

activation energy for interdiffusion at temperatures of between 650 and 725C was equal 

to 3.15eV. 

U.Voland, R.Cerny, P.Deus, D.Bergner, G.Fenninger: Crystal Research and Technology, 

1989, 24[11], 1177-85 

[446-72/73-040] 

Table 43 

Interdiffusion Parameters for InP/GaP Mixtures 

Temperature (C) D o (cm 2 /s) Q(eV) 

650 - 800 0.007 2.1 

650 - 700 1000 3.2 

675 - 725 400 3.1 

InP/InGaAs: Interdiffusion 

A study was made of the interfacial quality and thermal interdiffusion of quantum wells 

which had been grown using hydride vapor phase epitaxy. It was deduced that island and 

valley structures, with a height of one monolayer and a lateral extent which was of the 

order of one third of an exciton radius, existed at the interface. The interdiffusivity 

coefficient was estimated from the photoluminescence peak energy shift at 77K. Values 

of 2.5 x 10 -19 and 1.5 x 10 -18 cm 2 /s were deduced for temperatures of 700 and 750C, 

respectively. These values were more than 100 times higher than those for AlGaAs/GaAs 

quantum well structures, and more than 100 times smaller than those for InAlAs/InGaAs 

quantum well structures. 

K.Makita, K.Taguti: Superlattices and Microstructures, 1988, 4[1], 101-5 

[446-61-081] 

483

InSb 

Bi 

InSb: Bi Diffusion 

Bulk single crystals of In 1-x Ga x Sb 1-y Bi y (where x was between 0 and 0.21 and y was 

between 0 and 0.005) were grown onto InSb seed crystals by using a rotary Bridgman 

method. The quality of the crystals was assessed by using optical microscopic, X-ray 

topographic, 4-crystal X-ray diffractometric, electron-probe micro-analytical, energydispersive 

spectroscopic, and secondary-ion mass spectroscopic techniques. Due to 

segregation, the compositional ratio of Bi increased as the crystals grew. During the 

growth of InGaSbBi, Bi diffused into the InSb seed, and domains of InBi appeared. For 

comparison, InSb 1-y Bi y (where y was between 0 and 0.05) and In 1-x Ga x Sb (where x was 

between 0 and 0.16) were grown on InSb. It was found that Bi did not diffuse into InSb 

without Ga, whereas Ga diffused without Bi. The incorporation of Ga produced excess In 

and led to the formation of InBi domains. 

Y.Hayakawa, M.Ando, T.Matsuyama, E.Hamakawa, T.Koyama, S.Adachi, K.Takahashi, 

V.G.Lifshits, M.Kumagawa: Journal of Applied Physics, 1994, 76[2], 858-64 

[446-117/118-189] 

Cd 

InSb: Cd Diffusion 

A two-temperature zone method, applied to an InSb substrate plus Cd source system, was 

used. The diffusion profiles were determined by using capacitance-voltage measurements, 

and were similar to those which were habitually found for other III-V systems. Linear 

plots of junction-depth versus the square root of the diffusion time did not pass through 

the origin. This suggested that the usual interstitial-substitutional model and the 

conventional Boltzmann-Matano method could not be used to analyze the results. The 

diffusivity exhibited a maximum as a function of carrier concentration. 

S.L.Tu, K.F.Huang, S.J.Yang: Japanese Journal of Applied Physics, 1990, 29[3], 463-7 

[446-76/77-029] 

484

Ga InSb In 

Ga 

InSb: Ga Diffusion 

Bulk single crystals of In 1-x Ga x Sb 1-y Bi y (where x was between 0 and 0.21 and y was 

between 0 and 0.005) were grown onto InSb seed crystals by using a rotary Bridgman 

method. The quality of the crystals was assessed by using optical microscopic, X-ray 

topographic, 4-crystal X-ray diffractometric, electron-probe micro-analytical, energydispersive 

spectroscopic, and secondary-ion mass spectroscopic techniques. Due to 

segregation, the compositional ratio of Ga decreased, as the crystals grew. During the 

growth of InGaSbBi, Ga diffused into the InSb seed. For comparison, InSb 1-y Bi y (where y 

was between 0 and 0.05) and In 1-x Ga x Sb (where x was between 0 and 0.16) were grown 

on InSb. It was found that Bi did not diffuse into InSb without Ga, whereas Ga diffused 

without Bi. The incorporation of Ga produced excess In and led to the formation of InBi 

domains. 

Y.Hayakawa, M.Ando, T.Matsuyama, E.Hamakawa, T.Koyama, S.Adachi, K.Takahashi, 

V.G.Lifshits, M.Kumagawa: Journal of Applied Physics, 1994, 76[2], 858-64 

[446-117/118-189] 


Permeation of Ga was studied by placing an In-Ga-Sb solution in contact with InSb 

substrates under conditions of zero crystal growth. It was found that the permeation 

distances were equal to 770, 1270 and 2750µ at 380, 430 and 480C, respectively. It was 

deduced that the apparent diffusivity of Ga ranged from 3 x 10 -7 to 5 x 10 -6 cm 2 /s. Rapid 

permeation occurred when the substrate came into contact with the solution. 

M.Kumagawa, H.Ohtsu, E.Hamakawa, T.Koyama, M.Masaki, K.Takahashi, V.G.Lifshits, 

Y.Hayakawa: Materials Science and Engineering B, 1997, 44[1-3], 301-3 


A 10mm-thick InGaSb crystal was grown onto an InSb seed by using the rotary Bridgman 

method. It was found that Ga diffused rapidly into the seed and displaced some of the In. 

The apparent Ga diffusion coefficient was between 10 -8 and 10 -7 cm 2 /s. These values were 

much higher than the self-diffusion coefficients of In or Sb. Rapid diffusion occurred 

only when the substrate was in contact with the solution. The diffusion distance of Ga 

increased upon increasing the holding temperature or time. 

Y.Hayakawa, E.Hamakawa, T.Koyama, M.Kumagawa: Journal of Crystal Growth, 1996, 

163, 220-5 

[446-136/137-125] 

In 

344 InSb: In Diffusion 

Self-diffusion in Bridgman-type single crystals was studied, at temperatures ranging from 

400 to 500C, by using 114m In radiotracers. An anodic oxidation technique was used for 

485

In InSb Li 

serial sectioning, and the penetration profiles were fitted to an erf solution of the diffusion 

equations. It was found that the self-diffusion of In (table 44) could be described by: 

D(cm 2 /s) = 6.0 x 10 -7 [-1.45(eV)/kT] 

The migration enthalpy of In atoms was estimated to be equal to 0.66eV, and the 

corresponding formation enthalpy for an In vacancy was 0.79eV. 

A.Rastogi, K.V.Reddy: Journal of Applied Physics, 1994, 75[10], 4984-9 

[446-117/118-190] 

Table 44 

Diffusivity of 114m In in InSb 

Temperature (C) Diffusivity (cm 2 /s) 

414 1.5 x 10 -17 

435 3.6 x 10 -17 

449 4.8 x 10 -17 

457 5.6 x 10 -17 

477 8.9 x 10 -17 

500 2.4 x 10 -16 

345 InSb: In Grain Boundary Diffusion 

The self-diffusion of In was studied (table 45) in polycrystalline films by using neutron 

activation tracer scanning methods. The grain boundary diffusion parameters were 

evaluated at temperatures ranging from 200 to 400C. The data could be described by: 

D (cm 2 /s) = 1.17 x 10 -6 exp[-0.84(eV)/kT] 

The In diffused via grain boundaries within the temperature range which was studied. The 

grain boundary energy and its temperature dependence was also deduced. 

A.Rastogi, K.V.Reddy: Semiconductor Science and Technology, 1994, 9[11], 2067-72 

[446-119/120-219] 

Li 

InSb: Li Diffusion 

The lithiation of n-type monocrystalline (111)-oriented specimens, via direct reaction 

with n-butyllithium in hexane solution, was carried out at room temperature. This process 

was monitored by using X-ray diffraction and scanning electron microscope techniques. 

Resistivity and Hall coefficient data permitted the change in Li concentration to be 

evaluated. It was deduced that the Li diffusivity at 298K was 1.09 x 10 -8 cm 2 /s. 

G.Herren, N.E.Walsöe de Reca: Solid State Ionics, 1991, 47[1-2], 57-61 

[446-84/85-057] 

486

Li InSb Sb 

Table 45 

Grain Boundary Diffusivity of In in InSb 


218 2.67 x 10 -15 

248 7.89 x 10 -15 

282 2.42 x 10 -14 

302 4.91 x 10 -14 

353 1.78 x 10 -13 

390 4.53 x 10 -13 

413 7.75 x 10 -13 

Sb 

346 InSb: Sb Diffusion 

Self-diffusion in Bridgman-type single crystals was studied, at temperatures ranging from 

400 to 500C, by using 125 Sb radiotracers. An anodic oxidation technique was used for 

serial sectioning, and the penetration profiles were fitted to an erf solution of the diffusion 

equations. It was found that the self-diffusion of Sb (table 46) could be described by: 

D(cm 2 /s) = 5.35 x 10 -4 [-1.91(eV)/kT] 

The migration enthalpy of Sb atoms was estimated to be equal to 0.70eV, and the 

corresponding formation enthalpy for an Sb vacancy was 1.21eV. 

A.Rastogi, K.V.Reddy: Journal of Applied Physics, 1994, 75[10], 4984-9 

[446-117/118-190] 

Table 46 

Diffusivity of 125 Sb in InSb 

Temperature (C) Diffusivity (cm 2 /s) 

408 3.9 x 10 -18 

418 1.2 x 10 -17 

436 1.5 x 10 -17 

455 3.2 x 10 -17 

467 5.3 x 10 -17 

479 8.7 x 10 -17 

347 InSb: Sb Grain Boundary Diffusion 

The self-diffusion of Sb was studied (table 47) in polycrystalline films by using neutron 

activation tracer scanning methods. The grain boundary diffusion parameters were 

evaluated at temperatures ranging from 200 to 400C. The data could be described by: 

D (cm 2 /s) = 1.32 x 10 -4 exp[-1.11(eV)/kT] 

487

Sb InSb Te 

The Sb diffused via grain boundaries within the temperature range which was studied. 

The grain boundary energy and its temperature dependence was also deduced. 

A.Rastogi, K.V.Reddy: Semiconductor Science and Technology, 1994, 9[11], 2067-72 

[446-119/120-219] 

Table 47 

Grain Boundary Diffusivity of Sb in InSb 


218 5.15 x 10 -16 

248 2.41 x 10 -15 

256 3.44 x 10 -15 

282 1.09 x 10 -14 

310 3.34 x 10 -14 

350 1.34 x 10 -13 

Te 

InSb: Te Diffusion 

A PbTe source was used to grow n-type InSb by means of molecular beam epitaxy. Auger 

electron spectroscopic data showed that no surface segregation of Te occurred at doping 

levels of up to about 10 19 /cm 3 . Secondary ion mass spectrometry did not reveal the 

presence of Pb in the films, even at growth temperatures which were as low as 280C. This 

suggested that the Pb rapidly evaporated from the surface during growth. The secondary 

ion mass spectrometric depth profiles for Te revealed signs of solid-state diffusion at 

360C; with a diffusion coefficient of about 10 -13 cm 2 /s. 

D.L.Partin, J.Heremans, C.M.Thrush: Journal of Applied Physics, 1992, 71[5], 2328-32 

[446-86/87-046] 

488

Author Index 

Abe, Y. [440] 

Abernathy, C.R. [269, 294, 444, 458] 

Abrosimova, V.N. [292, 323] 

Acher, O. [439] 

Adachi, A. [422] 

Adachi, S. [484, 485] 

Aers, G.C. [421] 

Ahlborn, K. [464, 466] 

Ahlgren, T. [326] 

Alexandre, F. [398] 

Algora, C. [250, 355] 

Allen, E.L. [234, 239, 248, 307, 

327, 330, 342, 350] 

Alsina, F. [431] 

Alvarez, A.L. [211] 

Alwan, J.J. [242, 248] 

Amann, M.C. [373, 419, 438] 

Amaratunga, G. [318] 

Ambreé, P. [406, 407, 410, 419] 

Anderson, G.B. [248, 396] 

Anderson, S.G. [352] 

Andersson, M. [277, 304] 

Andersson, Y. [277, 304] 

Ando, M. [484, 485] 

Ansaldo, E.J. [372] 

Anthony, J.M. [308] 

Anthony, P.J. [238] 

Appelbaum, A. [474] 

Araujo, D. [356, 368] 

Araujo, G.L. [250, 355] 

Arent, D.J. [382] 

Armiento, C.A. [392] 

Arora, B.M. [428] 

Asahi, H. [420, 450] 

489 

Asashi, H. [433] 

Ashwin, M.J. [344] 

Baba, T. [274] 

Baba-Ali, N. [227, 229, 245, 328, 

390, 394] 

Babievskaya, I.Z. [460, 472] 

Baeumler, M. [426] 

Baiocchi, F. [335] 

Baird, R.J. [412] 

Baker, J.E. [230, 242, 244, 247, 248, 

252, 253, 257, 258, 263, 288, 

338, 339, 437, 448] 

Balk, P. [473] 

Bamba, Y. [236] 

Ban, Y. [263] 

Bandurko, V.V. [272] 

Banerjee, S. [428] 

Bao, X.J. [449] 

Baranowski, J.M. [399] 

Baratte, H. [281, 373] 

Barcz, A.J. [399] 

Baribeau, J.M. [325] 

Barker, R.C. [261] 

Baroni, S. [390] 

Bársony, I. [343] 

Bartels, A. [468] 

Baumann, F.H. [267] 

Baumeister, E. [247] 

Baumeister, H. [465] 

Baur, B. [482] 

Beall, R.B. [236, 246, 248, 280, 282, 

327, 333, 335, 349] 

Beaumont, B. [416] 

Beernink, K.J. [248, 262, 264]

Author Index 

Beggy, J.C. [261] 

Benchimol, J.L. [405] 

Bénière, F. [342] 

Ben-Tzur, M. [359, 473] 

Benz, K.W. [467] 

Ber, B.J. [440] 

Bergner, D. [483] 

Bernholc, J. [343, 347, 366, 367, 

368, 389] 

Besson, J.M. [313] 

Bhat, H.L. [450, 453] 

Bhat, K.N. [363] 

Bhat, R. [406, 415, 420, 480] 

Bhattacharya, P. [274, 405] 

Bhattacharya, P.K. [412] 

Bimberg, D. [297, 332, 352, 446, 447, 

453, 454, 462, 466, 469] 

Bisberg, J.E. [251, 358, 400] 

Biswas, D. [274] 

Bithell, E.G. [392] 

Blaauw, C. [468, 472, 474] 

Blanchard, B. [365, 366, 368, 370] 

Bockstedte, M. [301] 

Boeringer, D.W. [305] 

Boesker, G. [363] 

Bonapasta, A.A. [133] 

Bonkowski, J.E. [448] 

Booker, G.R. [423, 427] 

Borenstein, J.T. [311] 

Borodina, O.M. [313, 410, 464] 

Bösker, G. [298, 363, 364] 

Boudewijn, P.R. [437, 476] 

Boudreau, R. [245, 249, 339, 361] 

Bour, B.P. [262, 264] 

Bour, D. [314] 

Bour, D.P. [399] 

Bowen, T. [245, 249, 339, 361] 

Bradley, I.V. [397, 403, 414, 424, 429] 

Brandt, O. [465] 

Braunstein, G. [331] 

Bravman, J.C. [237, 285, 341, 343] 

Breitenstein, O. [340, 372, 398, 401] 

Brewer, J.H. [372] 

Briggs, A.T.R. [403, 414] 

Brillouet, F. [391, 398] 

Brion, H.G. [464, 466] 

Briones, F. [447] 

Bronner, W. [237] 

490 

Bruk, A.S. [313] 

Brum, J.A. [408, 412] 

Buchanan, M. [421] 

Budko, T.O. [476] 

Bühlmann, H.J. [456] 

Burke, P.T. [395] 

Bürkner, S. [396, 424, 426] 

Burnham, R.D. [234, 243, 244, 245, 

247, 249, 252, 256, 338] 

Burri, G. [368] 

Burton, R.S. [243, 245] 

Busygina, L.A. [440] 

Butherus, A.D. [335] 

Cahen, D. [477] 

Cai, D.Q. [395] 

Calawa, A.R. [388] 

Calle, F. [211, 465] 

Calleja, J.M. [432] 

Camassel, J. [431] 

Caneau, C. [266, 267, 415, 469, 480] 

Cannelli, G. [294] 

Cantelli, R. [294] 

Capizzi, M. [133, 294] 

Cardona, M. [302, 447] 

Cardone, F. [281, 373] 

Carlin, J.F. [456] 

Carr, N. [434] 

Caruso, R. [277] 

Castagné, J. [236, 246, 248, 280, 282, 

327, 333, 335, 349] 

Cerny, R. [483] 

Chan, L.Y. [359, 366, 473, 480] 

Chan, W.K. [406] 

Chand, N. [238] 

Chang, C.C. [406] 

Chang, C.S. [37] 

Chang, J.C.P. [233, 346] 

Chang, K.H. [274, 287, 345] 

Chang, K.J. [289] 

Chang, S.K. [470] 

Chang, T.C. [287] 

Chang, T.Y. [267] 

Chang, Y.A. [400] 

Chaplain, R. [342, 459] 

Charbonneau, S. [421] 

Chase, M.P. [302, 362] 

Chatterjee, B. [463] 

Chatterjee, S. [363]

Author Index 

Chavignon, J. [391] 

Chayahara, A. [319, 324, 325] 

Chen, A.C. [448] 

Chen, B. [343, 347, 389] 

Chen, C.H. [175] 

Chen, C.Y. [364] 

Chen, E.I. [253] 

Chen, H.R. [394] 

Chen, P. [398] 

Chen, S. [331] 

Chen, W.C. [37] 

Chen, Y. [364, 391] 

Chen, Y.C. [405] 

Cheng, C.L. [462] 

Cheng, K.Y. [448] 

Cheng, W.H. [474] 

Cheong, B.H. [289] 

Chester, M.J. [324] 

Chevallier, J. [238, 241, 294, 295, 

310, 312, 313, 314, 454, 460] 

Chi, P.H. [266, 267, 286, 309, 315, 341, 

364, 415, 459, 463, 467, 469] 

Chia, V.F.K. [234, 239, 248] 

Chin, A.K. [251, 355, 358, 400] 

Chiu, T.H. [336, 337, 338, 357, 386] 

Cho, K.I. [377, 422] 

Choa, F.S. [473] 

Choe, B.D. [83, 428] 

Choi, J.S. [470] 

Chow, D.H. [443] 

Chow, K. [372] 

Chow, K.H. [371] 

Christen, J. [332] 

Chu, S.N.G. [478, 479] 

Chui, T.H. [331] 

Cibert, J. [392] 

Clarke, R. [274] 

Clegg, J.B. [236, 246, 248, 280, 282, 

327, 333, 335, 349] 

Clemencon, J.J. [355] 

Cockayne, D.J.H. [395] 

Cockerill, T.M. [242, 248] 

Cohen, D.D. [350] 

Cohen, P.I. [231, 374, 375] 

Cohen, R.M. [17, 298, 315, 364, 

388, 393, 395, 483] 

Colas, E. [228, 254] 

Coldren, L. [250, 359] 

491 

Coldren, L.A. [244, 250, 455] 

Coleman, J.J. [242, 248] 

Conibeer, G.J. [451, 452] 

Corbett, J.W. [311] 

Cordero, F. [294] 

Corzine, S. [244, 250] 

Cox, H.M. [406] 

Cox, S.F.J. [372] 

Craford, M.G. [230, 263, 437] 

Craighead, H.G. [275] 

Cremer, C. [482] 

Crook, A. [242, 248] 

Cunningham, B.T. [256, 262, 288] 

Cunningham, J.E. [331, 336, 337, 338] 

Cunningham, T.J. [261] 

Cutlerywala, H. [364] 

Dabiran, A.M. [231, 374, 375] 

Dabkowski, F.P. [251, 358, 400] 

Dabrowski, J. [304, 342, 344] 

Dadgar, A. [466] 

Dagata, J.A. [324] 

Dallesasse, J.M. [243, 249, 257] 

Dandrea, R.G. [389] 

Danilewsky, A.N. [467] 

Dautremont-Smith, W.C. [294, 461, 464] 

Davies, G.J. [405] 

Däweritz, L. [384] 

De Cremoux, B. [361] 

De Souza, J.P. [281] 

De, T.J. [359] 

Deal, M.D. [234, 237, 239, 248, 282, 

283, 285, 286, 302, 303, 307, 318, 

327, 329, 330, 334, 341, 342, 343, 

350, 362] 

Deicher, M. [290] 

Deppe, D.G. [230, 252, 253, 256, 257, 

258, 263, 338, 339, 360, 437] 

Descouts, B. [391] 

De Temple, T.A. [242, 248] 

Deus, P. [483] 

Dhanasekaran, R. [278] 

Dieckmann, R. [270] 

Dieguez, E. [450] 

Dietrich, H.B. [286, 309, 341, 415, 

459, 463, 467, 469] 

Dildey, F. [419] 

Dion, M. [421] 

Dixon, R.H. [470]

Author Index 

Dodds, S.A. [372] 

Dos Passos, W. [274] 

Dreybrodt, J. [432] 

Dreyer, K. [386] 

Dubon-Chevallier, C. [405] 

Duke, C.B. [389] 

Dunsiger, S.R. [371] 

Dunstan, D.J. [429] 

Dutta, P.S. [450, 453] 

Duvarney, R.C. [372] 

Dzhumakulov, D.T. [321] 

Eastman, L.F. [283] 

Ebbinghaus, G. [419] 

Eberl, K. [302] 

Ebert, P. [448] 

Eccleston, R. [392] 

Eger, D. [468, 474, 477] 

Egger, U. [340, 372, 398, 401] 

Egilsson, T. [317] 

Eizenberg, M. [290] 

Elman, B. [392] 

El-Zein, N. [243, 249, 319] 

Emanuel, M.A. [248] 

Emelyanenko, J.S. [476] 

Emeny, M.T. [429] 

Emmerstorfer, B. [472] 

Emura, S. [433, 450] 

Enders, P. [480] 

Endicott, F.J. [347, 396, 399] 

Endoh, A. [236] 

Enquist, P. [359] 

Epler, J.E. [234, 256, 347] 

Erickson, J.W. [302, 447] 

Erofeeva, E.A. [321] 

Estle, T.L. [314, 371, 372] 

Fahy, M.R. [344] 

Fan, J.C. [287, 345, 427] 

Fareed, R.S.Q. [278] 

Farley, C.W. [308] 

Fathimulla, A. [286, 309, 341, 459, 

463, 467] 

Favennec, P.N. [459] 

Fedotov, A.B. [321] 

Fenninger, G. [483] 

Fewster, P.F. [344] 

Fischer, A. [302] 

Flat, A. [332] 

Fleischman, A.J. [373] 

492 

Fletcher, R.M. [230, 263, 437] 

Fons, P. [457] 

Fonstad, C.G. [406, 413, 414] 

Fontijn, G.M. [435, 436, 437, 438, 

439, 462, 463, 470, 475] 

Forchel, A. [432, 433] 

Fornari, R. [103] 

Forster, T. [382] 

Franke, D. [461] 

Franz, G. [373, 438] 

Franzheld, R. [398, 401] 

Freundlich, A. [348] 

Frova, A. [294] 

Fujii, E. [346] 

Fujii, M. [377] 

Fujii, N. [369] 

Fujii, S. [371] 

Fujii, T. [236] 

Fujii, Y. [324] 

Fujita, K. [340, 377, 426] 

Fukunaga, T. [252, 332] 

Furtado, M.T. [396, 409, 411, 417, 

418, 426] 

Furuya, A. [258] 

Gailhanou, M. [368] 

Gal, M. [395] 

Ganière, J.D. [245, 328, 356, 365, 

366, 368, 370] 

Gao, Y. [391, 398] 

García, S. [465] 

Gauneau, M. [342, 459] 

Gautam, D.K. [290, 365] 

Gautier, S. [404] 

Gavand, M. [294] 

Gavrilovic, P. [244, 247, 258, 338] 

Gaymann, A. [237] 

Genut, M. [290] 

George, T. [348] 

Geva, M. [238, 420, 478, 479] 

Ghaffari, M. [466] 

Giannelis, E.P. [351] 

Giannozzi, P. [311] 

Gibala, R. [274] 

Gibart, P. [241, 310] 

Gibbons, J.F. [360] 

Gilderman, V.K. [269] 

Gillin, W.P. [397, 399, 403, 414, 

424, 429, 430, 432]

Author Index 

Giovine, E. [294] 

Gisdakis, S. [319] 

Gislason, H.P. [293, 317] 

Glade, M. [472, 473] 

Glew, R.W. [431] 

Goldberg, R.D. [421] 

Gonda, S. [433, 450] 

González-Díaz, G. [465] 

Gorelenok, A.T. [440] 

Gösele, U. [237, 241, 252, 261, 277, 

280, 283, 291, 297, 298, 300, 301, 

305, 333, 336, 344, 345, 346, 354, 

360, 367, 372, 398, 401, 461, 481] 

Goto, K. [417] 

Goto, S. [238, 281, 379, 380, 387] 

Govorkov, A.V. [313] 

Grandjean, N. [387] 

Grant, J. [390] 

Grattepain, C. [348, 454] 

Grattepain, C.M. [238, 294] 

Grenet, J.C. [348] 

Griehl, S. [293] 

Griffin, P.B. [318] 

Grigorev, N.N. [361] 

Grodkiewicz, W.H. [335] 

Grote, N. [461] 

Gruen, N. [237] 

Gruenbaum, P.E. [452] 

Grundmann, M. [332] 

Gruska, B. [406, 407, 419] 

Grützmacher, D. [433, 473] 

Gudmundsson, J.T. [317] 

Guersen, R. [1] 

Guido, L.J. [234, 244, 247, 252, 256, 

259, 261, 262, 288, 338] 

Guillaume, J.C. [416] 

Gulwadi, S.M. [415, 469] 

Guoba, L.B. [278, 300] 

Gushchinskaya, E.V. [476] 

Gwilliam, R. [392, 397, 424] 

Gyuro, I. [432] 

Ha, J.S. [386] 

Ha, N.T. [357, 473, 478, 479] 

Haddara, Y.M. [285] 

Hafich, M.J. [265] 

Haga, D. [357] 

Haga, T. [456] 

Hailemariam, E. [266, 267, 405, 414] 

493 

Hallali, P.E. [281] 

Haller, E.E. [297, 302, 359, 366, 

372, 447, 473, 480] 

Hamakawa, E. [484, 485] 

Hamakawa, Y. [292] 

Hamilton, W.J. [443] 

Hamoudi, A. [441] 

Han, H.T. [245, 249, 339, 361] 

Hangleiter, A. [407, 410, 419] 

Hanson, C.M. [359, 473] 

Hara, N. [242, 306, 345, 390] 

Hara, T. [289] 

Harada, Y. [346] 

Harbison, J.P. [227] 

Harde, P. [461] 

Hardingham, C.M. [451, 452] 

Harper, J. [386] 

Harris, G.L. [283] 

Harris, J.J. [236, 246, 248, 280, 282, 

327, 333, 335, 349] 

Harris, J.S. [407, 411] 

Harrison, I. [227, 229, 230, 245, 

328, 390, 394] 

Hart, L. [344] 

Harun-ur Rashid, A.B.M. [351] 

Hashimoto, A. [252, 332] 

Hashimoto, H. [258, 284, 338] 

Haspeklo, H. [415, 475] 

Hata, M. [376, 387] 

Haukson, I.S. [317] 

Hawkins, R.L. [279] 

Hayakawa, Y. [484, 485] 

Haynes, T.E. [342] 

Heime, K. [433] 

Heine, K. [472] 

Heinecke, H. [482] 

Heinz, C. [451, 452] 

Heitz, R. [466] 

Helms, C.R. [340] 

Henini, M. [227, 229, 230, 245, 328] 

Heremans, J. [488] 

Hergeth, J. [473] 

Herms, M. [293] 

Hernandes, C.S. [351] 

Herren, G. [486] 

Herring, C. [314, 315] 

Herzinger, C.M. [242, 248] 

Herzog, L. [340]

Author Index 

Hesse, R. [354] 

Hettwer, H.G. [340, 353, 363, 364, 470] 

Heuken, M. [433] 

Heurtel, A. [312] 

Higashisaka, A. [289] 

Higuchi, M. [471] 

Hill, D.M. [352] 

Hillard, R.J. [449] 

Hillmer, H. [423] 

Hira, K. [319] 

Hirai, K. [323, 324, 325] 

Hirai, M. [340] 

Hitti, B. [371] 

Hiyamizu, S. [422] 

Ho, H.P. [227, 229, 230, 394] 

Hobbs, L. [468, 472] 

Hobson, W.S. [266, 267, 405, 414] 

Hockly, M. [423] 

Hoepfner, A. [319] 

Höfler, G.E. [243, 261, 319] 

Hofsäss, V. [423] 

Holland, O.W. [415, 469] 

Holmgren, D.J. [243, 245] 

Holonyak, N. [230, 234, 243, 244, 247, 

249, 252, 253, 256, 257, 258, 261, 

262, 263, 288, 319, 338, 339, 

360, 437] 

Homewood, K.P. [397, 399, 403, 414, 

424, 429, 432] 

Hong, C.Y. [389] 

Hong, J.M. [275] 

Hong, S.K. [440] 

Hong, W.P. [415, 469, 480] 

Hopkins, L.C. [279] 

Höpner, A. [471] 

Horikoshi, Y. [254, 315, 375, 377, 384] 

Horino, Y. [319, 325] 

Houde-Walter, S.N. [256, 258, 259, 260] 

Houdré, R. [456] 

Hovel, H.J. [373] 

Howard, A.J. [265] 

Howard, L.K. [397, 424, 429] 

Hsieh, K.C. [234, 243, 247, 252, 256, 

257, 258, 261, 319, 320, 360, 448] 

Hsieh, K.Y. [260, 409, 412] 

Hsu, C.C. [443] 

Hsu, J.T. [292] 

Hsu, L. [302, 447] 

494 

Hu, E. [250, 359] 

Hu, E.L. [175, 244, 250] 

Hu, J.C. [285, 286, 303] 

Hu, P. [393] 

Huang, J.H. [267] 

Huang, K.F. [484] 

Huang, L.J. [325] 

Hughes, O.H. [230] 

Hult, M. [277, 304] 

Humer-Hager, T. [288, 369] 

Hutchby, J.A. [359] 

Hwang, D.M. [415, 420] 

Hwang, Y. [355] 

Hwang, Y.L. [409, 412] 

Hyeon, J.Y. [466] 

Ichimura, K. [250] 

Iguchi, H. [233] 

Iikawa, F. [408, 412] 

Ikarashi, N. [274] 

Ikeda, H. [295] 

Ilegems, M. [456] 

Imamura, Y. [420] 

Inai, M. [286, 342] 

Inohara, H. [319] 

Inoue, Y. [299] 

Ishibashi, A. [263] 

Ishii, K. [236] 

Ishikawa, H. [236] 

Ishino, M. [418] 

Isu, T. [376, 380, 387] 

Ito, H. [242, 398, 401, 411, 481] 

Ito, R. [445] 

Ito, T. [255, 378] 

Itoh, H. [242] 

Itoh, M. [435, 441] 

Ivanov, S.V. [235, 281] 

Jach, T. [324] 

Jackman, T.E. [421] 

Jackson, T.N. [373] 

Jacobs, J.M. [438, 475] 

Jaeger, W. [363] 

Jafri, Z.H. [430] 

Jagadish, C. [395] 

Jäger, W. [298, 305, 353, 363, 364, 470] 

Jaklevic, J. [366, 480] 

Jaklevic, J.M. [359, 473] 

Jalil, A. [312] 

James, D.J. [326]

Author Index 

Jan, W. [331, 338] 

Janzén, E. [317] 

Jeanjean, P. [347] 

Jeong, B.S. [442] 

Jeong, W.G. [428] 

Jiménez, J. [103] 

Johnson, M.B. [284] 

Johnson, N.M. [314, 315, 407, 411] 

Jones, K.S. [307, 330, 341, 342, 350] 

Jordan, A.S. [238] 

Joullie, A. [455, 457] 

Joyce, B.A. [231] 

Juhel, M. [441] 

Kadono, R. [371, 372] 

Kahen, K.B. [239, 249, 334, 358, 

360, 367, 393] 

Kakinoki, H. [295] 

Kalem, S. [454] 

Kalish, R. [290] 

Kaliski, R.W. [234, 252, 256] 

Kamanin, A.V. [440, 477] 

Kamei, H. [457] 

Kamei, K. [426] 

Kamejima, T. [307] 

Kameneckas, J.V. [278, 300] 

Kaminska, M. [460] 

Kamiya, T. [431] 

Kanber, H. [283, 291, 320] 

Kaneko, Y. [242] 

Kang, B.K. [439] 

Kang, T.W. [389] 

Kano, H. [232, 387] 

Karenina, L.S. [269] 

Karlicek, R.F. [478] 

Karttunen, V. [312] 

Kasai, K. [306, 345] 

Kasu, M. [232, 380, 385, 386] 

Katahama, H. [426] 

Kataoka, M. [456] 

Kataoka, Y. [236] 

Katayama, M. [299] 

Katayama, Y. [379, 380, 387] 

Katoda, T. [351, 390] 

Katsumata, H. [113] 

Kavanagh, K.L. [316, 322, 346, 348] 

Kawabe, M. [280, 329] 

Kawaguchi, Y. [420] 

Kawai, T. [316] 

495 

Kawamura, Y. [420] 

Kawashima, M. [254, 375, 377, 384] 

Kazarina, N.N. [460, 472] 

Kazmierski, C. [476] 

Kazmierski, K. [361] 

Kehr, K.W. [374] 

Keinonen, J. [312, 326] 

Keller, R. [290] 

Keller, R.C. [340] 

Kempeners, M.N.C. [476] 

Ketata, K. [403, 404, 405] 

Ketata, M. [403, 404, 405] 

Khald, H. [455, 457] 

Khoo, G.S. [382] 

Khreis, O.M. [399] 

Kidoguchi, I. [263] 

Kiefl, R.F. [371, 372] 

Kim, I. [428] 

Kim, J.G. [446] 

Kim, J.H. [439] 

Kim, M.S. [333] 

Kim, S.G. [450] 

Kim, S.K. [389] 

Kim, S.T. [442] 

Kim, T.S. [308] 

Kim, T.W. [389] 

Kim, Y. [333, 395] 

Kim, Y.H. [389] 

Kimura, T. [369, 431] 

Kirchner, P.D. [346] 

Kish, F.A. [243, 247, 249, 258, 339] 

Kishi, M. [351] 

Kishimoto, D. [381, 385, 456] 

Kisielowski, C. [388] 

Kitada, T. [422] 

Klein, P.B. [286, 309, 341, 459, 

463, 467] 

Klem, J.F. [265] 

Klöber, J. [293] 

Klockenbrink, R. [468] 

Knecht, A. [297, 352, 446, 447, 453, 

454, 462, 469] 

Ko, K.Y. [331] 

Kobayashi, J. [284, 338] 

Kobayashi, N. [232, 380, 385, 386] 

Kobayashi, T. [481] 

Koch, M.W. [288] 

Koehler, K. [237]

Author Index 

Koenraad, P.M. [284, 343] 

Kohnke, G.E. [288] 

Kohzu, H. [289] 

Kolbas, R.M. [260, 355, 409, 412] 

Koltsov, G.I. [284, 459] 

Komeno, J. [306, 345] 

Kondo, E. [376] 

Kondo, S. [435, 441] 

König, U. [415, 475] 

Koo, J.Y. [386] 

Kopev, P.S. [235, 281] 

Kopf, R.F. [244, 279] 

Koponen, I. [312] 

Korbutyak, D.N. [293] 

Korobkov, N.N. [292, 323] 

Koshiba, S. [232, 387] 

Koszi, L.A. [461, 464] 

Koteles, E.S. [392] 

Kothiyal, G.P. [412] 

Koumetz, S. [403, 404, 405] 

Koyama, T. [484, 485] 

Koza, M. [415, 420] 

Koza, M.A. [480] 

Kozhukhova, E.A. [410, 464] 

Kozlovskii, V.V. [292, 323] 

Krames, M.R. [253] 

Kraus, G.T. [351] 

Kräutle, H. [297, 352, 446, 447, 453, 

454, 462, 469] 

Krauz, P. [391, 398, 402, 410, 413, 441] 

Kreitzman, S.R. [372] 

Kreller, D. [472] 

Krieghoff, T. [471] 

Kringhøj, P. [469] 

Kristjánsson, S. [317] 

Kubena, R. [393] 

Kucera, J. [321] 

Kudykina, T.A. [361] 

Kuech, T.F. [235, 240] 

Kühn, G. [354, 355, 471] 

Kuhn, J. [423] 

Kuisl, M. [415, 475] 

Kumagai, M. [323, 324] 

Kumagawa, M. [484, 485] 

Kumagaya, M. [319, 325] 

Kumar, V. [453] 

Kuo, C.P. [230, 263, 437] 

Kuo, J.M. [244, 279] 

496 

Kuo, T.Y. [331] 

Kupka, R.K. [391] 

Kurishima, K. [481] 

Kuriyama, Y. [276, 301] 

Kuroda, S. [242] 

Kusano, C. [238, 281] 

Kuttler, M. [466] 

Kuwamoto, H. [474] 

Kuzuhara, M. [307] 

Kwiatkowski, S. [399] 

Kwon, H.K. [83] 

Kwon, O. [439, 440] 

Kwon, Y.S. [251] 

Kwon1, S.D. [83] 

Ky, N.H. [245, 328, 356, 365, 

366, 368, 370] 

Lagally, M.G. [448] 

Lahiri, I. [1, 233] 

Lakdawala, V.K. [292] 

Landgren, G. [253, 317, 353, 466] 

Landheer, D. [325] 

Lanzillotto, A.M. [246, 334] 

Larkins, E.C. [396, 424, 426] 

Laruelle, F. [393] 

Lashkevich, L.N. [293] 

Lau, W.M. [325] 

Launay, F. [361, 439] 

Launay, P. [403, 404, 405] 

Launois, H. [391] 

Lauterbach, C. [469] 

Lawrence, D.J. [239, 251, 357, 358, 400] 

Ledentsov, N.N. [235, 281] 

Lee, C. [333, 342] 

Lee, C.C. [237, 285, 341, 342, 343] 

Lee, C.P. [287, 345, 394, 427] 

Lee, E. [386] 

Lee, E.H. [440] 

Lee, J.C. [235, 240] 

Lee, J.H. [260, 409, 412] 

Lee, J.I. [470] 

Lee, J.K. [440] 

Lee, J.L. [280, 329] 

Lee, J.Y. [440] 

Lee, K.H. [329, 437, 438, 440] 

Lee, S. [386] 

Lee, S.C. [37, 287] 

Lee, S.H. [394, 427] 

Lee, S.T. [251, 261, 305, 331, 357, 400]

Author Index 

Lee, Y.T. [439, 440] 

Lefebvre, F. [404] 

Leitch, A.W.R. [310, 311] 

Lengel, G. [386] 

Lennard, W.N. [325, 477] 

Leon, R.P. [460] 

Leonard, D. [455] 

Lester, S.D. [308] 

Lesunova, R.P. [269] 

Leycuras, A. [348] 

L'Haridon, H. [459] 

Li, E.H. [433] 

Li, W. [482] 

Li, W.M. [315] 

Li, Y.J. [275, 307] 

Liau, Z.L. [465] 

Lichti, R.L. [314, 371, 372] 

Lifshits, V.G. [484, 485] 

Likonen, J. [326] 

Liliental-Weber, Z. [388, 460] 

Lim, H.J. [470] 

Lim, H. [83] 

Lin, H.H. [287] 

Lin, Z. [352] 

Lindeberg, I. [277, 304] 

Linnarsson, M. [317] 

Linnarsson, M.K. [253, 317, 353, 357] 

Liou, D.C. [287, 345] 

Liu, B.D. [287] 

Liu, D.G. [287, 345] 

Liu, H. [446] 

Lo, V.C. [339] 

Lo, Y.C. [260] 

Logan, R.A. [478, 479] 

Logan, R.C. [294] 

Lomasov, V.N. [292, 323] 

Long, J. [462] 

Long, N.J. [423, 427] 

Lopata, J. [294, 449, 461, 464] 

López, M. [379] 

Lord, S.M. [407, 411] 

Lösch, R. [423] 

Loural, M.S.S. [396, 418, 426] 

Ludowise, M.J. [230, 263, 437] 

Ludwig, M.H. [446] 

Luftman, H.S. [266, 267, 279, 357, 

405, 414] 

Luken, K.M. [277] 

497 

Lunardi, L.M. [244] 

Luo, Y.S. [383] 

Lusson, A. [454] 

Lyahovitskaya, V. [477] 

Lyubomirsky, I. [477] 

MacFarlane, A. [371] 

Machayekhi, B.[238, 241, 294, 310, 

314] 

Machayekhi, D. [313] 

Maehara, Y. [426] 

Magee, C.W. [246, 316, 322, 348] 

Magerle, R. [290] 

Maier, M. [237, 396, 424] 

Major, J.S. [256, 262, 288] 

Makarov, V.V. [284, 459] 

Makita, K. [483] 

Malin, J.I. [448] 

Malkovich, R.S. [320] 

Mallard, R.E. [423, 427] 

Malyshev, S.A. [476] 

Mani, H. [455, 457] 

Mannoh, M. [263] 

Marbeuf, A. [294] 

Marcon, J. [403, 404, 405] 

Marek, H.S. [477] 

Marenkin, S.F. [460, 472] 

Marfaing, Y. [312] 

Marschall, P. [415, 475] 

Marti, A. [250, 355] 

Mártil, I. [465] 

Martin, J.M. [266, 267, 465] 

Martin, P. [404] 

Martin, R.W. [434] 

Masaki, M. [485] 

Masseli, K. [473] 

Massies, J. [387] 

Masut, R.A. [408, 412] 

Matsuhata, H. [457] 

Matsui, T. [406, 440, 460] 

Matsui, Y. [418] 

Matsumoto, S. [435, 441] 

Matsumura, N. [456] 

Matsushita, A. [371] 

Matsushita, S. [346] 

Matsuyama, T. [484, 485] 

Matyi, R.J. [247, 360] 

Matz, R. [482] 

Meglicki, Z. [350]

Author Index 

Mehrer, H. [353, 363, 470] 

Mei, P. [227, 228, 415, 420] 

Meier, H.P. [382] 

Melchior, H. [361] 

Melloch, M.R. [1, 233] 

Melman, P. [392] 

Méndez, B. [450] 

Mendonça, C.A.C. [386] 

Menéndez, J. [390] 

Merkulov, A.V. [440] 

Merlin, R. [274] 

Merz, J. [393] 

Merz, J.L. [244, 245, 249, 250, 

339, 359, 361] 

Mestres, N. [274, 432] 

Meyer, J.W. [316, 322, 348] 

Micallef, J. [433] 

Migitaka, M. [295, 296] 

Mihashi, Y. [417] 

Miles, R.H. [443] 

Milnes, A.G. [449] 

Miloche, M. [241, 310, 314, 460] 

Min, S.K. [333] 

Minervini, A.D. [253] 

Mishima, T. [238] 

Mitchell, I.V. [421] 

Mitev, P. [259] 

Mitra, S. [274, 306] 

Mitsuhara, M. [435, 441] 

Mitsui, S. [369] 

Miyauchi, E. [284, 338] 

Miyoshi, S. [445] 

Mizuta, M. [447] 

Mochizuki, A. [289] 

Mochizuki, K. [238, 281, 286, 289, 

377, 422] 

Mochizuki, Y. [447] 

Moise, T.S. [261] 

Mokina, I.A. [440, 477] 

Molinari, E. [390] 

Montie, E.A. [475] 

Moon, D.C. [442] 

Moore, W.J. [279] 

Morgenstern, T. [355] 

Mori, H. [324] 

Mori, M. [242] 

Morishita, Y. [379, 380, 387] 

Morita, T. [284, 338] 

498 

Moriya, N. [290] 

Morrow, R.A. [277, 311, 322] 

Moseley, A. [434] 

Motisuke, P. [408, 412] 

Mui, D.S.L. [245, 249, 339, 361, 455] 

Mukai, K. [395, 432] 

Mukai, S. [242] 

Müller, G. [270] 

Mulpuri, S. [415, 469] 

Murata, M. [431] 

Murray, J.J. [307, 327, 330, 334, 

348, 350] 

Murray, R. [335] 

Murrell, D.L. [416] 

Musbah, O.A. [400] 

Muto, S. [456] 

Nadella, R.K. [266, 267] 

Nagahara, M. [445] 

Nagamine, K. [371] 

Nagamune, Y. [232, 387] 

Nagano, J. [276, 301] 

Nakajima, K. [403, 408, 421] 

Nakamura, M. [296] 

Nakamura, T. [286, 289] 

Nakamura, Y. [232, 250, 387, 401] 

Nakanishi, H. [319] 

Nakano, Y. [290, 365] 

Nakashima, K. [420, 433] 

Nakata, Y. [456] 

Nam, D.W. [244, 247, 256, 338, 360] 

Nam, E.S. [439] 

Näser, A. [466] 

Nashelskii, A.Y. [410, 464] 

Nassibian, A.G. [350] 

Navratil, K. [321] 

Nazar, L. [415, 420] 

Neave, J.H. [231] 

Nekado, Y. [296] 

Nelson, R.W. [459] 

Neugebauer, J. [445] 

Neumann, R. [319] 

Newman, R.C. [322, 335, 344] 

Ng, I. [433] 

Nicholas, R.J. [434] 

Niedermayer, C. [372] 

Niklas, J.R. [293] 

Nilsson, S. [382] 

Nishida, K. [325]

Author Index 

Nishinaga, T. [276, 302, 374, 375, 376, 

377, 380, 381, 385, 422, 

455, 456, 467] 

Nishino, T. [292] 

Nishiyama, K. [371] 

Noack, M. [374] 

Noda, T. [232, 387] 

Noge, H. [232, 387] 

Nolte, D.D. [1, 233] 

Nomura, Y. [379, 380, 387] 

Norcott, M. [281] 

Nordell, N. [253, 317, 353] 

Nörenberg, H. [384] 

Norman, A.G. [427] 

Northrup, J. [344] 

Northrup, J.E. [299, 304, 306, 342, 345] 

Novikova, V.V. [410, 464] 

Nowak, E. [354, 355, 471] 

Nozaki, S. [327, 348] 

Nozaki, T. [307] 

Nukui, T. [242] 

Nutt, H.C. [326] 

O’Reilly, E.P. [396, 424] 

Ogasawara, Y. [316] 

Ogata, H. [406, 460] 

Ogihara, M. [401] 

Oh, J.E. [274] 

Oh, Y.T. [389] 

Ohfuji, S. [276, 301] 

Ohishi, T. [440] 

Ohnaka, K. [263] 

Ohnishi, H. [340] 

Ohno, T. [255] 

Ohori, T. [306, 345] 

Ohsawa, J. [295, 296] 

Ohtsu, H. [485] 

Ohtsuka, K. [440] 

Ohtsuka, K.I. [406, 460] 

Oikawa, H. [289] 

Ojala, P. [253, 317, 353] 

Okamoto, Y. [467] 

Olmsted, B.L. [256, 258, 259, 260] 

Omelyanovskii, E.M. [313, 410, 464] 

Omura, E. [244, 250, 359, 417] 

Onabe, K. [445] 

Ong, C.K. [382] 

Ono, H. [274] 

Osentowski, T.D. [230, 263, 437] 

499 

Osgood, R.M. [321] 

Oshinowo, J. [432, 433] 

Ossart, P. [391, 398, 402, 410, 413] 

Ostling, M. [277, 304] 

Osvald, J. [388] 

Otsuka, N. [418] 

Ougazzaden, A. [441] 

Ourmazd, A. [267] 

Oyanagi, H. [457] 

Paine, B.M. [283] 

Pajot, B. [294, 295] 

Pak, K. [316] 

Pakhomov, A.V. [313, 410, 464] 

Palguev, S.F. [269] 

Palma, A. [379] 

Pan, H. [482] 

Paoli, T.L. [234, 252, 256, 347] 

Parguel, V. [459] 

Park, H.H. [437, 438, 439, 440] 

Park, H.L. [442, 470] 

Park, R.M. [446] 

Parker, S.M. [462] 

Partin, D.L. [488] 

Pascual, J. [431] 

Pashkova, O.N. [460, 472] 

Paska, Z.F. [357] 

Pass, C.J. [234, 239, 248] 

Pavesi, L. [245, 311, 328, 356, 368] 

Pearton, S.J. [63, 266, 267, 269, 277, 

294, 311, 405, 414, 444, 

449, 458, 461, 464] 

Pechman, R.J. [383] 

Peiner, E. [468] 

Perenboom, J.A.A.J. [343] 

Perley, A.P. [266, 267, 405, 414] 

Perrin, S.D. [432] 

Peskov, N.V. [380] 

Peterson, D.L. [239, 393] 

Petrich, G.S. [231, 374, 375] 

Petroff, P.M. [275, 307, 392, 393, 455] 

Pétursson, J. [317] 

Peyre, H. [431] 

Pfeiffer, L. [246] 

Pfeiffer, L.N. [390] 

Pfeiffer, W. [290] 

Pfister, J.C. [366] 

Pfiz, T. [372] 

Phillips, C. [392]

Author Index 

Pilkuhn, M.H. [407, 410, 419] 

Pinzone, C.J. [478] 

Piqueras, J. [450] 

Pisarev, A.A. [272]] 

Planel, R. [347, 391] 

Plano, W.E. [230, 234, 244, 247, 256, 

257, 258, 263, 338, 360, 437] 

Ploog, K. [384] 

Plummer, J.D. [234, 239, 248, 285, 286, 

302, 303, 307, 318, 330, 

350, 362] 

Polyakov, A.J. [313] 

Polyakov, A.Y. [410, 449, 464] 

Ponce, F.A. [347, 396, 399] 

Poole, P.J. [421] 

Poroikov, J.A. [460, 472] 

Potter, T.J. [412] 

Pozsgai, I. [293] 

Prescha, T. [310, 311] 

Primig, R. [482] 

Pross, P. [290] 

Pudenzi, M.A.A. [351] 

Quillec, M. [402, 410, 413] 

Quintana, V. [355] 

Rahbi, R. [238, 241, 294, 314] 

Rai-Choudhury, P. [449] 

Räisänen, J. [312, 326] 

Rajatora, M. [326] 

Rajeswaran, G. [239, 358, 367, 393] 

Ralston, J.D. [396, 424, 426] 

Ramasamy, P. [278] 

Rao, E.V.K. [391, 398, 402, 410, 

413, 441] 

Rao, K. [441] 

Rao, M.V. [266, 267, 286, 309, 341, 

415, 459, 463, 467, 469] 

Rao, P.R.S. [363] 

Rao, S.S. [403, 414] 

Rastogi, A. [486, 487, 488] 

Ravich, V.N. [460, 472] 

Razeghi, M. [439] 

Reddy, K.V. [486, 487, 488] 

Reddy, V. [342] 

Redinbo, G.F. [275] 

Rees, P.K. [326] 

Régrény, A. [342] 

Reinhart, F.K. [245, 328, 356, 365, 

366, 368, 370] 

500 

Ren, H.W. [302] 

Renner, D. [474] 

Rentschler, J.A. [267] 

Reynolds, C.L. [420] 

Reynolds, S. [360] 

Richard, T.A. [243, 249] 

Rihet, Y. [459] 

Ringel, S.A. [463] 

Riseman, T.M. [372] 

Robbins, V.M. [257] 

Robert, J.L. [347] 

Roberts, C. [236, 246, 248, 280, 

327, 349] 

Robinson, D. [270] 

Robinson, H.G.[282, 283, 318, 341, 342] 

Robinson, W. [393] 

Rodionov, A.I. [321] 

Roedel, R.J. [364, 459] 

Roos, G. [315, 407, 411] 

Rose, B. [460] 

Roth, A.P. [408, 412] 

Rothemund, W. [396, 424, 426] 

Rubart, W.S. [307, 330, 350] 

Rucki, A. [353, 363, 364, 470] 

Rudra, A. [456] 

Ruf, T. [447] 

Rytova, N.S. [310] 

Ryu, S.W. [428] 

Sacilotti, M.A. [408, 409, 411, 412, 

417, 418] 

Sacks, R.N. [261] 

Sadana, D.K. [281] 

Sai, N. [394] 

Saito, D. [316] 

Saito, R. [431] 

Sajoto, T. [335] 

Sakaguchi, M. [319, 323, 324] 

Sakai, S. [251] 

Sakaki, H. [232, 387] 

Sakakibara, K. [418, 471, 475, 479] 

Sakalas, A.P. [278, 300] 

Salemink, H.W.M. [284] 

Sallese, J.M. [456] 

Salmeron, M. [233] 

Sandrik, R. [388] 

Sano, N. [422] 

Santos, M. [246, 334, 335] 

Sapriel, J. [391, 398]

Author Index 

Sargünas, V.R. [278, 300] 

Sasa, S. [236] 

Satho, M. [319, 325] 

Sato, E.A. [409, 411, 417, 418] 

Sauer, N.J. [279] 

Scarrott, K. [427] 

Schad, R.G. [281] 

Schade, U. [479, 480] 

Schaff, W.J. [283] 

Schauer, S.N. [459] 

Scheffler, M. [301] 

Schier, M. [419] 

Schlaak, W. [297, 352, 446, 447, 453, 

454, 462, 469] 

Schlachetzki, A. [468] 

Schlapp, W. [423] 

Schlesinger, T.E. [235, 240, 243, 

245, 449] 

Schlotter, N.E. [406] 

Schmidt, H.J. [472] 

Schmidt, M.T. [321] 

Schneider, M. [391] 

Schneider, W. [372] 

Scholz, R. [372, 398] 

Schowalter, L.J. [381] 

Schubert, E.F. [244, 246, 279, 336, 

337, 478] 

Schultz, M. [277, 372, 398, 401] 

Schulz, K.J. [400] 

Schumann, B. [354, 355, 471] 

Schumann, D. [371] 

Schumann, H. [466] 

Schützendübe, P. [384] 

Schwab, C. [314, 371, 372] 

Schwartz, B. [335] 

Schwartz, C.L. [228, 415, 420] 

Schwarz, S.A. [227, 228, 415, 420, 480] 

Schweizer, H. [423] 

Scilla, G.J. [373] 

Scott, E.G. [405, 423] 

Sealy, B.J. [429, 467] 

Sekiguchi, Y. [418, 471, 475] 

Seko, M. [418, 475] 

Sellitto, P. [347] 

Semprini, E. [379] 

Serreze, H.B. [477] 

Seshadri, S. [259, 261] 

Seta, M. [450] 

501 

Shacham-Diamand, Y. [351] 

Shah, D.M. [406] 

Shahar, A. [254] 

Shapira, Y. [352] 

Sharma, V.K.M. [451, 452] 

Shayegan, M. [246, 334, 335] 

Sheets, J. [316, 322, 348] 

Shen, X.Q. [375, 380, 381, 385, 

455, 456] 

Shepherd, F.R. [468, 474, 477] 

Shi, S.S. [245, 249, 339, 361] 

Shiba, Y. [426] 

Shichijo, H. [247, 360] 

Shieh, T.H. [287] 

Shiina, K. [306, 345] 

Shimogaki, Y. [365] 

Shimomura, S. [422] 

Shinoda, A. [286, 342] 

Shintani, Y. [251] 

Shiomi, A. [323, 324] 

Shiraishi, K. [255, 376, 378] 

Shiraki, Y. [445] 

Shitara, T. [231, 376, 377, 422] 

Shiu, W.C. [433] 

Shmidt, N.M. [440, 477] 

Sicart, J. [347] 

Siefert, R.L. [383] 

Siegel, W. [293] 

Siethoff, H. [464, 466] 

Silveira, J.P. [447] 

Simes, R. [393] 

Simons, D.S. [266, 267, 286, 309, 315, 

341, 364, 415, 459, 463, 467, 469] 

Sin, Y.K. [355] 

Singh, J. [274, 405] 

Singh, S. [335] 

Skoryatina, E.A. [320] 

Skromme, B.J. [415, 420] 

Skudlik, H. [290] 

Slotte, J. [326] 

Smirnov, V.M. [272] 

Smith, A.D. [403, 414] 

Smith, F.T. [251, 357, 400] 

Smith, R.S. [326] 

Smith, S.C. [243, 245, 249] 

Sobotta, H. [354] 

Södervall, U. [298, 364] 

Solmon, H. [270]

Author Index 

Somogyi, K. [241, 293, 310] 

Song, J.I. [480] 

Soni, R.K. [450] 

Sonoda, T. [369] 

Speier, P. [432] 

Spence, J.P. [367] 

Spencer, M.G. [283] 

Spiller, G.D.T. [405] 

Springthorpe, A.J. [472, 477] 

Srivastava, A.K. [428] 

Stähl, K. [277, 304] 

Stam, M. [449] 

Stark, J.B. [336] 

Stavola, M. [461, 464] 

Steckl, A.J. [398] 

Sternitzke, M. [270] 

Stevenson, D.A. [282, 318, 327, 329, 

334, 437, 438] 

Stillman, G.E. [256, 262, 288] 

Stobbs, W.M. [392] 

Stollenwerk, M. [433] 

Stolwijk, N.A. [298, 353, 363, 364, 470] 

Storz, F.G. [386] 

Streetman, B.G. [308] 

Stutius, W. [258] 

Stutzmann, M. [454] 

Suehiro, H. [242] 

Sugawara, M. [395, 432] 

Sugg, A.R. [243, 247, 249, 256, 258] 

Sugimoto, H. [440] 

Sugiura, H. [435, 441] 

Sun, D. [262, 264] 

Sun, J.Z. [339] 

Sundaram, V.S. [452] 

Suzuki, T. [374, 376] 

Swaminathan, V. [238, 277, 420, 

461, 464] 

Swanson, M.L. [477] 

Swart, J.W. [351] 

Syfosse, G. [313] 

Szafranek, I. [262] 

Szafranek, M. [262] 

Tada, K. [290, 365] 

Tadayon, B. [283] 

Tadayon, S. [283] 

Taguti, K. [483] 

Takahashi, K. [484, 485] 

Takahashi, S. [417] 

502 

Takamiya, S. [369] 

Takamori, A. [258, 284, 338] 

Takamori, T. [284, 338] 

Takano, H. [319, 323, 324, 325] 

Takebe, T. [286, 342, 377] 

Takeda, A. [364] 

Takeda, S. [289] 

Takiguchi, T. [417] 

Takizawa, J. [450] 

Talamo, A. [379] 

Tamura, S. [319, 325] 

Tan, T.Y. [237, 241, 252, 261, 277, 

283, 291, 297, 298, 300, 301, 303, 

305, 331, 333, 336, 344, 345, 346, 

360, 367, 372, 401, 461] 

Tanaka, M. [302, 374, 376] 

Tang, T.K. [242, 248] 

Tanigawa, S. [280, 329, 371] 

Taninaka, M. [401] 

Tanoue, H. [450] 

Tasker, P.J. [283] 

Taylor, M. [423] 

Taylor, S. [456] 

Taylor, S.J. [416] 

Tejwani, M.J. [283] 

Tell, B. [336, 337, 338] 

Terada, S. [346] 

Terauchi, Y. [251] 

Tews, H. [288, 319, 369] 

Theys, B. [191, 238, 241, 294, 310, 

313, 314, 454, 460] 

Thibierge, H. [398, 402, 410, 413, 441] 

Thijs, P.J.A. [435, 436, 439] 

Thomas, L.M. [292] 

Thompson, J. [434] 

Thornton, R.L. [234, 248, 252, 256, 

396, 399] 

Thrush, C.M. [488] 

Thrush, E.J. [427] 

Thundat, T. [381] 

Tian, Q. [271] 

Tjaden, D.L.A. [437, 438, 475, 476] 

Tokuda, Y. [299] 

Tokumitsu, E. [357] 

Tomassini, N. [379] 

Tomioka, T. [236] 

Tomita, N. [422] 

Tomlinson, W.J. [254]

Author Index 

Towers, M. [326] 

Trafas, B.M. [383] 

Tramontana, J.C. [399] 

Treat, D. [399] 

Treat, D.W. [262, 264] 

Treichler, R. [247, 288, 369, 419, 465] 

Trequattrini, F. [294] 

Tsai, C.M. [427] 

Tsai, K.L. [394, 427] 

Tsang, J.S. [394, 427] 

Tsang, W.T. [386, 473] 

Tsu, R. [305] 

Tsuchida, N. [296] 

Tsuchiya, M. [232, 275, 307, 387] 

Tsugami, M. [369] 

Tsushi, A. [250] 

Tu, C.W. [244, 294] 

Tu, S.L. [484] 

Tuck, B. [227, 229, 230, 245, 328, 390] 

Tulchinsky, D.A. [311] 

Tweet, D.J. [457] 

Uekusa, S. [113] 

Uematsu, M. [277, 280, 281, 287, 

354, 398, 401] 

Uesugi, F. [417] 

Ullrich, H. [297, 352, 446, 447, 453, 

454, 462, 469] 

Umeno, M. [348] 

Unger, B. [479] 

Urban, K. [363, 448, 470] 

Usami, A. [299] 

Uskov, W.A. [321] 

Vaccaro, P. [340] 

Van Berlo, W.H. [253, 317, 353, 466] 

Van de Walle, C.G. [445] 

Van de Wijgert, W.M. [435, 436, 439] 

Van der Stadt, A.F.W. [343] 

Van der Vleuten, W.C. [284] 

Van Dongen, T. [438, 475] 

Van Gieson, E. [382] 

Van Gurp, G.J. [435, 436, 437, 438, 

439, 475, 476] 

Van Uitert, L.G. [335] 

Varava, A.V. [272] 

Vaudry, C. [459] 

Vawter, G.A. [244, 250, 359] 

Venkatesan, T. [227, 228, 415, 420] 

Vesely, E.J. [256] 

503 

Veuhoff, E. [247, 465] 

Vicek, J.C. [406, 413, 414] 

Vieu, C. [391] 

Virkar, A.V. [271] 

Vitkauskas, A.A. [278, 300]] 

Viturro, R.E. [259] 

Voland, U. [483] 

Völkl, J. [464, 466] 

Vook, D.W. [360] 

Vriezema, C.J. [463, 470] 

Wada, K. [280, 281, 287, 354] 

Wada, M. [418, 471, 475, 479] 

Wada, N. [251] 

Wada, O. [258] 

Wada, T. [299, 364] 

Wager, J.F. [275, 300] 

Wagner, J. [426] 

Wake, D. [405] 

Wakejima, A. [422] 

Walker, J. [315] 

Walsöe de Reca, N.E. [486] 

Walter, W. [382] 

Walters, R.J. [125] 

Walukiewicz, W. [359, 366, 473, 480] 

Wandel, K. [407, 410, 419] 

Wang, C. [366, 367, 368, 389] 

Wang, E.G. [448] 

Wang, K.W. [462] 

Wang, L. [297, 302, 447] 

Wang, X.S. [383] 

Warren, A.C. [346] 

Watanabe, A. [376] 

Watanabe, M. [242] 

Watanabe, N. [252, 332] 

Watanabe, T. [286, 340, 342, 377] 

Weaver, J.H. [352, 383] 

Webb, R.P. [429] 

Weber, E.R. [233, 348, 460] 

Weber, J. [310, 311] 

Wei, L. [280, 329] 

Weill, G. [313] 

Weimer, M. [386] 

Wenzl, H. [374] 

Werner, J. [361] 

Werner, P. [277, 372, 398, 401] 

West, K.W. [246, 390] 

Weyer, G. [469] 

Wheeler, C.B. [459]

Author Index 

Whelan, J.M. [283, 291, 320] 

Whitlow, H.J. [277, 304] 

Whitney, P.S. [406, 413, 414] 

Wichert, T. [290] 

Wicks, G.W. [259, 288] 

Wielsch, U. [407, 419] 

Wilkie, J.H. [467] 

Willén, B. [357] 

Williams, J. [258] 

Williams, M.D. [386] 

Williams, P. [459] 

Willoughby, A.F.W. [451, 452] 

Wilson, R.G. [63, 269, 444, 449, 458] 

Wisser, J. [472] 

Wittorf, D. [470] 

Wolf, H. [290] 

Wolfram, P. [461] 

Wolk, J.A. [447] 

Wolter, J.H. [284, 343] 

Wong, S.L. [434] 

Wood, A. [434] 

Wood, C.E.C. [283, 288] 

Woodall, J.M. [1, 233, 346] 

Woodbridge, K. [236, 246, 248, 280, 

327, 349] 

Wu, A.T. [348] 

Wu, C.H. [243, 319, 320] 

Wu, M.Y. [287] 

Wu, X.S. [244, 250, 359] 

Wu, Z. [321] 

Wurzinger, P. [288, 369] 

Xia, H. [325] 

Xu, F. [352] 

Xu, J. [394] 

Xu, Z. [394] 

Yaguchi, H. [445] 

Yajima, H. [242] 

Yakimenko, I.J. [440] 

Yakobson, B. [389] 

Yakobson, S.V. [410, 464] 

Yamaguchi, H. [254, 315, 375, 377, 384] 

Yamamoto, T. [286, 342, 377] 

Yamamura, S. [431] 

Yamazaki, S. [395, 432] 

Yang, B. [317] 

Yang, B.H. [293, 317] 

Yang, E.S. [443] 

Yang, K. [381] 

504 

Yang, S.J. [484] 

Yang, X. [394] 

Yang, Y.F. [443] 

Yang, Y.N. [383] 

Yesis, L. [335] 

Yi, J.Y. [386] 

Yokota, K. [319, 323, 324, 325] 

Yokoyama, N. [456] 

Yoneda, M. [250] 

Yonezu, H. [316] 

Yoo, H.J. [251] 

Yoon, H.W. [227] 

Yoon, I.T. [442] 

You, H.M. [261, 305, 344, 345, 346] 

Young, E.W.A. [462, 463, 470] 

Yu, C.F. [321] 

Yu, K.M. [359, 366, 460, 473, 480] 

Yu, S. [237, 241, 252, 291, 297, 

298, 305, 333] 

Yu, S.J. [433] 

Yuan, S. [395] 

Yugo, S. [431] 

Yurchuk, S.J. [284, 459] 

Yurre, T.A. [440] 

Yutani, A. [325] 

Zahraman, K. [241, 310] 

Zalm, P.C. [463, 470] 

Zavada, J.M. [63, 269, 444, 458] 

Zehr, S.W. [474] 

Zhang, K. [405] 

Zhang, P. [394] 

Zhang, Q. [389] 

Zhang, Q.M. [343, 347, 366, 367, 368] 

Zhang, S. [276] 

Zhang, S.B. [299, 306, 345] 

Zhang, T. [355] 

Zhdanov, S.K. [272] 

Zheng, B. [394] 

Zheng, J.F. [233] 

Zheng, L.R. [331] 

Zielinski, E. [432] 

Zimmermann, H. [367, 461] 

Zolper, J.C. [63, 265] 

Zou, J. [395] 

Zou, W.X.[244, 245, 249, 250, 339, 361] 

Zrenner, A. [335] 

Zucker, E.P. [252] 

Zundel, T. [310]

Author Index 

Zwicknagl, P. [288, 369] 

Zydzik, G.J. [335] 

Zypman, F. [305] 

Zytkiewicz, Z.R. [372] 

505

Keywords (abstracts) 

2-dimensional [230, 231, 232, 343, 

374, 375, 376, 377, 384, 385, 

386, 392, 421, 422] 

3-dimensional [316, 386, 392, 455] 

adatoms [254, 375, 376, 378, 381, 

382, 445] 

amorphization [297, 307, 329, 341, 349, 

446, 449, 453, 462] 

amorphous [316, 319, 322, 323, 

324, 325, 341, 347, 453, 466] 

amphoteric [244, 247, 333, 358, 366, 

473, 479] 

anelastic [294] 

anions [385, 414, 433, 436, 437, 438] 

anisotropy [232, 234, 256, 374, 

384, 385, 482] 

annealing [224, 227, 231, 233, 234, 235, 

236, 240, 242, 243, 245, 252, 253, 

256, 257, 258, 259, 260, 261, 263, 

264, 265, 266, 267, 268, 271, 274, 

275, 276, 277, 278, 280, 281, 282, 

283, 284, 285, 288, 289, 290, 291, 

292, 296, 299, 300, 301, 304, 305, 

306, 307, 309, 312, 316, 317, 318, 

319, 320, 321, 324, 325, 327, 329, 

331, 332, 334, 336, 337, 338, 340, 

341, 342, 343, 344, 346, 348, 349, 

351, 352, 353, 357, 359, 360, 362, 

363, 365, 366, 369, 373, 374, 388, 

389, 390, 391, 392, 393, 394, 395, 

396, 397, 398, 401, 402, 403, 404, 

405, 407, 408, 409, 410, 412, 413, 

414, 415, 417, 420, 421, 422, 424, 

426, 427, 428, 430, 431, 432, 433, 

434, 437, 438, 439, 444, 446, 447, 

448, 449, 450, 453, 454, 456, 458, 

459, 461, 462, 463, 464, 465, 467, 

468, 469, 470, 471, 472, 474, 475, 

482, 483] 

annihilation [287] 

anodic [364, 485, 487] 

anodization [395] 

anomaly [331, 332] 

antisites [264, 265, 296, 350, 355, 396, 

[424, 450] 

arrays [470] 

band-edge [244, 247, 338, 395, 424] 

band-gap [238, 241, 260, 294, 312, 

338, 340, 395, 424, 426, 483] 

bands[316, 317, 340, 391, 447, 465, 479] 

barrier [231, 243, 249, 274, 276, 

304, 311, 342, 344, 384, 386, 397, 

406, 413, 418, 427, 430, 433, 445, 

448, 460, 463, 474] 

bombardment [278, 300, 351, 383] 

boundaries[224, 270, 276, 286, 301, 325, 

387, 440, 486, 488] 

capless [234, 256, 257, 281, 288, 

291, 320] 

carrier [236, 243, 249, 250, 251, 258, 

287, 288, 289, 294, 295, 299, 306, 

319, 320, 326, 338, 339, 340, 345, 

346, 350, 358, 368, 388, 393, 400, 

417, 433, 451, 452, 454, 462, 470, 

471, 473, 476, 478, 484] 

cathodoluminescence [259, 274, 275, 

306, 332, 382, 392, 401, 450] 

507

cation [232, 233, 254, 385, 386, 414, 

419, 427, 433, 436, 437, 438, 439] 

channelling [325, 341, 366, 479] 

Czochralski [285, 293, 322, 331, 339] 

dangling-bond [386] 

deep-level [293, 307, 312, 316, 

340, 396, 424, 449, 459, 460] 

defect-free [443] 

deposition [231, 247, 251, 254, 261, 

276, 286, 288, 289, 301, 324, 332, 

335, 348, 357, 372, 373, 374, 375, 

377, 380, 383, 386, 395, 400, 420, 

421, 424, 426, 442, 468, 471, 472, 

473, 474, 481] 

diffusion-controlled [290] 

diffusion-induced [253, 283, 300, 340, 

363, 364, 368, 369, 370, 426] 

diode [250, 311, 449] 

dislocations [224, 230, 274, 293, 306, 

332, 341, 342, 344, 348, 360, 363, 

399, 426, 440, 470] 

dislocation-free [251] 

di-vacancy [414] 

donor-acceptor [442, 447, 468, 472, 474] 

donors [238, 241, 249, 257, 258, 

259, 289, 294, 295, 297, 298, 303, 

304, 306, 307, 308, 310, 311, 312, 

313, 314, 315, 328, 329, 330, 333, 

338, 340, 344, 345, 346, 347, 350, 

360, 369, 370, 371,388, 398, 406, 

410, 437, 438, 442, 447, 449, 459, 

460, 461, 464, 465, 468, 470, 471, 

472, 474, 475, 476] 

DX [246, 334, 347] 

EL2 [226, 276, 277, 293, 311, 373, 448] 

electromigration [278, 372, 421] 

epilayer [236, 286, 287, 316, 329, 

344, 471] 

epitaxy [227, 232, 234, 237, 239, 

244, 248, 250, 253, 274, 281, 283, 

286, 287, 289, 315, 316, 317, 326, 

329, 335, 337, 345, 346, 347, 348, 

353, 355, 357, 368, 369, 376, 377, 

379, 380, 381, 382, 384, 387, 389, 

395, 406, 407, 408, 411, 413, 416, 

417, 418, 419, 421, 422, 423, 424, 

427, 436, 438, 447, 452, 453, 455, 

456, 457, 462, 468, 478, 482, 483, 488] 

etching [253, 276, 301, 317, 324, 

343, 353, 365, 372, 373, 386, 402, 

409, 413, 455, 472, 475] 

exciton [244, 247, 287, 338, 392, 

441, 483] 

faults [287, 345, 440] 

Fickian [403, 414] 

films [234, 239, 242, 244, 248, 250, 

251, 276, 290, 301, 316, 319, 320, 

322, 324, 334, 335, 347, 354, 355, 

357, 359, 373, 381, 395, 400, 401, 

415, 419, 424, 438, 440, 445, 450, 

451, 469, 475, 486, 487, 488] 

first-principles [343, 347, 378, 

386, 389, 444, 445] 

fluence [278, 392, 449, 450] 

free-carrier [454] 

Frenkel [245, 275, 300, 327] 

hetero-epitaxial [348, 445, 463] 

hetero-interface [250, 258, 274, 316, 

348, 406, 413, 440, 468] 

heterojunction [237, 238, 254, 286, 

359, 383, 410, 419, 420, 440, 442, 

480, 481, 482] 

heterostructure [234, 238, 252, 256, 

262, 287, 388, 392, 420, 426, 427, 

438, 439] 

homo-epitaxial [381, 463] 

hysteresis [302] 

in-diffusion [234, 237, 239, 248, 252, 

258, 267, 276, 277, 280, 281, 286, 

291, 292, 294, 297, 309, 321, 339, 

341, 346, 354, 363, 366, 367, 389, 

398, 408, 411, 415, 416, 419, 459, 

461, 463, 464, 465, 467, 469] 

interstitial [230, 236, 237, 238, 250, 253, 

278, 282, 283, 284, 285, 289, 300, 

301, 303, 304, 312, 317, 318, 339, 

342, 344, 353, 355, 357, 359, 360, 

361, 362, 364, 366, 367, 368, 369, 

370, 371, 373, 386, 387, 405, 406, 

407, 409, 411, 416, 417, 418, 419, 

437, 438, 442, 460, 462, 470, 471, 

472, 473, 474, 475, 476, 478, 479, 480] 

interstitialcy [261, 298, 434, 483] 

interstitial-substitutional [227, 229, 250, 

508

261, 297, 355, 361, 366, 401, 406, 

407, 419, 437, 438, 451, 452, 460, 

473, 475, 476, 478, 479, 480, 484] 

interstitial-type [470] 

islands [335, 383, 384, 426, 455] 

lattice-mismatch [389, 408, 412, 463] 

light-emitting [250] 

Lindhard-Scharff-Schiott [334] 

line-width [395, 421] 

luminescence [235, 240, 274, 281, 292, 

340, 356, 390, 396, 424, 427, 447, 

450, 454, 470, 474] 

metalorganic [231, 232, 247, 253, 261, 

286, 288, 289, 306, 317, 332, 345, 

348, 353, 357, 380, 385, 386, 387, 

408, 411, 416, 417, 418, 419, 420, 

426, 427, 428, 430, 435, 440, 462, 

465, 466, 468, 470, 471, 472, 473, 

474, 478, 480, 481, 482] 

microcathodoluminescence [312] 

microcrystalline [321] 

mid-gap [276] 

migration [231, 236, 244, 245, 246, 248, 

253, 254, 255, 262, 264, 267, 269, 

280, 283, 285, 291, 292, 301, 304, 

312, 314, 315, 317, 320, 334, 335, 

336, 340, 342, 344, 350, 353, 355, 

374, 375, 377, 378, 380, 381, 382, 

383, 385, 388, 402, 409, 410, 411, 

413, 414, 417, 418, 422, 442, 445, 

446, 448, 455, 459, 464, 471, 474, 

486, 487] 

misfit [236, 426, 439, 440] 

mismatch [290, 420, 436, 443, 468] 

misorientation [231, 237, 281, 374, 

375, 376, 386, 422, 482] 

mobility [231, 241, 246, 283, 294, 295, 

310, 343, 372, 373, 374, 375, 381, 

421, 468] 

models [256, 278, 321, 336, 360, 367, 

404, 467, 476] 

moiré [426] 

monolayer [231, 232, 287, 315, 345, 374, 

375, 386, 483] 

Monte Carlo [231, 239, 249, 254, 289, 

341, 380, 385, 392] 

Mössbauer [468] 

multi-layers [267, 370, 372, 377, 

383, 448, 482] 

muons [370, 371, 372] 

muonium [313, 314, 370, 371, 372] 

native [230, 243, 245, 249, 257, 

258, 259, 263, 276, 281, 298, 299, 

301, 306, 321, 322, 345, 358, 366, 

437, 444, 445, 448, 452, 473, 479] 

negative-U [311, 445] 

neutron [320, 486, 487] 

non-equilibrium[260, 279, 280, 285, 305, 

331, 354, 365, 366, 367, 389, 404, 

440, 480, 481] 

n-type [233, 238, 250, 251, 253, 257, 

258, 266, 287, 294, 295, 298, 299, 

303, 305, 310, 311, 312, 313, 314, 

316, 317, 319, 321, 322, 339, 340, 

344, 345, 346, 355, 357, 363, 365, 

371, 373, 388, 393, 394, 400, 410, 

415, 429, 437, 438, 442, 444, 452, 

460, 461, 463, 464, 467, 468, 469, 

474, 475, 476, 480, 483, 486, 488] 

offsets [242, 247] 

organometallic [228, 254, 283, 315, 416, 

457, 478] 

out-diffusion [237, 252, 266, 267, 268, 

274, 281, 286, 287, 290, 291, 297, 

299, 305, 307, 344, 345, 346, 363, 

369, 398, 405, 415, 435, 441, 443, 

444, 450, 458, 462, 466, 469, 470, 

471, 472, 474] 

oxidation [243, 249, 321, 324, 340, 

485, 487] 

passivated [310, 311, 312, 407, 411, 

449, 461, 464] 

passivation [241, 263, 310, 312, 407, 

411, 444, 447, 449, 452, 453, 454, 463] 

patterned [242, 247, 254, 377, 382, 445] 

p-doped [253, 300, 311, 317, 353] 

Pendellösung [443] 

permeation [485] 

photo-electron [299, 324] 

photo-emission [276, 301, 351, 479] 

photo-induced [292] 

photoluminescence [236, 239, 245, 248, 

253, 258, 259, 260, 262, 264, 277, 

283, 286, 287, 288, 292, 293, 316, 

509

317, 322, 327, 340, 350, 356, 365, 

366, 368, 369, 370, 389, 390, 391, 

392, 394, 395, 396, 403, 407, 408, 

409, 411, 412, 414, 417, 421, 422, 

424, 425, 426, 427, 428, 429, 430, 

431, 432, 433, 435, 436, 439, 441, 

442, 448, 450, 456, 465, 466, 470, 

474, 479, 483] 

photonic [259] 

Poole [421] 

positron [280, 329] 

pre-exponential [275, 299, 300, 464, 466] 

protons [292, 312, 322, 323] 

p-type [233, 237, 238, 243, 253, 257, 

263, 264, 265, 267, 279, 280, 283, 

287, 288, 289, 294, 297, 298, 299, 

300, 303, 305, 316, 317, 319, 321, 

322, 328, 340, 345, 354, 358, 360, 

363, 365, 373, 388, 393, 394, 403, 

406, 410, 413, 414, 415, 418, 420, 

429, 444, 450, 460, 461, 463, 464, 

467, 472, 473, 474, 483] 

quantum-well [234, 239, 243, 249, 252, 

253, 254, 256, 260, 262, 274, 287, 

306, 369, 370, 391, 392, 393, 395, 

396, 399, 422, 425, 426, 427, 430, 

431, 433, 434, 435, 436, 439, 441, 456] 

radiation-enhanced [272] 

radiotracers [485, 487] 

Raman [283, 306, 345, 350, 351, 368, 

390, 391, 426, 431, 432, 433, 449] 

roughening [464] 

roughness [380, 384] 

self-interstitials [237, 240, 252, 277, 279, 

280, 283, 291, 297, 298, 300, 301, 

303, 333, 354, 367, 368, 394, 404, 

461, 470, 481] 

simulation [231, 239, 249, 277, 278, 

282, 285, 286, 289, 301, 303, 318, 

339, 342, 343, 353, 360, 362, 367, 

378, 380, 385, 386, 388, 392, 398, 

403, 404, 407, 419, 422, 427, 461, 479] 

sintering [270, 361] 

stoichiometric [296, 446, 453, 462] 

superlattices [227, 228, 229, 230, 232, 

233, 234, 235, 237, 239, 240, 241, 

243, 244, 247, 248, 249, 252, 255, 

256, 257, 258, 259, 261, 283, 287, 

297, 300, 302, 305, 315, 335, 343, 

345, 347, 360, 366, 367, 388, 389, 

390, 391, 393, 394, 397, 398, 401, 

415, 419, 420, 427, , 433, 439, 443, 

447, 448, 482] 

temperature-dependent [239, 293, 316, 

321, 351, 430, 451] 

tracer [374, 461, 486, 487] 

transients [284, 285, 292, 293, 307, 312, 

314, 316, 317, 318, 373, 398, 449, 

459, 463] 

tunnelling [231, 232, 233, 284, 376, 380, 

382, 383, 384, 385, 386, 448] 

twinning [224, 274, 306] 

ultra-thin [390, 456] 

undersaturation [283, 297, 300, 301, 

346, 365] 

up-hill [282, 285, 317, 318, 339] 

vacancies [230, 239, 243, 245, 248, 

249, 252, 253, 256, 257, 258, 259, 

260, 263, 275, 276, 277, 280, 282, 

285, 286, 291, 292, 294,297, 298, 

299, 300, 301, 303, 304, 305, 307, 

308, 315, 316, 317, 318, 325, 326, 

327, 328, 329, 330, 331, 333, 334, 

336, 338, 339, 340, 341, 342, 343, 

344, 345, 346, 350, 351, 356, 357, 

359, 360, 361, 366, 367, 368, 369, 

370, 374, 383,385, 388, 389, 390, 

392, 393, 394, 396, 397, 398, 399, 

404, 406, 407, 410, 414, 415, 419, 

421, 424, 428, 429, 430, 437, 451, 

460, 470, 471, 475, 478, 479, 480, 

481, 486, 487] 

wafers [239, 277, 284, 292, 295, 

318, 321, 338, 358, 361, 354, 364, 

420, 428,464, 465, 471, 473, 

476, 479] 

X-ray [274, 276, 277, 278, 286, 290, 

296, 299, 300, 301, 304, 324, 325, 

343, 348, 351, 352, 358, 366, 368, 

388, 399, 417, 420, 422, 423, 427, 

428, 435, 436, 439, 443, 446, 447, 

453, 454, 461, 469, 473, 477, 479, 

483, 484, 485, 486] 

510

Diffusion in GaAs and other III-V Semiconductors - Stephen J. Pearton

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