Damage formation and annealing studies of low energy ion implants ...

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collision cascade and are mobile at room temperature, in combination with the operation of dynamic annealing. The migration and trapping of these point defects at the oxide interface is the origin of the formation of the amorphous, near-surface damage layer. A similar process in which the surface represents a trap capturing interstitials (of which there is an excess over vacancies (27)), can also explain the formation of the damage layer observed for the case of the low dose, low energy As implants. In this early stage amorphisation will be nucleation limited (14, 16). At a later stage in the process, the planar growth of the amorphous layer into the underlying bulk material is observed. This phenomenon is due not only to the accumulation of mobile interstitials originating from greater depths at the advancing amorphous/crystalline interface, but also to the expansion of disordered regions due to the increasing density of coverage of the Si crystal matrix by the growing number of collision cascades that ultimately overlap. The type of planar growth seen here at room temperature has also been observed at elevated temperatures in studies of amorphisation / recrystallisation of the Si matrix under higher energy ion bombardment (11, 14-16). Similar processes to those described above were considered to be responsible for the planar growth. Considering next the development of the As profile and its correlation with the displaced Si distribution, Figure 5.2 shows that at doses below 1 × 10 14 cm -2 the implanted As profile is significantly shallower than the TRIM calculated one which is also drawn in the figure. However for increasing doses the As profile peak shifts to greater depths, from 3.6 nm at 4 × 10 14 cm -2 to 5.5 nm for 1.8 × 10 15 cm -2 , as indicated by the arrows in the figure. At the same time the As profile broadens towards the surface, as is clearly visible for doses of 4 × 10 14 cm -2 and above. The latter behaviour can be explained on the basis of range shortening due to the changing composition of the matrix through As accumulation, collisional redistribution and / or collective mass flows to take account of the accommodation of the implanted ions. Sputtering can be largely ignored, as a TRIM calculation shows that 2.5 keV implanted to a dose of 1.8 × 10 15 cm -2 causes the removal of less than 0.3 nm oxide. The incident ions will of course cause disorder to the native oxide which remains unnoticed by MEIS. Figure 5.2 clearly shows that the movement of As to greater depth with increasing dose occurs in parallel with the planar growth of the amorphous layer. This indicates that the two effects are closely related. On the basis of ballistics arguments, there is no doubt that a substantial fraction of the As is initially stopped beyond the amorphous /crystalline interface. Since MEIS does not detect any clear presence of As at these depths, the possibility of the dopant having moved into substitutional positions and therefore being 111
no longer “visible” in MEIS has to be considered. This was checked in a control experiment using SIMS in which all As present is detected. Figure 5.3 shows the SIMS depth profiles, all taken under identical sputter erosion conditions, for As implanted to a dose of 8 × 10 13 cm -2 into a virgin Si crystalline sample and to doses of 5 × 10 13 cm -2 and 1 × 10 14 cm -2 into a self-ion pre-amorphised Si sample. Although for shallow implants SIMS may give rise to small peak shifts, these will be the same in all these cases. The TRIM calculated As profile is again included for reference. The approximate peak positions for the two cases are indicated by arrows. It is clear that there is a good agreement between the As implant and the TRIM calculated profile, as would be expected. Significantly however, the SIMS results shows a clear shift “inwards” by ~1 nm of the maximum of the distribution, for the implant into the crystalline compared to the ones into the pre-amorphised Si sample. This corroborates the MEIS evidence and confirms that the shallower As peak observed for low implant doses for which the Si matrix is not yet fully amorphised is real. The effect of ion beam mixing limited resolution of SIMS is likely to cause the tail on the SIMS profiles in Figure 5.3 beyond the TRIM profile. As Concentration (cm -3 ) 1E20 1E19 SIMS 2.5 KeV As implant into crystalline and amorphous Si Xtal Amorph 1E18 0 5 10 15 20 25 Depth (nm) 112 As 5E13 amorph As 1E14 amorph As 8E13 x-tal As TRIM Figure 5.3 SIMS depth profiles taken under identical sputter erosion conditions, for As implanted to a dose of 8 × 10 13 cm -2 into a virgin Si crystalline sample and to doses of 5 × 10 13 cm -2 and 1 × 10 14 cm -2 into a self-ion pre-amorphised Si sample. The TRIM calculated As profile is shown for comparison 0.01 1E-3 TRIM (A.U.)
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Damage formation and annealing stud
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3.2.2.6 Other models 40 3.2.3 Other
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Chapter 7 Interaction between Xe, F
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List of Figures Figure 1.1 a) Schem
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Figure 4.13 Variation in the kinema
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Figure 6.10 MEIS energy spectra for
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Figure 7.8 Combined MEIS Xe depth p
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Abbreviations and Symbols a/c amorp
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Abstract The work described in this
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Chapter 7 5 M. Werner, J.A. van den
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terminal (Vg), current cannot flow
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This is an approximate average leve
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the active channel, adjacent to the
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To continue to improve devices ther
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produces a device quality regrown l
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technique of channelling Rutherford
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22 J.S Williams. Solid Phase Recrys
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and the probability of scattering t
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importance for many atomic collisio
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M1, V0, E0 Figure 2.2 Elastic scatt
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2.3.1 Models for inelastic energy l
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dE/dx (ev/Ang) 10 1 Inelastic Energ
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dE/dx (eV/Ang) 125 100 75 50 25 0 2
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Figure 2.5 Results of TRIM simulati
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Chapter 3 Damage and Annealing proc
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the Si/SiO2 interface, consuming th
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On the basis that by creating an in
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Figure 3.4 Structure of crystalline
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a Si atom will suffer little angula
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3.2.2.5 Homogeneous model (Critical
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Sputtering and atomic mixing play a
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and is approximately 25 times faste
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elevant dopants later. For equal co
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nearest neighbour distance (52). By
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Category I defects are produced whe
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thermal annealing (600 - 700 °C an
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Figure 3.11 Relationship between im
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defect pairs due to Coulomb attract
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⎛ 〈 C ⎞ ⎛ ⎞ I 〉 〈 C V
Page 79 and 80: 27 R.D. Goldberg, J. S. Williams, a
Page 81 and 82: 67 H. Bracht. Diffusion Mechanism a
Page 83 and 84: Hall effect measurements were carri
Page 85 and 86: energy than one scattered from an a
Page 87 and 88: epresents a small improvement over
Page 89 and 90: (dE/dx)out multiplied by the path l
Page 91 and 92: they are small compared to the diff
Page 93 and 94: ackscattering (27). This fact forms
Page 95 and 96: Figure 4.7 a) Plot of a Gaussian di
Page 97 and 98: similar to the width of the error f
Page 99 and 100: UP Ion Beam SPIN Rotation Sample Sc
Page 101 and 102: Kinematic factor (K) 1.0 0.8 0.6 0.
Page 103 and 104: Figure 4.14 Illustration of the dou
Page 105 and 106: 4.2.2.4 Interpretation of spectra A
Page 107 and 108: with are comparatively small, ~ 0.5
Page 109 and 110: Inelastic energy loss (eV/Ang) 32 2
Page 111 and 112: iterative procedure is carried out
Page 113 and 114: Yield (couts per 5µC) 300 250 200
Page 115 and 116: SIMS experiments were also carried
Page 117 and 118: MEIS, using the scattering conditio
Page 119 and 120: 4.5 Sample production Samples have
Page 121 and 122: an N2/O2 environment to maintain an
Page 123 and 124: 38 M. Anderle, M. Barozzi, M. Bersa
Page 125 and 126: damage evolution behaviour observed
Page 127 and 128: Yield (counts per 5 µC) 250 200 15
Page 129: essentially a “zero dose” profi
Page 133 and 134: yield (cts / 5µC) 500 400 300 200
Page 135 and 136: 5.4 Conclusion MEIS analysis with a
Page 137 and 138: Chapter 6 Annealing studies 6.1 Int
Page 139 and 140: 6.2.2.2 Results and Discussion Figu
Page 141 and 142: theory predictions and X-ray fluore
Page 143 and 144: implantation conditions are those u
Page 145 and 146: a) b) c) Yield (counts per 5 µC) Y
Page 147 and 148: greater than MEIS. SIMS is not sens
Page 149 and 150: attributed to the interference betw
Page 151 and 152: The as-implanted sample, with a bro
Page 153 and 154: a) b) Yield (counts per 5 µC) Yiel
Page 155 and 156: interface, as evidenced by the high
Page 157 and 158: duration, is observed. MEIS results
Page 159 and 160: Yield (counts per 5µC) 500 400 300
Page 161 and 162: ack edges of the Si peaks are very
Page 163 and 164: underneath the SiO2 layer, iii) it
Page 165 and 166: R s (Ω/sq) 950 900 850 800 750 60
Page 167 and 168: As concentration (at/cm 3 ) 1E22 1E
Page 169 and 170: R s (Ω/sq) 950 900 850 800 750 70
Page 171 and 172: Following annealing it was observed
Page 173 and 174: ∆a/a (x 10 -3 ) 4,0 epi550 3,5 3,
Page 175 and 176: Yield (counts per 5 uC) 350 300 250
Page 177 and 178: (FWHM). Concomitantly, As in the re
Page 179 and 180: Yield (counts per 5 µC) 450 400 35
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The higher temperature anneals carr
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ecomes steeper for the sample annea
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Figure 6.28 Schematic illustrations
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and the 2D picture in Figure 6.31b)
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6.5 Conclusion In summary, in this
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20 L. Capello, T. H. Metzger, M. We
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egarding B profiles relevant to the
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Yield (counts per 5 µC) 400 300 20
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TRIM AU 0.04 0.03 0.02 0.01 TRIM si
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a) F profile PAI 3 keV BF2 b) F pro
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Yield (counts per 5 µC) 20 15 10 5
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the corresponding PAI sample, yield
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amorphous matrix, (16) i.e. local c
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21 M. Anderle, M. Bersani, D. Giube
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stopped at depths beyond the observ
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the role of each individual element
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