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Editorial Panel
Paolo ANTONA
Materials Consultant.
Pietro APPENDINO
Professor, Material Technology and Applied Chemistry; Faculty of Engineering,
Turin Polytechnic.
Marcello BADIALI
R&D Director, Teksid Aluminum.
Giuseppe CAGLIOTI
Professor, Solid State Physics ; Faculty of Engineering, Milan Polytechnic.
Enrico EVANGELISTA
Professor, Metallurgy; Faculty of Engineering, Ancona Polytechnic.
Luca Paolo FERRONI
Marketing Director, Teksid Aluminum.
Merton C. FLEMINGS
Head Department of Materials Science and Engineering, M.I.T., Cambridge, Mass.
Sergio GALLO
Invited Professor, Applied Chemistry & Metallurgy; Faculty of Engineering, Turin
Polytechnic. Past President Teksid S.p.A.
Cinzia MODENA
Editorial Coordinator, Metallurgical Science & Technology, Teksid Aluminum.
Claudio MUS
Materials Consultant.
Walter NICODEMI
Professor, Siderurgy & Siderurgical Technology; Faculty of Engineering, Milan Polytechnic.
Editor in chief: Marcello BADIALI
Registrazione presso il Tribunale di Torino n. 3298 del 12 maggio 1983
Associato alla Unione Stampa Periodica Italiana
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the approval of the Editorial Panel.

E. Gariboldi, D. Ripamonti,
L. Signorelli, G. Vimercati,
F. Casaro
S. Seifeddine, T. Sjögren,
I. L. Svensson
M. El Mehtedi, L. Balloni,
S. Spigarelli, E. Evangelista,
G. Rosen, B.H. Lee, C.S. Lee
Metallurgical
Science and Technology
A journal published
by Teksid Aluminum
twice a year
Vol. 25 No. 1, July 2007
FRACTURE TOUGHNESS AND
MICROSTRUCTURE IN AA 2XXX
ALUMINIUM ALLOYS
VARIATIONS IN MICROSTRUCTURE AND
MECHANICAL PROPERTIES OF CAST
ALUMINIUM EN AC 43100 ALLOY
COMPARATIVE STUDY OF HIGH
TEMPERATURE WORKABILITY OF ZM21
AND AZ31 MAGNESIUM ALLOYS
PAGE
3
12
23

EDITORIAL
The present issue marks an important goal for Teksid Aluminum: the 25th
anniversary of Metallurgical Science and Technology (MS&T). Twenty five years
of life represent a remarkable milestone which is not commonly reached
by a scientific journal. MS&T is one of the few specialized editorial products
edited by an industrial company. It has been always delivered world wide,
free of charge and with continuity.
This moment deserves a word on the nature of the Journal, on its intrinsic,
permanent features as well as on its evolution and future perspectives.
The birth of MS&T originated from the initiative of professor Sergio Gallo,
former President of Teksid S.p.A., together with full professors Aurelio
Burdese of the Politecnico di Torino and Walter Nicodemi of the Politecnico
di Milano. They suggested to earmark part of the financial resources destined
to the image promotion activities of the company for a specialized scientific
review. Their proposal was warmly agreed by Antonio Mosconi, the then
Managing Director of Teksid, who, in the first issue’s introduction, wrote:
“The aim of Metallurgical Science and Technology is to initiate a dialogue,
already part of the scene in other countries, between all those
resources that are committed to the search for new frontiers in
metallurgical techniques, so as to derive the greatest synergies from
them.
Metallurgical Science and Technology then enters the scene as an
instrument whereby the dialogue between metallurgical science and
metalworking industry can be taken to greater depths.
‘Academic’ metallurgy indeed, has often shown signs of hankering
after pure science, estranged from industry. It should not be forgotten
that its final objective remains practical application.
A journal such as this is not as ambitious as to imaging that it can
entirely fill the goal. Its aim is rather to inculcate the habit of
comparison and so cast a permanent bridge between the two parties.
Its very existence shows that the need to set up a more systematic
and continuous two-way flow of information is seen by both sides as
an indispensable prelude to its development.”
Since these targets and perspectives were indicated, several changes have
occurred in the life of Teksid Group, the sponsor and editor of the journal.
The original name of the company, Teksid, evokes a core business centered
on cast iron foundry. Founded in 1978 as a spin-off of Fiat metallurgical
activities, its experience rooted in the very beginning of the Italian industrial
development which started-up in the early century. Steel and cast-iron
production were then the company’s core business focus. The global trend
toward “energy conservation” triggered by the recurrent oil crises of the
seventies and eighties induced Teksid to share and favor the efforts aiming
at vehicle weight reduction: the interest of the Company then gradually
Metallurgical Science and Technology
shifted to light metal alloys – like magnesium and,
above all, aluminum – as base materials. Teksid was
then split into divisions depending on the material
their industrial activity was dedicated to (iron,
aluminum and magnesium). In 2002 the aluminum
division was sold to the private equity market to
give birth to Teksid Aluminum, today’s publisher
of MS&T.
It is known that today’s market dynamics, especially
in the car production industry, is much and
dramatically faster than in the past years, and Teksid
Aluminum had to react to the changing scenario
through a severe restructuring and divestiture
process, which largely resized its traditional
footprint of an international and industry leading
company.
The evolution experienced by Teksid Aluminum
and the radical changes occurred during the last
decades in science, in technology and consequently
in our society, have left the objectives of the journal
unchanged, and its contents have evolved
coherently with the scientific and technologic
trends of metallurgy.
In the Company strategy, these changes have been
accompanied by a constant search for innovative
and competitive product solutions, by a constant
effort to enable the costumers enjoy the company’
s collaboration, from the concept stage to the
delivery of unfinished or completely machined
products and by a constant attention the best
combination of technical performance and
product quality.
This evolution has been reflected by MS&T too.
This is not the place for presenting an exhaustive
list of the topics covered by the journal in all
these years. However probably some readers
will remember academic and technological
contributions dedicated to innovative forming
techniques – such as lost-foam or semisolid
forming, – or to the challenges derived by the
emergence of new structural materials such as cast

magnesium alloys, or to thermomechanical
treatments – like liquid hot isostatic pressing, –
and to new experimental methods for testing metal
properties and performance – thermoelastoplastic
yield stress vs. conventional yield stress, and much
more.
Since its birth, this journal has hosted contributions
of both academic research and metalworking
industry from all over the world. Perhaps the
‘academicians’ have cooperated more
enthusiastically than the industrial researchers to
the “two-way flow of information” originally
wished, but these attitudes is to be attributed both
to the confidential nature of industrial research
and to the old academic caveat “publish or perish”
still alive in the web era..
In all these years Teksid Aluminum has treasured the MS&T bridging role
between science and industrial expertise in the R&D programs of the Group.
Furthermore the metallurgic community could enjoy a 25-year long support
to scientific divulgation. We can assert that the targets, originally indicated
by Mr. Mosconi, have been reached. Nowadays MS&T papers published since
1983 can be found in university libraries, academic institutions and industrial
R&D laboratories all over the world. Notably, all the issues published since
the year 2000 can be also retrieved comfortably from the Company website
(www.teksidaluminum.com).
The topics such as those mentioned above and unpredictable new ones –
specifically related to light metals and alloys - will continue to arouse interest
of contributors and readers of MS&T for many years to come, as the
increasing cooperation and subscriptions testify.
In a joyful anniversary like this one, we cannot conclude this editorial other
than by congratulating with Metallurgical Science & Technology for its longevity
and scientific maturity, and by wishing Happy Birthday and a long life!
Metallurgical Science and Technology
Editorial Panel

FRACTURE TOUGHNESS AND
MICROSTRUCTURE IN AA 2XXX
ALUMINIUM ALLOYS
1 E. Gariboldi, 1 D. Ripamonti, 1 L. Signorelli, 1 G. Vimercati, 2 F. Casaro
1 Politecnico di Milano, Dipartimento di Meccanica, Milano (MI)
2 Varian Vacuum Technologies, Leinì (TO)
Abstract
The paper presents the toughness properties of forgings made of two AA 2xxx series
aluminium alloys with different microstructural conditions. Fracture toughness tests in
crack opening mode I were performed on compact tension specimens machined from
the forgings in different orientations. The tests were performed both at room temperature
and at 130°C.
Fracture toughness properties were related to microstructural and fractographic features
of the alloys in order to discuss on their failure mechanisms. The effect of the coarse
intermetallic phases within grains or at their boundaries in the different conditions was
underlined. The testing temperature, within the range here investigated, neither affected
fracture toughness properties nor failure mechanisms.
KEYWORDS
Al alloy AA2014, Al alloy AA2618, fracture toughness, microstructure.
3 - Metallurgical Science and Technology
Riassunto
Il lavoro illustra le proprietà di tenacità alla frattura misurate
in forgiati di grandi dimensioni realizzati con due leghe di
alluminio della serie 2xxx in differenti condizioni
microstrutturali. Dai forgiati sono stati ricavati provini CT
con differenti orientazioni sui quali sono state condotte
prove di tenacità a frattura secondo il modo I di apertura
della cricca. Le proprietà ottenute sono state correlate
con la microstruttura riscontrata nei campioni e completate
con analisi frattografiche atte ad individuare i meccanismi
di cedimento. È stato così messo in luce l’effetto delle
diverse microstrutture ed in particolare delle particelle
grossolane di fasi intermetalliche presenti a bordo grano
o all’interno dei grani che differenziano i forgiati nelle leghe
esaminate di composizione più complessa.

1. INTRODUCTION
Aluminium-alloy forging is currently used to
manufacture structural components of relatively
large and complex shape. The plastic deformation
imparted to the material can positively affect its
microstructure by promoting recrystallization
cycles and a greater homogeneity of alloying
elements. However, it should be considered that
in large size forgings, the relatively low amount of
plastic strain given to the alloy cannot completely
refine the structure and intermetallic particles as
in other small-size wrought products such as
extruded bars or rolled sheets. In addition, the
slower quenching rates experienced by large
forgings result in lower mechanical properties
achieved after the subsequent aging process. Large
differences in cooling rate between surface and
centre of large forgings during solution annealing
also result in remarkable residual stresses, that are
often relieved by inserting a plastic deformation
step after quenching and before the aging
treatment [1]. This method also modifies the
precipitation sequence and kinetics of the alloy.
The above described effects significantly affect the
tensile and fracture properties of aluminium alloy
forgings. Focusing the attention on the fracture of
aluminium alloys, it was reported that toughness
is strictly related to the presence of coarse
particles, 0.1 to 10 µm in diameter, that can be
either non-equilibrium particles formed during
2. MATERIALS INVESTIGATED
The present investigation was carried out on three
forgings having a roughly cylindrical shape with a
diameter of 250 mm, made of aluminium
alloys AA2014 (Al4CuSiMg) and AA2618
(Al2Cu1.5MgNi). The parts had been forged from
extruded bars of diameter 190 mm with different
manufacturing cycles.
Two forged samples of the AA2014 alloy were
produced by forging in two steps at 390°C. The
samples were then heat treated to T6 temper by
different parameters. A first sample, hereafter
referred to as 2014-A forging was solution
annealed at 505°C for 6 hours, water quenched
and artificially aged at 160°C for 14 hours, following
the usual industrial heat treatment route. In the
case of the forging in AA2618 alloy (hereafter
referred to as 2618 forging), the solution annealing
at the usual temperature for this material, 530°C,
lasted 1 hour, and it was followed by water
quenching and by artificial aging at 190°C for 20
solidification or inclusions from insoluble impurities [2]. These particles
crack easily as the matrix deforms within the plastic flow zone at the crack
tip and causes the typical ductile fracture mode where crack propagates
via coalescence of voids. The amount, size and distribution of these second
phase particles is thus relevant for the fracture toughness properties and
even material with comparable tensile properties can display significantly
different fracture toughness properties.
In addition to the abovementioned effect, the role played by submicrometer
particles (0.01 to 0.5 µm in diameter) need to be considered. In the case of
the same volume fraction and microstructural features of coarse particles,
a substantial modification of toughness can be observed in age hardenable
aluminium alloys varying the amount and characteristics of fine hardening
particles. The behaviour is complex and the fine hardening particles are
responsible for it. It is well known that the presence of particle having
suitable distribution and size enhances the resistance of peak aged alloys to
deformation and thus tends to reduce the extension of the plastic zone,
positively affecting toughness [2, 3]. On the other hand, the lower strainhardening
capacity of the material in the peak aged condition with respect
to the underaged condition, also gives rise to local plastic instabilities that
significantly contribute to reduce the material toughness [2]. Further, where
grain boundary precipitate free zones are observed, strain localization in
these regions and intergranular ductile fracture can occur [3, 4]. In these
cases the fracture toughness depends on the spacing and size of the voidnucleating
particles at grain boundaries [4, 5]. The higher fracture toughness
displayed by alloys in the underaged with respect to over-aged condition as
well as the transition towards intergranular ductile fracture mode as
overaging proceeds confirm this latter effect [2, 4].
The aim of the present paper is to contribute to a better understanding of
the correlation between microstructure, tensile and toughness properties
of aluminium forgings as a result of different thermomechanical cycles,
focusing in particular to the role of large second phase particles.
hours, following a common industrial practice. The 2014-B forging was
produced and heat treated as for 2618 material, leading to a substantially
overaged matrix and to intermetallic phases distribution rather different
than that of 2014-A forging.
Light optical microscopy observations and Vickers hardness tests (0.98N
load) were performed to evaluate the general microstructural features.
Tensile test specimens were machined from the forgings in the hoop direction
in regions characterized by a homogeneous structure and hardness.
Two sets of Compact Tension (CT) fracture toughness specimens were
machined with cracks laying in diametral planes of the forgings. In the first
set, the crack propagation direction was radial (CR direction according to
ASTM E399-90 and B645-02 [6, 7]) while in the second set it was longitudinal
(CL direction). The specimens had a thickness B of 20 mm [6, 7].
Tests were carried out on a MTS 810 universal testing machine, equipped
with an environmental chamber suitable for test temperatures up to 250°C.
Before fracture toughness tests, a precracking stage was performed at test
temperature until a total crack length (machined notch and fatigue crack)
of about 20 mm was reached. Precracking was performed under load control
(sinusoidal cycles at 10 Hz frequency) while crack length was monitored
via the elastic compliance technique by measuring the Crack Opening
Displacement (COD). The stress intensity factor during pre-cracking
decreased linearly with crack length from 11 to 8 MPa•m 1/2 . Fracture
4 - Metallurgical Science and Technology

toughness tests were carried out imposing a displacement rate of 0,025
mm/s, monitoring the COD and applied load until the occurrence of unstable
crack propagation.
In a second stage of the research study, the toughness of the materials at
130°C was investigated. The J IC integral was measured according to the
single-specimen technique and the ASTM E1820-01 standard [8]. The J IC
parameter was evaluated instead of K IC since the material toughness was
expected to be greater than that measured at room temperature. The tests
were performed monitoring crack length via the compliance method at
fixed steps of COD. In order to reduce the strain accumulated under load
during each load step due to creep effects at 130°C, the holding period
under constant applied load (P) for the adjustment of the crack length was
fixed to 1 s. J Q was estimated according to the mentioned standard as the
intercept between the power-law curve (fitting the experimental data within
the range stated by ASTM E1820-01) and the straight line passing to the
3. RESULTS
3.1 MICROSTRUCTURE
All forgings were characterized by grains elongated along the main plastic
flow path experienced during forging. Coarse intermetallic particles were
also present in the microstructure, as expected for these alloys. More
specifically, the 2014-A material (figure 1a) was characterized by elongated
grains with a transversal size of about 50 µm and by large intermetallic
particles aligned in the flow direction: globular Al 2 Cu (θ) particles (bright
particles in figure 1a) and blocky shaped clustered particles containing Fe,
Mn, Si and Cu (darker particles in the same figure). In regions where these
particles were observed small equiaxed grains were often detected
A B
C D
5 - Metallurgical Science and Technology
point ∆a = 0.2 mm, P = 0 N with a slope
corresponding to a double flow stress (this latter
corresponding to the average between the yield
and the ultimate tensile stress).
Validity requirements of J 1C tests were not met in
all cases since unstable crack propagation or popin
phenomena occurred in some samples. In these
cases the K Q parameter was evaluated from the
P-COD curves.
Fracture surfaces were observed by scanning
electron microscopy (SEM). Metallographic
sections were also cut perpendicularly to the crack
plane to observe the crack propagation path by
optical microscopy and SEM.
suggesting the local recrystallization effect induced
by these intermetallics.
The direction of plastic flow was less evident in
the 2014-B samples, that displayed rather elongated
grains of 50-100 µm in transverse size together
with smaller equiaxed grains in the regions
containing large intermetallic particles (figure 1b).
These were in lower amount with respect to the
previous alloy. Particles with dark appearance at
grain boundaries (figure 1b) were of complex
chemical composition. On the contrary globular
Al 2 Cu particles had a bright appearance in
Fig. 1:
Typical microstructure of
the investigated forgings.
2014-A (a), 2014-B (b),
2618 at low (c) and
higher (d) magnification.
The longitudinal axis of
forgings is horizontal in
all the micrographs.

micrographs and were observed both inter- and
intragranularly.
The 2618 sample was characterized by very large
grains slightly elongated in the flow direction. Their
size in the transverse direction was of the order
of 500 µm, as shown in figure 1c. As expected,
considering the composition of the AA2618 alloy
[1], several intragranular FeNiAl 9 particles were
observed together with globular Al 2 Cu and (CuFe)
Al 3 precipitates.
The microstructure and hardness of the forgings
were in all cases homogeneous in the regions
where tension and toughness specimens were
sampled.
3.2 MECHANICAL BEHAVIOUR
The average tensile properties in hoop direction
of the investigated materials at room temperature
and at 130°C are summarized in Table I. At room
temperature, the 2014-A forging showed the
highest tensile strength and the lowest ductility.
The increase in test temperature from 20 to 130°C
led to a significant reduction in the ultimate tensile
strength (UTS) and 0,2% proof stress (YS) and to
a corresponding significant increase in reduction
of area (Z) and elongation (A) at fracture.
The results of the fracture toughness tests are
summarized in Table 2 while representative load
vs. COD graphs are shown in Figure 2. Here, the
lines needed for the evaluation of K Q are added
A) B)
6 - Metallurgical Science and Technology
Table 1: Tensile properties of
the investigated alloys
YS [MPa] UTS [MPa] A [%] Z [%]
Temperature 20°C 130°C 20°C 130°C 20°C 130°C 20°C 130°C
2014-A 421 386 454 402 1,8 9,3 3,7 15,8
2014-B 357 341 425 393 7,3 8,4 14,2 16,9
2618 350 345 416 409 6,4 8,6 10,8 12,3
Table 2: Fracture toughness of the materials
investigated (K Q or K IC according to ASTM B645
and K JIC accordiing to ASTM E1820)
Specimen Temperature 2014-A 2014-B 2618
direction [°C] MPa . m Ω MPa . m Ω MPa . m Ω
CL 20 19,3 (K IC ) 23,0 (K Q ) 22,4 (K Q )
CL 20 20,1 (K IC ) 24,9 (K Q ) 23,3 (K Q )
Average CL 20 19,7 (K IC ) 24,0 (K Q ) 22,9 (K Q )
CL 130 18,8 (K IC ) 23,4 (K JIC ) 21,9 (K Q )
CR 20 23,4 (K IC ) 24,8 (K Q ) 26,9 (K Q )
CR 20 22,8 (K IC ) 24,4 (K Q ) 26,7 (K Q )
Average CR 20 23,1 (K IC ) 24,6 (K Q ) 26,8 (K Q )
CR 130 21,4 (K JIC ) 23,4 (K JIC ) 25,6 (K Q )
to the experimental curves
The condition of plain strain crack propagation in CT specimen fracture
can be met for sufficiently large specimens. The minimum specimen thickness
Fig. 2: Load vs. COD plots of the samples tested at 20°C.
A) 2014-A forging, CL specimen; B) 2014-B forging, CR specimen

Table 3: F values for the materials and specimen
orientation investigated
Specimen Temperature 2014-A 2014-B 2618
direction [°C]
CL 20 9.4 4.8 4.9
CL 20 8.7 4.1 4.5
Average CL 20 9.1 4.5 4.7
CL 130 9.1 4.2 4.9
CR 20 6.4 4.1 3.4
CR 20 6.8 4.3 3.4
Average CR 20 6.6 4.2 3.4
CR 130 7.4 3.3 3.6
Table 4: J IC values (N/mm) of the investigated
forgings tested at 130°C
Specimen Temperature 2014-A 2014-B 2618
direction [°C] [N/mm] [N/mm] [N/mm]
CL 130 - 7.23 -
CR 130 5.96 7.00 -
(B) to be adopted can be calculated, at the end of the tests, according to
the two standard above considered:
B min =5·(K Q /YS) 2 according to ASTM B645 (1)
B min =2,5·(K Q /YS) 2 according to ASTM E399 (2)
A) B)
Fig. 2: Load vs. COD plots of the samples tested at 20°C.
A) 2014-A forging, CL specimen; B) 2014-B forging, CR specimen
7 - Metallurgical Science and Technology
Thus, the minimum thickness required by the E399
standard is half of that required by the B645
standard for aluminium alloys. The agreement to
plain strain conditions and the possible deviation
from the minimum value of B can be evaluated
using a parameter F, defined as:
F=B·(YS/K Q ) 2 (3)
Where F greater than 5 means that the
requirements for plain strain condition are met
for both standards and valid K IC can be obtained.
2,5

according to the ASTM E1820 standard. As far as
the AA2618 alloy that resulted to be the toughest
material at room temperature, at 130°C it showed
unstable crack propagation (Figure 3b) that
prevented the evaluation of J IC. In the cases of valid
J IC , K JIC was computed, assuming plain strain
conditions, by using the following correlation
proposed by ASTM E1820 standard:
3.3. FRACTOGRAPHIC
OBSERVATIONS
Fractographic observations were performed at
two main locations on fracture surface of each
specimen in the crack propagation region: the first
next to, the latter at 4 mm from the boundary
with fatigue precrack. Selected images at low
magnification of the first location for the different
alloys and specimen orientation are reported in
figure 4. In the unstable crack propagation region,
the 2014-A forging showed the typical features of
ductile fracture, generated by nucleation of dimples
mainly from the coarser intermetallic particles
fractured in a brittle way (figure 5a). Within the
matrix, regions with dimples of far smaller size
were also visible, laying on well-defined planes
A)
C)
K JIC = [(E/(1-v 2 ))•J IC ] 0.5 (4)
K JIC values are also listed in Table 2 for comparison purpose. Examination of
this table clearly shows that toughness of 2014-B forging is greater than
that of the same alloy in forging 2014-A. Further, it can be stated that in all
the examined samples, crack propagation in CL orientation is favoured
with respect to that in CR orientation.
(figure 5b). These can be correlated to the presence of fine particles that in
some cases were observed inside the dimples in high magnification images
that suggest ductile intergranular fracture.
The fracture surfaces of the 2014-B forging did not show significant
differences from 2014-A material, with relatively extended regions of
microdimples (figure 7a) that could be correlated to the presence of coarse
grain boundary particles (figure 6b). The size and distribution of these
particles corresponded to those observed at grain boundaries (figure 4c).
The 2618 forging, showed a transgranular ductile fracture mode (figures
4d, 7b). Dimples observed on fracture surfaces nucleated from the
homogeneously distributed FeNiAl 9 particles (figures 6c and 6d).
The fracture surfaces of specimens tested at 130°C of forgings made of
AA2014 alloy were similar to those tested at room temperatures (figure
8a). On the contrary, as the temperature increased, small-size microdimples
were observed on the fracture surface of the forging made of AA2618 alloy
(compare figure 7b and 8b).
Fig. 4: Fracture surface of cracks propagated at room temperature in the boundary region between fatigue precracking and unstable crack
propagation region. 2014-A forging, CL (a) and CR (b) directions. 2014-B forging; CR direction (c) and 2618, CR direction (d).
B)
D)
8 - Metallurgical Science and Technology

A) B)
A)
C)
Fig. 5: Fracture surface of cracks propagating at room temperature in 2014-A forging in CL (a) and CR (b) direction.
Fig. 6: Microstructure in the region of unstable crack propagation on longitudinal section of CT specimens. a) 2014-A forging, CL direction;
b) 2014-B forging, CL direction; c) and d) 2618 forging, CL and CR directions, respectively.
B)
D)
A) B)
Fig. 7: Fracture surface sampled in CL direction of 2014-B forging (a) and of 2618 forging (b) tested at room temperature.
9 - Metallurgical Science and Technology

A) B)
Fig. 8: Fracture surface of cracks propagated at 130°C in specimen with CL orientation of 2014-A (a) and 2618 (b) forgings.
4. DISCUSSION
In the examined forgings crack propagated in a
ductile manner, in some cases by an intergranular
ductile mode, linking the early fractured coarse
intermetallic particles, located at inter or
intragranular position in the different materials
investigated.
Following the approach proposed by Hahn and
Rosenfield [2] for aluminium alloys, it can be stated
that crack propagates when the size of extensive
plastic strain formed ahead of the crack tip
corresponds to the average coarse interparticle
distance (δ) Further, according to these authors,
K and d are correlated as follows:
δ =(0.5K 2 )/(E•YS) (5)
In the case of 2618 forging the δ values estimated
using room temperature tensile and toughness
properties were 11 and 15 µm for specimen
orientations CL and CR, respectively, comparable
to the average interparticle size between the
intragranular particles (mainly FeNiAl 9 particles)
in longitudinal and radial direction, respectively. The
preferred particle orientation and their tendency
to align axially in the forging regions where
specimens were sampled corresponds to a greater
CONCLUSIONS
The toughness properties in opening mode I of
forgings made of AA2014 and AA2618 aluminium
alloys with different microstrctural conditions were
presented.
Fracture toughness ranged between 19 and 26
MPam 0.5 , depending on the material and specimen
direction. Fracture toughness was in general lower
interparticle distance in the crack propagation direction of CR specimens.
As previously presented, in the specimens obtained from forgings made of
AA2014 alloy the crack proceeds in a ductile manner but along an
intergranular path, intercepting the oriented coarse aligned intermetallic
particles axially. The values of δ computed from equation 5 for these two
forgings are 6.7 and 9 µm for CL and CR specimen of 2014-A forging,
respectively. In the case of the last forging (2014-B), δ equals 12 µm for
both CL and CR directions. Thus, also in the case of AA2014 alloy, d is
comparable to the interparticle distance along the crack path. It can be
thus stated that fracture toughness is correlated to the mean interparticle
distance along the crack path.
The above observations well agree with the simple correlation proposed
by Hahn and Rosenfield [2] for the case of forged Al-alloy parts (where
fracture is initiated by cracking of large intermetallic particles once reached
critical stress/strain levels in the region of extensive deformation at the
crack tip [3]) and fracture toughness is correlated to the mean interparticle
distance along the crack path.
Thus, forgings of the same heat treatable alloy, even in the same heat
treatment condition and with comparable hardness and tensile properties
can show significantly different fracture behaviour, depending on intermetallic
particle population. Especially, as the mean interparticle distance along the
crack path between the coarser intermetallic particles increases, the
nucleation of voids at these particles is shifted at higher applied loads and
toughness is improved.
The role of coarse intermetallic particles and the need to reduce their
amount, to optimize their size and distribution in forgings (taking into account
the most critical crack propagation directions in the components) in order
to increase material toughness is therefore highlighted.
for specimen sampled in CL direction. In the examined forgings crack
propagated in ductile manner, in some cases by an intergranular ductile
mode, linking in any case the voids formed at the early fractured inter- or
intragranular coarse intermetallic particles. No significant difference in
toughness nor in the fracture mode was observed when tests were
performed at room temperature or at 130°C.
Fracture toughness properties were related to the presence and distribution
of the coarse intermetallic phases within grains or at their boundaries in
the different investigated alloys.
10 - Metallurgical Science and Technology

The simple correlation proposed by Hahn and Rosenfield between fracture
toughness and the mean interparticle size was found to be applicable when
the interparticle distance along the crack path (different for different sampling
direction in forgings where these particle were aligned along flow directions)
is taken into account. An observation and quantitative assessment of the
REFERENCES
[1] J.S. Robinson, R.L. Cudd, J.T. Evans. Creep resistant aluminium alloys
and their applications. Material Science and Technology, 19 (2003) pp.
143-155.
[2] G.T. Hahn, A.R. Rosenfield. Metallurgical Factors Affecting Fracture
Toughness of Aluminum Alloys. Metall. Trans. A, (1975) pp. 653-668.
[3] T. Kobayashi. Strength and fracture of aluminium alloys. Materials Science
Engineering, A286 (2000) pp. 333-341.
[4] A.P. Reynolds, E.P. Crooks. The effect of thermal exposure on the fracture
behaviour of aluminium alloys intended for elevated temperature
service. in “Elevated temperature on fatigue and fracture”, ASTM STP
1297. Williamsbourg, Virginia(USA), (1997) 191-205.
[5] B.Q. Li, A.P. Reynolds. Correlation of grain-boundaries precipitates
parameters with fracture youghness in an Al-Cu-Mg-Ag alloy subjected
to long-term thermal exposure. Journal of Materials Science, 33 (1998)
pp. 5849-5853.
[6] ASTM B 645-02 Practice for plane-strain fracture toughness testing of
aluminium alloys.
[7] ASTM E 399-99 Standard test method for plane-strain fracture
toughness of metallic materials.
[8] ASTM E1820-01. Standard test method for measurement of fracture
toughness.
11 - Metallurgical Science and Technology
distribution of coarse intermetallic phases on
optical micrographs can be a useful tool to check
fracture toughness properties of different forgings
in their peak aged condition when yield strength
is known.

VARIATIONS IN MICROSTRUCTURE AND
MECHANICAL PROPERTIES OF CAST
ALUMINIUM EN AC 43100 ALLOY
Salem Seifeddine, Torsten Sjögren and Ingvar L Svensson
Jönköping University, School of Engineering, Component technology - Sweden
Abstract
The microstructure and mechanical properties of a gravity
die and sand cast Al-10%Si-0.4%Mg alloy, which is one of
the most important and frequently used industrial casting
alloys, were examined. Tensile test samples were prepared
from fan blades and sectioned through three positions
which experienced different cooling rates. Furthermore,
the inherent strength potential of the alloy was revealed
by producing homogeneous and well fed specimens with a
variety of microstructural coarseness, low content of oxide
films and micro-porosity defects, solidified in a laboratory
environment by gradient solidification technology. The
solidification behaviour of the alloy was characterized by
thermal analysis. By means of cooling curves, the
solidification time and evolution of the microstructure was
recorded. The relation between the microstructure and
the mechanical properties was also assessed by using quality
index-strength charts developed for the alloy. This study
shows that the microstructural features, especially the ironrich
needles denoted as β-Al FeSi, and mechanical
5
properties are markedly affected by the different processing
routes. The solidification rate exerts a significant effect on
the coarseness of the microstructure and the intermetallic
compounds that evolve during solidification, and this
directly influences the tensile properties.
KEYWORDS
Aluminium cast alloys, gravity- and sand casting,
gradient solidification technique, thermal analysis,
porosity, quality index.
Riassunto
In questo lavoro sono state esaminate la microstruttura e le proprietà meccaniche di
una delle leghe più comuni nell’industria delle leghe leggere, la lega Al-10%Si-0.4%Mg,
colata per gravità sia in conchiglia che in sabbia. Le provette di trazione sono state
ricavate da pala di ventilatore e sono state sezionate in tre posizioni caratterizzate da
diverse velocità di raffreddamento. Inoltre è stato valutato il valore potenziale di resistenza
a trazione ottenibile della lega, mediante produzione di campioni caratterizzati da diversa
morfologia strutturale, basso contenuto di film ossidi microporosità, solidificati in
ambiente controllato con metodologia di solidificazione a gradiente. Il comportamento
in solidificazione è stato caratterizzato mediante analisi termica. Per mezzo di curve di
raffreddamento, si è determinato il tempo di solidificazione e l’evoluzione della
microstruttura. La relazione tra la microstruttura e le proprietà meccaniche è stata
determinata usando specifiche carte di correlazione indice di qualità – resistenza. Questo
studio dimostra che le caratteristiche microstrutturali, in particolare i grani aciculari ad
alto tenore di ferro indicati come β-Al5FeSi, e le proprietà meccaniche sono
marcatamente condizionate dai differenti percorsi preparativi. La velocità di
raffreddamento esercita una significativa influenza sulla morfologia microstrutturale e
sui componenti intermetallici che si sviluppano durante la solidificazione, influenzando
così direttamente le proprietà di resistenza a trazione.
12 - Metallurgical Science and Technology

INTRODUCTION
The mechanical properties of cast aluminium alloys are very sensitive to
composition, metallurgy and heat treatment, the casting process and the
formation of defects during mould filling and solidification. The coarseness
of the microstructure and the type of phases that evolve during solidification
are fundamental in affecting the mechanical behaviour of the material. The
mechanical properties obtained from gradient solidified samples are assumed
to represent the upper performance limits of a particular alloy, due to the
low level of defects obtained from the favourable melting and solidification
procedure.
Different processing routes offer different qualities and soundness of
commercial components. The true potential of most alloys is seldom
approached, and there exists a lack of data in the literature describing the
variations in microstructure and related mechanical properties of alloy EN
AC 43100. The present investigation was therefore carried out in order to
Fig. 1: Fan blade cast in sand- and gravity moulds.
13 - Metallurgical Science and Technology
elucidate and assess the influence of the casting
process and resultant variations of microstructure
on the mechanical properties, by extracting tensile
specimens from Al-10%Si-0.4%Mg commercial cast
components manufactured by two different casting
methods; sand and gravity die casting. In addition,
the inherent strength potential of this casting alloy
will be investigated by manufacturing tensile test
specimens using gradient solidification technology.
The mechanical properties which have been
measured in this work are: ultimate tensile strength,
yield strength and fracture elongation. The
microstructural features that have been evaluated
are the secondary dendrite arm spacing (SDAS),
pore fraction and precipitated phases such as the
length of the iron-rich β-phase, Al 5 FeSi.
MATERIALS AND
EXPERIMENTS
Cast component
The components that have been studied are fan
blades, as shown in figure 1, which are produced
commercially in a great variety of shapes and sizes,
and by different casting methods. This particular
component is considered to be of interest to study
and analyse due to the widespread use of different
casting methods to manufacture fan blades, which
in this case were manufactured by sand mouldand
gravity die casting.
The locations of the specimens extracted for study
are shown in figure 1; from the top (T), middle
(M) and bottom (B) parts of the fan blade. These
locations were chosen due to their different
solidification times (i.e. thicknesses), which results
in different coarseness of the microstructure,
intermetallic phases and defects.
Cast alloy
Table 1: Chemical composition of the casting alloy
Composition Elements (wt. %)
The commercial blades and specimens produced
by gradient solidification experiments were cast
using alloy EN AC 43100. The standard
composition of the alloy and the actual
composition are given in table 1. Strontium was
Alloy Standard Si Fe Cu Mn Mg Zn Ti Cr Ni Al
EN AC-43100 SS-EN 1706 9.0-11.0 0.55 0.10 0.45 0.20-0.45 0.10 0.15 - 0.05 Bal.
EN AC-43100 Actual sample 9.96 0.47 0.11 0.25 0.35 0.097 0.018 0.01 0.01 Bal.

added in the form of Al-10%Sr master alloy at a
level of 150- 200 ppm.
Sand and gravity die casting
The sand moulds used were made by hand. As a
pattern, a fan blade cast in a gravity mould was
used, which is as already mentioned, one of the
methods normally used to manufacture this size
and type of fan blade. The sand mould consisted
of chemical bonded sand.
Gradient solidification
The same alloy as that used for the sand and gravity
die cast components, shown in table 1, was cast
into rods for further solidification studies
in the gradient solidification equipment. The
gradient solidification technique allows the
production of a high quality material with a low
content of oxide films, porosity and shrinkage
related defects, and produces a homogenous and
A)
well-fed microstructure throughout the entire specimen. In this work a
resistance-heated furnace with an electrically driven elevator was used, see
figure 2a. Three different growth velocities, v, were used, 0.03 mm/s, 0.3
mm/s and 3 mm/s, which correspond respectively to an SDAS of about 50,
20 and 7 µm. At each velocity, three specimens were made. The specimens
cast in the gradient solidification equipment were protected by argon in
order to avoid oxidation.
Thermal analysis
In order to study the thermal history of the solidified components,
thermocouples were inserted at three different positions in the sand cast
fan blade, as indicated in the picture in figure 1. To record the cooling curves
a data acquisition device from National Instruments was used. In each of
the sand moulds three thermocouples were placed. The thermocouples
were of S-type, which means that the different materials in the thermocouple
wires are platinum and platinum-rhodium (90Pt-10Rh). When manufacturing
the thermocouples, the wires are inserted in a two-hole Al 2 O 3 -tube to
separate the wires, which are connected to a plug. The thermocouples
were protected from the molten metal by a quartz glass tube. The maximum
operating temperature for the S-type thermocouple is over 1500°C [1].
The recording rig is illustrated in figure 2b.
B)
Fig. 2: . Illustration of the instruments that have been used during the investigation. While a) demonstrates the gradient solidification equipment,
b) presents the data acquisition rig that is connected to the sand moulds.
14 - Metallurgical Science and Technology

RESULTS OF THE
MICROSTRUCTURE AND
MECHANICAL PROPERTIES
INVESTIGATION
Since specimens were extracted from fan blades
with different thickness and at different sections
it was necessary to machine them in order to
obtain specimens suitable for tensile testing. The
gradient solidified specimens have also been
prepared in the same manner, and the dimensions
of the samples used for tensile testing are shown
in figure 3. In order to analyse the microstructure
and measure the different microstructural features,
a light microscope, image processing software and
a scanning electron microscope, SEM, were used.
Microstructure development
The microstructure coarseness is defined as the secondary dendrite arm
spacing, SDAS, which is also a function of the local solidification time, and
quantifies the scale of the microstructure and its constituents. Examining
the microstructure of specimens from both sand and gravity die casting, it
has been observed that there is a wide scattering of SDAS and consequently
variations of sizes and shapes of the intermetallic compounds within the
same component. The needle-shaped β-phase, Al 5 FeSi also seems to coexist
with the α-phase, Al 15 (Fe,Mn) 3 Si 2 , which has the morphology of Chinese
script. As depicted in figure 4, the dendrites are equiaxed and the eutectic
is finely modified. Comparing the microstructure formations, it can be noticed
that in the case of gravity die casting the microstructure is finer, due to the
higher cooling rate. Furthermore, the different coarsenesses of
microstructure, the phases and their morphologies are comparable to the
sand cast microstructure.
It is relevant to note that even if Mn has been added to the alloy to act as
A) B)
Fig. 3: A schematic illustration of the tensile specimen (dimensions in mm).
Fig. 4: . a) Illustration of the microstructure of a gravity die cast specimen while b) shows a sand cast microstructure. The black spot is a pore.
15 - Metallurgical Science and Technology
an Fe-corrector and to promote the formation of
Chinese script, see table 1, iron-rich Al 5 FeSi-needles
are still formed. These needles act as local stress
raisers in the matrix and are considered to be
deleterious for the mechanical performance, which
is why their formation and growth should be taken
into consideration and as much as possible
suppressed.
In order to understand the local variations in
microstructure, and to understand the formation
sequences and development of constituent phases
in the sand cast component, thermal analysis was
implemented. When manufacturing the sand cast
fan blades at the foundry, data for the cooling
curves was recorded. In each of the sand moulds,
three thermocouples were used, as shown in figure 1.

With this composition, see table 1, and casting
technique, three different cooling curves have been
obtained and are presented in figure 5a. The figure
also shows that there exists clear variations in local
solidification times which are directly related to
the variations in component thickness, see also
table 2. These variations in cooling time lead to
major local differences in microstructures which
affect the tensile performance of the cast
components. The SDAS for the sand cast and
gravity die cast component varied between 47-61
µm and 20-45 µm respectively.
Figure 5b shows that the freezing range is
approximately between 595°C and 545°C. From
the different cooling curves it is possible to see at
which temperature and time the dendritic network
starts to grow, at which temperature and time the
eutectic starts to grow and also the total
solidification time, from pouring the melt until the
fraction solid has reached 100%.
When the value of the derivative exceeds zero,
figure 5b, the temperature is actually increasing
due to the release of latent heat due to phase
transformations. This is most evident in the case
where a positive derivative is observed at about
A) B)
Mechanical properties
The results of the tensile tests from the
commercial fan blades and the laboratory
prepared samples are shown in the diagrams in
590°C when the dendritic network starts to grow. The next positive value
of the derivative appears when the eutectic starts to grow, at about 570°C.
When the temperature is around 550°C precipitations of a complex eutectic
with intermetallics and a hardening phase such as Mg 2 Si starts to precipitate
from the remaining melt. From the different thermocouples the following
solidification times have been determined:
Table 2. Solidification times for the different parts
of the sand cast fan blade.
Thermocouple Thinnest part Mid part Bottom
(T) (M) part (B)
Solidification time (s) 260 470 610
Material thickness (mm) 12 20 28
Examining the gradient solidified specimens, a significant discrepancy between
the different solidification rates is observed. The higher the cooling rate
the finer the microstructure, as illustrated in figure 6. The iron-rich
β-phase Al 5 FeSi is found to be very short and very well dispersed for the
two highest solidification velocities; 3 and 0.3 mm/s, with a corresponding
SDAS ≈ 7 and 23 µm respectively, in comparison to the lowest velocity
0.03 mm/s, with SDAS ≈ 47 µm, where long iron-bearing phases were found
randomly distributed in the eutectic regions.
Fig. 5: The thermal history of the sand cast component is registered in figure a) while in b) an analysis of the cooling curve
obtained from the bottom section is performed.
figure 7. The two different casting methods are compared in figure 7a, where
each point represents a different position on the fan blade. A wide scatter
in data is clearly observed but it can be noticed that the properties of the
thinner parts seem to approach higher combinations of stress and strain
values.
16 - Metallurgical Science and Technology

In order to simulate the processes of the cast components, tensile test
samples of the same alloy and corresponding microstructural coarseness
have been produced by using gradient solidification. By producing a more
homogenous microstructure with a lower level of defects, the material’s
inherent potential was revealed, see figure 7b. The scatter in data in figure
7b is ascribed to different levels of defects, and the higher stress-strain
A) B)
Fig. 6: Phases and microstructures for a) SDAS = 47 µm obtained at 0.03 mm/s and b) SDAS = 7 µm obtained at 3 mm/s.
A) B)
µ
µ
Fig. 7: The graph in a) illustrates the ultimate tensile strength and elongation to fracture for the gravity die and sand casting specimens where B is
the bottom (thickest) section, M the middle and T the top (thinnest) part of the blade. The graph in b) presents the tensile properties for the
gradient solidified specimens with a range of cooling rates.
17 - Metallurgical Science and Technology
combinations are the limits for what this alloy
might perform under ideal conditions.
When quantifying the SDAS, five measurements
were taken at three randomly chosen positions of
the specimen. From these data a mean value was
calculated. At high solidification rates the SDAS
µ
µ
µ

appears to be very small and the microstructure
has been found to be more uniform. As noticed,
the overall effect of obtaining a small SDAS is an
improvement in both ductility and tensile strength,
see figure 8.
When measuring the length of the iron-rich Al 5 FeSi
needles (β-phase), the 15 longest needles observed
in the microstructure were measured, considering
the entire cross-section area of the sample, and a
mean value was assessed. It should be mentioned
that the needles were not chemically analyzed, but
A) B)
µ
µ
µ
µ
µ
µ
were identified according to their morphologies and colours. A correlation
was found between the length of the β-phase and the mechanical properties.
An increase in needle length results in a reduction of tensile strength and
elongation as can be seen in figure 9. Note that the iron-bearing β-phase
was too short to be measured or not present at all in the specimen solidified
at a rate of 3 mm/s, SDAS = 7 µm. It is also seen that even if these needles
are relatively shorter in some cases than in corresponding samples, they
exhibited lower properties. The premature fractures must therefore be
due to other defects such as oxide film inclusions or other brittle and
undesirable phases such as Al 8 Mg 3 FeSi 6 . These intermetallics are harmful to
the mechanical properties since cracking of these phases reduces the ductility,
Fig. 8: The graphs a) and b) illustrate the influence of SDAS on the tensile properties.
A) B)
µ
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µ
µ
µ
18 - Metallurgical Science and Technology
µ
µ
µ
µ
µ
µ
Fig. 9: The influence of the β-phase needle length on a) ultimate tensile strength and b) the elongation to fracture.
µ
µ
µ
µ
µ

µ
µ
µ
µ
µ
µ
Fig. 10: This graph shows the relation between the β-phase length and the SDAS.
and their formation reduces the Mg available in the melt to form Mg 2 Si
which has a hardening effect and improves both the yield and ultimate
tensile strengths.
Generally, larger SDAS is associated with lower cooling rates and also with
longer β-phase needles. The local solidification time influences the growth
of the intermetallics that develop and grow during the solidification process,
and hence a relation between SDAS and the β-phase length is observed as
shown in figure 10.
When a tensile sample contains pores it is reasonable to assume that the
region of porosity will yield first due to the reduced load bearing area
concentrating the stress near the voids. The area fraction of porosity was
determined by measuring the area just below and at the fracture surface.
A) B)
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Fig. 11: The graphs a) and b) illustrate the influence of porosity on the tensile properties.
19 - Metallurgical Science and Technology
The first mentioned was measured by an image
analyzer and the latter was photographed and
observed by scanning electron microscopy, SEM.
This method enables an examination of the
distribution and size of the pores.
In the present study, it has been observed that the
cooling rate, SDAS in this case, seem to govern
the length as well as the percentage area fraction
of porosity to some extent, see figure 11. It appears
that a reduction in SDAS results in smaller average
pore size and a reduced area fraction of porosity.
It is obvious that as the solidification time is low,
less time will therefore be available for the diffusion
of the hydrogen into the interdendritic regions
which results in small sizes and fraction of pores.
But comparing the porosity formation, the gradient
solidified samples with SDAS ~ 23 µm are
associated with largest percentage area fraction
of porosity which might be due to the arrestment
of premature pores as they become entrapped by
the advancing solidification front.
The discrepancy could be derived from the theory
that involves the harmfulness of the oxide films,
which suggests that when an oxide particle is
approached by the solidification front, the particle
experiences the hydrogen-rich environment
produced by the rejection of gas from the
advancing solid. Furthermore, the access of gas by
diffusion into the air pocket in the gap of the oxide
particles, the pore will start to form and grow. As
the freezing rate is slow and the particle is poorly
wetted by the melt, time will therefore be available
for more hydrogen to diffuse resulting in pore
expansion and growth.
µ
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In the case of the larger SDAS ~ 47 µm, it is
therefore reasonable to assume that due the slow
cooling conditions many air pockets formed within
the liquid or due to interaction with oxide particles
have been driven in front of the solidified front
without being engulfed and in that case not been
detected in the gradient solidified specimens due
to the solidification mode. Another reason might
be the longer time available for the hydrogen to
diffuse and move in front of solid-liquid interface
out of the sample into the surrounding
environment.
The scatter in the data is too great to derive any
clear correlation between porosity level and the
tensile strength and ultimate elongation, see figure
11 a and b. According to literature, it is stated [2]
that the static tensile properties such as ultimate
tensile strength, yield strength and elongation to
A) B)
DISCUSSION
Generally, the mechanical properties of metals are
widely influenced by their microstructures. When
defining the microstructure many parameters have
to be taken into account, and these include among
others the secondary dendrite arm spacing, the
size and shape of precipitated phases such as silicon
and iron-bearing phases, grain size, porosity, etc.
In this investigation, many of the parameters
mentioned above have been measured and
correlated to the mechanical properties.
Constituents that may have deleterious effect on
the mechanical properties are defects such as
oxide films or undesired phases and inclusions.
A comparison of the results of the tensile testing
with the casting method and the corresponding
fracture, were all decreased with an increased degree of porosity. On the
contrary, according to [3-7], the reduction in tensile properties has almost
no correlation with the average volume fraction of porosity. In fact, the
decrease in tensile properties was attributed to the length and/or the area
fraction of defects in the fracture surface. Development of a high fraction
of porosity may also be due to Fe in the melt. The long, needle-shaped iron
intermetallic phase formed is expected to cause severe feeding difficulties
during solidification. The morphology of the β-phase blocks the interdendritic
flow channels, which is why it is proposed that higher iron contents in the
alloy are associated with higher levels of porosity.
SEM examination was also performed in order to study what kinds of pores
are frequently found on the fracture surfaces in this work. Figure 12a
illustrates a pore seen in a sand cast specimen and 12b is seen in a gradient
solidified specimen at a solidification rate of 3 mm/s.
Worth to indicate is that the melt hydrogen content has not been measured
and assumed in this case to be constant for all the alloys since they have
been produced under similar conditions.
Fig. 12: Illustration of pores. a) Showing sphericity of a possible gas pore b) The cavity is observed at the edge of the
fractured surface shows a possible gas-shrinkage pore.
specimens with three different solidification rates obtained by gradient
solidification shows clearly that the tensile behaviour and production method
is related by three different stress-strain curves. The material exhibiting
SDAS ≈ 7 µm results in the highest stress-strain combinations as shown in
the diagram in figure 13. The main reason for this behaviour is the finer
microstructure, a more homogeneous distribution of the intermetallic phases
and lower levels of porosity. Even though there are no specimens with
corresponding SDAS extracted from commercial components available in
this study, this data clearly shows the inherent potential strength of the
alloy.
The stress-strain values for the sand cast specimens follow the curve for
SDAS = 47 µm (obtained by gradient solidification), but fracture occurs at
lower stress and strain levels, see figure 13. The stress-strain curve for
SDAS = 23 µm (by gradient solidification) corresponds well to the stressstrain
values of the gravity die cast specimens.
The reason why the sand and gravity die cast specimens fractured much
earlier than the gradient solidified specimens, and exhibited different
20 - Metallurgical Science and Technology

µ
µ
µ
µ
µ
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µ
Fig. 13: This graph shows tensile test curves for the gradient solidified specimens and
includes mechanical properties data for the gravity die and sand cast specimens.
combinations of strength and fracture elongation values may be due to the
formation of brittle and coarse intermetallics such as the iron-bearing
β-phase and an inhomogeneous microstructure. Even if similar SDAS has
been obtained from the gradient solidification equipment, the variations in
the β-phase needles, local coarsenesses of microstructures and sizes of
intermetallics are likely to be different due to microsegregation in the
castings. Segregation of elements is likely to occur in castings and this will
affect the development and solidification sequences of the microstructure
and lead to local variations in mechanical properties. At the thinner part of
the fan blades the β-phase needles are likely to be short and have no time
to grow and coarsen as is the case in thicker parts where the local iron
A) B)
µ
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21 - Metallurgical Science and Technology
level might also be different. It has also been shown
that a finer microstructure with smaller SDAS is
beneficial for the mechanical properties. A finer
microstructure means larger areas of grain
boundary leading to more difficulties for
continuation of slip and dislocation movement
during deformation, hence yielding stronger alloys.
A simple tool that might be used to predict the
impact of changes in the chemical composition,
solidification rates and routes, microstructure and
heat treatment on the tensile performance is by
implementing a quality index, Q. Its practical value
originates from the fact that the mechanical
“quality” of an alloy is expressed by a single
numerical parameter, whose value can be read off
a quality index chart or calculated through a simple
equation relating the elongation to fracture, s , the f
ultimate tensile strength, UTS, and d, which is an
empirical parameter chosen to be around 150 MPa
for EN AC 43100 (Al-7%Si-0.4%Mg) to make Q
more or less independent of the yield strength of
the alloy, as seen in equations 1-2, [8-10]:
Q = UTS + d * log (100*s ) (1)
f
n
σ = Kε
(2)
where σ is the true flow stress, K is the alloy’s
strength coefficient, ε is the true plastic strain and
n the strain hardening exponent.
In order to produce the charts in figure 14, a mean
value for the strength coefficient, K, has been
calculated. For the commercial cast components,
figure 14a, K is approximately 410 MPa and for
the gradient solidified materials K is around 430
MPa, figure 14b. The strain hardening exponent, n,
is numerically equal to the necking onset strain
Fig. 14: A quality index chart for the alloy studied obtained for gravity- and sand cast samples is presented in a) while b) represents the gradient
solidified specimens. The mean value of the strength coefficient is a little higher for the gradient solidified specimens.
µ
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and can be used to define the relative ductility
parameter, q, see equation 3.
q = s f /n (3)
Defining q = 1 means the onset of necking
represents samples that fail either at the onset of
necking or beyond, while q < 1 identifies lower
quality or less ductile samples. For instance q =
0.5, 0.2, 0.1 identifies samples that fail at 50%, 20%
and 10% of their maximum uniform strain,
respectively.
As observed, the points corresponding to
SDAS = 52 µm are located near the bottom of
the chart, which may be due to their coarser
CONCLUSIONS
As an outcome from the current investigation, the
dendritic cell size, SDAS, and the length of ironbearing
phase whether as needles or Chinese
script, governs the tensile strength and ductility of
Al-Si-Mg alloys. The early fracture of sand and
gravity die cast specimen can be related to the
length of the iron-bearing phase, oxide films
entrapment, and inhomogeneous microstructure
and to the Si particle size and morphology, which
determine the rate of particle cracking with the
applied strain.
• A finer microstructure, small and well
distributed iron-rich compounds and low levels
of porosity exert a beneficial impact on the
resulting mechanical properties.
• The locally increased levels of iron due to
ACKNOWLEDGEMENTS
REFERENCES
1. Metals Handbook; ASM International, 1998, pp.
663-665.
2. M. Avalle, G. Belingardi, M.P. Cavatorta and R.
Doglione, International Journal of Fatigue Vol. 24
, (2002), pp. 1-9.
3. C.H. Cáceres and B.I. Selling , Materials Science
and Engineering A, Vol. 220, Issues 1-2, (1996),
pp. 109-116.
4. M. Harada, T. Suzuki and I. Fukui , Journal of the
Japan foundrymen´s society Vol. 55, No. 12,
(1983), pp. 47-50.
microstructure and the length of the β-phase. As the strength of the gravity
die cast specimens increases, the data points shift toward the upper right
corner.
Regarding the gradient solidified samples, the data points for the
SDAS = 7 µm specimens, with the finer microstructure, exhibit the highest
values of Q and q. These values seem to be the only ones that reach the
limit imposed by the necking line q = 1.
As the SDAS and the level of deleterious microstructural compounds are
decreased, the strength and ductility of the alloy is increased and so does
the Q, allowing for a possible comparison between different alloys and
processing routes. Furthermore, the Si particles, playing a role as reinforcing
phases, decrease in size and the shape becomes more fibrous, which has an
overall beneficial effect on the mechanical properties.
possible segregations might have resulted in longer and larger fractions
of β-phase needles. The main reason for the remarkable reduction of
the mechanical performance of commercial components is proposed
to be due to the presence of the locally induced stresses by the needles
which break, link and propagate the cracks in an unstable manner. The
higher the solidification rate the shorter the length of the needles and
the sounder the material.
• The gradient solidification method shows that the EN AC 43100 alloy
used for commercial cast components processed by sand and gravity
die casting, has an improvement potential to provide sounder castings
with substantial ductility.
• The addition of Mn did not alter the morphology of iron-bearing βphase
into a more compact Chinese script. Instead, these iron
intermetallics seem to be coexisting in the microstructure.
• Porosity is not found to be the primary parameter controlling the tensile
strength and ductility, but it is worth bearing in mind that the reduction
of the load bearing area generally has a harmful influence on the overall
properties.
The authors acknowledge Johan and Jan Tegnemo at Mönsterås Metall AB for their invaluable help and support for making the
experiments in the foundry and Nordic Industrial Fund for the financial support within the Nordic Project STALPK.
5. M.K. Surappa, E. Blank and J.C. Jaquet, Scripta Metallurgica, Vol. 20, No. 9,
(1986), pp. 1281-1286.
6. J.G. Conley, J. Huang, J. Asada and K. Akiba, One Hundred Third Annual
Meeting of the American Foundrymen’s Society; Rosemont, IL; USA, (1998).
pp. 737-742.
7. C.H . Cáceres, Scripta Materialia Vol. 32 No.11, (1995), pp. 1851-1856.
8. C.H. Cáceres, I.L. Svensson, J.A. Taylor, International Journal of Cast Metals
Research, Vol. 15, No. 5, (2003), pp. 531-543.
9. C.H. Cáceres, International Journal of Cast Metals Research, Vol.12, (2000),
pp. 367-375.
10. M. Drouzy, S. Jacob and M.Richard, AFS International Cast Metals Journal,
Vol. 5, No. 2, (1980), pp. 43-50.
22 - Metallurgical Science and Technology

COMPARATIVE STUDY OF HIGH
TEMPERATURE WORKABILITY OF ZM21
AND AZ31 MAGNESIUM ALLOYS
1 M. El Mehtedi, 1 L. Balloni, 1 S. Spigarelli, 1 E. Evangelista, 2 G. Rosen, 3 B.H.Lee, C.S.Lee
1 Dipartimento di Meccanica, Università Politecnica delle Marche, Ancona, Italy
2 ALUBIN, Kiryat Bialik, Israel
3 Department of Materials Science and Engineering, Pohang University, Pohang, Korea
Abstract
High temperature regime, 300-450°C for Mg-Al-Zn alloys, is currently used in primary
processing, such as rolling and extrusion, as well as for secondary operation like forging.
The knowledge of temperature and strain rate proper combination (processing window)
as well as the microstructure evolution occurring during hot deformation clarifies the
relationships between forming variables and final properties of components. Numerous
data on AZ31 and few other Mg-Al alloys, produced by laboratory testing, are available
in the scientific and technical literature. The ZM21, Mg-2Zn-1Mn, by contrast, is
characterized by absolute lack of scientific data. In the alloy the addition of manganese,
by suppressing the formation of beta phase, increases the solidus temperature that
results in the larger processing window than in AZ31. The benefit requires extensive
analysis aimed at optimizing the deformation variables that affect the microstructure
refinement under dynamic and static recrystallization. The high-temperature plastic
deformation and the microstructure evolution of the ZM21 were thus investigated in
the temperature range between 200 and 500°C and results were analysed and compared
with those of a conventional heat-treated AZ31.
23 - Metallurgical Science and Technology
Riassunto
Il regime delle alte temperature 300-450°C, è comunemente
usato per le operazioni primarie di deformazione plastica,
come la laminazione e l’estrusione, come pure per quelle
secondarie come la forgiatura o lo stampaggio delle leghe
Mg-Al-Zn. La conoscenza della combinazione adatta
temperatura-velocità di deformazione (finestra di processo)
e della evoluzione della microstruttura fa luce sulle relazioni
tra variabili di deformazione e proprietà dei componenti.
Numerosi dati, ottenuti con prove di compressione o
torsione, sono disponibili sulla lega AZ31 e su pochi altre
leghe della classe Mg-Al. Sulla lega ZM21 (Mg-2Zn-1Mn)
non sono invece disponibili dati scientifici. Nella lega
l’addizione del Mn sopprime la formazione della fase beta
ed aumenta la temperatura del solidus, producendo una
finestra di processo più ampia di quella dell’AZ31. Questi
aspetti richiedono analisi approfondite sono rivolte
all’ottimizzazione delle variabili di deformazione che poi
governano l’affinamento della microstruttura per effetto
della ricristallizzazione sia dinamica che statica. La
deformazione plastica ad alta temperatura e l’evoluzione
della microstruttura della lega ZM21 sono state analizzate
nell’intervallo di temperatura 200-500°C e i risultati sono
confrontati con quelli di una lega convenzionale trattata
termicamente.

1. INTRODUCTION
Lightweight constructions are becoming more and
more important for the automotive industries due
to increase of fuel prices and legislative
requirements like CAFÉ in the US, and the directive
or the control of CO 2 emissions in the EU [1,2].
Magnesium alloys for their unique combination of
light weight, high specific strength and stiffness and
high recycling capability are candidate to replace
steel, iron and, in some cases, also aluminum parts.
The growing interest in magnesium alloys is
expected to increase in the next years [3]. The
use of components manufactured via die-, sand-,
mould- and rheo-casting is fairly well introduced
in a wide number of applications which are ranging
from steering wheels, instrumental panels, gear box
housing and even to hybrid engine blocks. The
scenario for the automotive industry is that the
use of magnesium alloys for many components,
mainly as castings, is further expanding in the short
term, whereas the application in body and chassis
components, including sheet and extrusions, is
anticipated in the medium term [4].
A driving force for the introduction of wrought
alloys is their property enhancement, since casting
alloys are affected by porosity, which would be of
particular importance for structural and crashrelevant
parts. Initial applications for this category
are expected for extrusions and forgings, to be
followed by sheets. These requirements may
address researches towards the modification of
known alloys as well as the development of new
ones aiming at improving both mechanical
performances, also by ageing (T6), and high
temperature stability. The growing interest in use
of magnesium alloys is expected to increase in the
next years, and to be sustained by research [5].
From recent magnesium alloys conferences [6], the
following items stand out as limiting or retarding
factors for the use of wrought alloys: 1) poor cold
forming properties; 2) limited number of wrought
alloys and lack of processing data. As far as the
item 1) is concerned, it is important to note that
magnesium alloys are difficult to work at low
temperature (

EXPERIMENTAL
The AZ31 experimental result data belong to the authors’ wide database
obtained in the last 3 years by testing a batch of AZ31 with different initial
microstructure: ingot, extruded, heat treated, overaged. These investigations
aimed at clarifying the microstructure features that significantly affect the
high temperature deformation of the alloy; the data were partly published
[11,12] and used to identify the processing window (strain rate and
temperature conditions) that is successfully used for industrial forming
operations. Warm forming experiments were also carried out to assess the
AZ31 sheet formability [13]. The present data were obtained by testing
the specimens machined from rolled AZ31 heat treated at 385°C for 14 h,
at the Pohang University.
The ZM21 alloy was produced and extruded by Alubin, Israel. The specimens
for torsion testing, with gauge radius R, and length L, 5 and 10 mm respectively,
were machined from extruded rods. Torsion tests were carried out in air
on a computer-controlled torsion machine, under equivalent strain rates
ranging from 10 -3 to 5 s -1 and temperatures from 200°C to 400°C. Specimens
were heated 1°C/s by induction coils and maintained 5 minutes before
torsion to stabilize the testing temperature that was measured by
thermocouple in contact with the gauge section. Specimens just after the
fracture were rapidly quenched with water jets to avoid microstructure
RESULTS AND DISCUSSION
The initial microstructure of the alloys before testing is displayed in Figure
1. The heat-treated AZ31 exhibited an equiaxed and fine microstructure.
The microstructure of the extruded ZM21 is composed by duplex grain
size, the small and the large grains were 20 and 60 µm respectively.
Figures 2 and 3 show typical equivalent stress vs. equivalent strain curves
obtained by testing at 200 and 400°C the rolled and heat treated AZ31 and
the extruded ZM21, respectively. The flow curves shape of the rolled AZ31
does not differ appreciably from that observed in the cast [11,12]. At
A) B)
25 - Metallurgical Science and Technology
modifications during slow cooling.
The von Mises equivalent stress, σ, and equivalent
strain, ε, were calculated using the relationships:
σ =
3 M
( 3 + m'+
n'
)
2 π R
3
2 π N R
ε =
2)
3 L
where N is the number of revolutions, M is the
torque, m’ (strain rate sensitivity coefficient) at
constant strain is ∂ log M / ∂ log N,
•
and n’(strain
hardening coefficient) at constant strain rate is
∂ log M / ∂ log N; at the peak stress, clearly n’=0.
To reveal the microstructure, longitudinal sections
at the periphery of the sample gauge, i.e. in the
point where the strain and strain rates assumed
the values calculated by equations 1) and 2), were
investigated by means of optical microscopy.
200°C, under all strain rates, the flow curves of
both alloys exhibit well defined σ y , ~60 MPa for
AZ31 and ~20 MPa for ZM21, a continuous
increase of work hardening up to σ max , which is
larger in AZ31, followed by fracture that in ZM21
occurs at larger strain. At 400°C, AZ31 the flow
curve, after yielding and very limited workhardening,
gets a steady state region, typical
behaviour of dynamic recovery (DRV), up to
fracture. The ZM21 flow curves display yielding
Fig. 1: Microstructure of the rolled and heat-treated AZ31 (a) and of the extruded ZM21 (b).
1)

and rapid hardening up to σ peaks located at ε~0.7,
irrespective of the different strain rates, followed
by a continuous softening towards the fracture
that occurs at e larger than in AZ31; at 400°C and
strain rate 0.05 s -1 the maximum strain value is
2.5. The flow curve exhibiting a peak followed by
softening in stationary state is indicative of dynamic
recrystallization (DRX) that nucleates in proximity
of the peak producing new stable grains whose
size controls the softening and the σ of stationary
state.
In the present instance, the absence of stationary
state and the decrease of σ up to fracture may be
indicative of three processes promoting the
continuous softening: (i) the onset of DRX around
the peak, producing new small grains, (ii) the fast
new-grain growth at the critical size for (iii) crack
nucleation and growth.
A significant decrease in the strain to failure in
torsion can be observed (Figure 4), in particular
in the high-temperature region; an additional
difference is the observation of a well defined
yielding in ZM21, (Figure 3) i.e. of the presence of
a short strain range without load increase, an effect
A)
B)
A)
B)
Fig. 3: Equivalent stress vs. equivalent strain curves obtained in torsion for
ZM21 at 200 (a) and 400°C (b).
26 - Metallurgical Science and Technology
Fig. 2: Equivalent stress vs. equivalent strain curves obtained in
torsion for AZ31 at 200 (a) and 400°C (b).
probably due to a sudden increase in the number of
dislocations unpinned from their solute atmosphere
(either Zn or Mn).
The temperature and strain rate dependence of the
peak stress are shown in Figure 5. The experimental
data are well described by the equation:
ε& = A[ sinh(
ασ)
] n
exp( −Q
/ RT)
3)
where A and α are material parameters, n is a
constant, R is the gas constant, T is the absolute
temperature, and Q is the apparent activation energy
for high-temperature deformation process. In the
AZ31 a best-fitting procedure gave α=0.017 MPa-1 and n= 4.4 for temperatures between 200°C and
400°C. The activation energy, calculated by plotting
sinh(as) as a function of 1/T, was Q=140 kJ/mol. The
value is relatively close to the one of self-diffusion of
Mg (135 kJ/mol) or for diffusion of Al in Mg (143 kJ/
mol). In the case of the ZM21, α was close to 0.049
MPa-1 , and n ranged between 2.6 and 4.4 (average
n=3.3). The activation energy for high temperature
deformation was close to 200 kJ/mol, higher than in
the case of AZ31.

A) B)
Fig. 4: Equivalent strain to failure for AZ31 (a) and ZM21 (b).
A) B)
ασ
Fig. 6: Zener-Hollomon parameter as a function of peak-flow stress for AZ31 and
ZM21.
27 - Metallurgical Science and Technology
Fig. 5: Peak flow stress dependence on applied stress as a function of temperature for AZ31 (a) and ZM21 (b).
ασ
α
ασ
Figure 6 plots the Zener-Hollomon parameter, Z,
defined as
[ ( ) ] n
sinh
Z = ε&
exp( Q / RT)
= ασ
4)
as a function of peak stress, with Q= 170 kJ/mol
and α=0.034 MPa-1 (mean values of those obtained
experimentally for AZ31 and ZM21). Analysis of
Figure 6 clearly demonstrates that the peak flow
stresses are higher in the AZ31 than in the ZM21.
Figure 7 shows typical microstructure of the AZ31
quenched after torsion testing. At 200°C, and a
strain rate of 10-2 s-1 , the structure is mostly
composed by slightly elongated grains (Figure 7a),
with very fine recrystallized grains on several grain
boundaries. At 300°C, under the same strain rate,
the recrystallized fraction is substantially higher
than at the lower temperature, and the average
size of the recrystallized grains is larger (Figure

A) B)
C) D)
Fig. 7: Microstructure of the AZ31 tested at: 200°C-10-2 s-1 (a) 300°C-10-2 s-1 (b) and 400°C- 1 s-1 (c); (d) 400°C-10-3 s-1 .
7b). Fully recrystallization and localised grain
growth at 400°C, in the high strain rate range,
resulted in the presence of some large grains
heavily deformed by twinning, albeit a relatively
large fraction of the structure is composed by fine
equiaxed grains. Under lower strain rate, the
structure is composed by a homogeneous
distribution of relatively coarse equiaxed grains
(Figure 7d).
The mechanism of dynamic recrystallization (DRX)
has been analysed by McQueen and Konopleva
[15]; twinning in fact takes place before the other
mechanisms, except basal glide that occurs in
favourably oriented grains. At low strains, twinning
favourably reorients grains for slip. As deformation
reaches a sufficiently high value, DRX starts where
high misorientations have been created by
accumulation of dislocations, i.e. where slip has
occurred on several slip system (near grain
boundaries and twins). The new fine grains form a
mantle (or a “necklace”) along grain boundaries
and deform more easily than the grain core, thus repeatedly undergoing
recrystallization [12].
Figure 8 shows typical micrographs of the microstructure of ZM21 tested
in torsion. The structure, even at 300°C, is largely unrecrystallized, with
elongated grains; at 400°C, the grains are mostly equiaxed, but elongated
structures still persist. The strong decrease in flow stress after the peak
should be attributed to the late onset (at strain close to ε=0.7-0.8) of
recrystallization and coarsening phenomena, that in this alloy exhibit quite
different kinetics than in the case of the AZ31. Only at the highest
temperature (500°C), the microstructure underwent complete
recrystallization and grain growth, a process that resulted in the presence
of coarse (50 µm) grains.
The above analysis gives some interesting preliminary indications on the
high-temperature formability of the extruded ZM21 alloy. In particular:
1. the intrinsic ductility, i.e. the equivalent failure strain in torsion, is lower
in ZM21 than in AZ31;
2. the fraction of recrystallized structure, at all the investigated
temperatures, is lower in ZM21 than in AZ31;
3. the maximum equivalent stress, i.e. the working load, is higher in AZ31
than in ZM21; yielding has been observed in ZM21, thus suggesting a
solute strengthening role played by either Zn or Mn;
28 - Metallurgical Science and Technology

A)
C)
Fig. 8: Microstructure of the ZM21 tested at: 300°C-5 s-1 (a) 400°C- 5x10-2 s-1 (b) and 500°C- 5 s-1 (c).
4. the different composition of the ZM21 enables the use of higher
working temperature, up to 500°C, i.e. well above the upper limit
allowed for AZ31. On the other hand, even though strain to failure is
relatively high at 500°C, the microstructure undergoes extensive grain
CONCLUSIONS
The high temperature workability of the extruded ZM21 alloy has been
investigated by torsion tests between 200 and 500°C. Both mechanical and
microstructural results were compared with those observed in a rolled
AZ31 alloy; the comparison indicated that the strain to fracture in torsion
were significantly lower in ZM21 than in AZ31, even though in general in
the former alloy the ductility exceeded the minimum allowable value of
ε=1. These differences in mechanical response were attributed to a lower
tendency to dynamic recrystallization; while in rolled AZ31 early dynamic
recrystallization was observed at temperatures as low as 200°C,
microstructural observations clearly suggested that in the extruded ZM21
the lower limit for DRX was shifted toward significantly higher temperatures.
B)
29 - Metallurgical Science and Technology
growth, that precludes the fine grain size that
has been observed to be a prerequisite for
optimum sheet formability.
ACKNOWLEDGEMENTS
The authors greatly acknowledge the Ministries
of Foreign Affairs of Italy and Korea who in the
frame of “2004-06 Joint Scientific and Technological
Research Program” promoted the research by
funding the researchers’ mobility.
The valid support of D. Ciccarelli in experimental
activity has been appreciated.

REFERENCES
1) CAFÉ, http:/www.nhtsa.dot.gov
2) Communication from the Commission to the
Council and European Parliament:
Implementing the Community strategy to
reduce CO emissions from cars. Third annual
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report (reporting year2001), COM (2002)
693.
3) EU Directive on the End of Life of Vehicles,
2000/53/EC.
4) S. Schumann, Mat. Sci. Forum 488-489 (2005),
1-8.
5) Ya Zhang, X. Zeng, C. Lu, W. Ding, Y. Zhu, Mat.
Sci. Forum, 488-489,(2005), 123.
6) Magnesium Technology 2006, A.A. Luo et al,
eds. TMS, 2006.
7) F.J. Humphreys, M. Hatherly: Recrystallization
and related annealing phenomena. Pergamon
Press, Oxford, 1996.
8) J. Koike, Mat. Sci.Forum, 419-422 (2003), 199.
9) H. Utsunomiya, T. Sakai, S. Minamiguchi, H. Koh, Magnesium Technology
2006, A.A. Luo et al. eds. TMS, 2006, 201.
10) J. Bohlen, J. Swoistek, W.H. Sillekens, P.J. Vet, D. Letzig, K.U. Kainer,
Magnesium Technology 2005, TMS, 2005, 241.
11) S. Spigarelli, M. El Mehtedi, E. Evangelista, J. Kaneko, Metall.Sci.Techn. 23
(2005) 11-17.
12) S. Spigarelli, M. El Mehtedi, M. Cabibbo, E. Evangelista, J. Kaneko, A. Jäger,
V. Gartnerova, mater. Sci.Eng. A 462 (2007), 197.
13) A. Forcellese, M. El Mehtedi, M. Simoncini, S. Spigarelli, Key Engineering
Materials 344 (2007) 31.
14) S. Spigarelli, E. Evangelista, M. El Mehtedi, L. Balloni, Mechanical And
Microstructural Aspects Of High Temperature Formability Of AZ31
Sheets, Proceeding 10 ESAFORM 2007, Zaragoza, Spain, 1305.
15) H.J. McQueen, E.V. Konopleva, in Magnesium Technology 2001, J. Hryn
ed., TMS, 2001, p. 227.
30 - Metallurgical Science and Technology

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