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Effect of a short solution treatment time on microstructureand mechanical properties of modified
Al–7wt.%Si–0.3wt.%Mg alloy
D.L. Zhang a,*, L.H. Zheng b, D.H. StJohn b
a Department of Materials and Process Engineering, University of Waikato, Private Bag 3105, Hamilton, New Zealand b CRC for Cast Metals Manufacturing (CAST), Department of Mining, Minerals and Materials Engineering,
University of Queensland, Queensland 4072, Australia
Abstract
Microstructural change caused by a short solution treatment and the corresponding change in tensile properties and impact
energy of a strontium modified Al–7wt.%Si–0.3%Mg cast alloy were studied. It was found that a solution treatment of 10 min at 540
or 550 °C is sufficient for the a-aluminium phase to homogenise and achieve the maximum level of magnesium and silicon as
predicted by the solubility and alloy composition limits. A solution treatment of 30 min causes spheroidisation, coarsening and an
increase in inter-particle spacing of the eutectic silicon particles leading to a significant improvement in ductility and impact re-
sistance. Compared with a standard 6 h solution treatment, solution treatment of 30 min at 540 or 550 °C is sufficient to achieve
more than 90% of the maximum yield strength and more than 95% of the maximum UTS and the maximum average elongation to
fracture. However, only 80% of the maximum impact energy can be attained by the short solution treatment. The values of the
ductility and impact energy pass through a minimum between 1.5 and 10 min of solution treatment time indicating that solution
treatments of less than 10 min should be avoided.
Ó 2002 Elsevier Science Ltd. All rights reserved.
Keywords: Aluminium casting alloys; Heat treatment; Mechanical properties; Microstructural change; Castings
1. Introduction
When cast components for structural applications
such as alloy wheels are manufactured using Al–Si–Mg
based casting alloys (typically A356 and A357), T6 heat
treatment is in most cases an essential step in the man-
ufacturing process. The T6 heat treatment provides two
beneficial effects: an improved ductility and fracture
toughness through spheroidisation of the eutectic silicon
particles in the microstructure and a higher alloy yieldstrength (YS) through the formation of a large number
of fine b00 precipitates which strengthen the soft alu-
minium matrix. The first benefit is realized through the
solution treatment (normally at a temperature around
540 °C) while the second benefit is achieved through
the combination of solution treatment, quenching and
artificial ageing (at a temperature in the range of 140–
170 °C) [1–3]. In the casting industry, it is often specified
that a cast component should be solution treated for 6 h
at 540 °C.
While the benefit of T6 heat treatment is accepted, the
additional cost and production time associated with
such a treatment is also substantial. Taking a cast
component made by a low pressure die casting process
as an example, the casting process normally takes less
than 10 min, while a typical T6 heat treatment cycle may
take more than 10 h. This means that shortening thetotal time of the T6 heat treatment cycle has a major
impact on productivity and manufacturing cost. For this
reason, there is a strong interest in establishing the
feasibility of shortening the solution time. Shivkumar
et al. [4] demonstrated that for permanent mould cast
test bars of a modified A356 alloy, a solution treatment
of 50 min at 540 °C is sufficient to attain more than 90%
of the maximum YS, more than 95% of the ultimate
tensile strength (UTS) and nearly 90% of the maximum
elongation for a given ageing condition. In agreement
with the observation on the tensile property change,
Journal of Light Metals 2 (2002) 27–36
www.elsevier.com/locate/lightmetals
* Corresponding author. Tel.: +64-7-838-4783; fax: +64-7-838-4835.
E-mail address: [email protected] (D.L. Zhang).
1471-5317/02/$ - see front matter Ó 2002 Elsevier Science Ltd. All rights reserved.
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they also observed that the magnesium and silicon
contents in the a-aluminium dendrites reached the max-
imum equilibrium level according to the alloy compo-
sition and the distribution of magnesium and silicon
became homogeneous within 50 min of solution treat-
ment at 540 °C. It is therefore clear that for permanent
mould castings of a modified A356 alloy, a solution
treatment of 50 min is sufficient to attain at least 90% of
the maximum tensile properties of the alloy. Although
the maximum tensile properties are not achieved, the
dramatic reduction of the solution time may offer an
opportunity to increase productivity and reduce cost
while maintaining the performance of the components.
Full modification of the as-cast structure appears to
be an essential precondition for the feasibility of a short
solution treatment. It has been well established that the
fibrous eutectic silicon phase in the modified structure is
fragmented and spheroidised much more rapidly than
the plate shaped silicon particles in the unmodified
structure [1,4–8]. With a completely unmodified micro-structure, solution treatment of 1–2 h at 540 or 550 °C
has little effect on the morphology of the eutectic silicon
particles, while with a fully modified microstructure the
effect of short solution time of 1–2 h on the eutectic
silicon particles is much more significant [4,5,7,8].
In this study, the effect of even shorter (<30 min)
solution treatment time at 540 or 550 °C on the micro-
structure and properties of a fully modified Al–7wt.%Si–
0.3wt.%Mg alloy was investigated. This study focuses
on examining the changes in microstructure and me-
chanical properties during the very early stages of so-
lution treatment (0–30 min).
2. Experimental method
The sample bars used in all the solution treatment
experiments except those for producing impact test sam-
ples, were cut from the rim region of an Al–7wt.%Si–
0.3wt.%Mg alloy wheel cast by low pressure die casting.
The alloy was modified by using 0.015 wt.% Sr.
The actual composition of the alloy was determined to
be Al–6.7wt.%Si–0.26wt.%Mg–0.12wt.%Fe–0.02wt.%Ti–
0.015%Sr by using the inductively coupled plasma
technique. The approximate dimensions of the bars were12 Â 15 Â 55 mm3. The bars were solution treated in a
high temperature salt bath held at 540 or 550 °C. The
advantage of using a salt bath was that the heating rate
was high, and thus the isothermal holding time formed
the major part of the solution treatment, which ranged
from 2 to 30 min. The temperature of the bars was
monitored during solution treatment by using a ther-
mocouple embedded in one of the bars and a comput-
erised data logging system. Fig. 1 shows a typical
heating curve of the samples during solution treat-
ment. From the heating curve, it is observed that it took
approximately 30 s to heat the samples from room
temperature to 505 °C. Due to fast heating and fairly
low temperature, the solution treatment effect experi-enced by the samples during this heating stage is negli-
gible, so this heating time was discounted. The solution
treatment time quoted was the net time that the sample
temperature is above 505 °C. Fig. 1 shows that shortly
after reaching 505 °C the samples reach 540 or 550 °C.
After each solution treatment, the samples were
quenched in a hot water bath held at 60 °C. The quen-
ched samples were then immediately aged in an air cir-
culated furnace at 140 °C for 4 h. This is a typical
underageing treatment used in the manufacture of
wheels. To prepare for tensile testing, round specimens
were machined from the heat treated samples. The ten-
sile tests were performed using a screw driven Instron
tensile testing machine. The cross-head speed used was 1
mm/min. The strain was measured by using an extenso-
meter attached to the sample and with a measuring
length of 10 mm. The 0.2% proof stress was used as the
measure of yield stress. Five samples were tested for
each heat treatment condition.
The microstructure of the as-cast and heat treated
samples was examined using an optical microscope and
quantitatively analysed using an image analyser (QM
750). To quantify the microstructural change during
solution treatment, the emphasis of the image analysis
was placed on the aspect ratio, equivalent circle diam-eter of the eutectic silicon particles and the inter-particle
spacing. Each measurement included 800–1200 particles
obtained from several areas inside the eutectic regions.
The iron-rich intermetallic and Mg2Si phases were de-
liberately excluded from the measurements. In order to
easily observe changes to the silicon morphology the
data were grouped around the selected values of 2 for
the aspect ratio and 1 lm for the particle size. The
magnesium and silicon contents in the a-aluminium
matrix were measured using an electron probe micro-
analyser (JEOL JXA-8800L) with an operating voltage
Fig. 1. Sample temperature as a function of time during solution
treatment in a salt bath held at 540 °C.
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than 2 as a function of solution time at 540 °C. By in-
creasing solution time from 1.5 to 19.5 min, the fraction
of the silicon particles with an aspect ratio of less than 2
increased slightly from 0.63 to 0.71. A further increase
of the solution time did not result in any further increase
in the number fraction of the silicon particles with a low
aspect ratio of less than 2. This shows that prolonged
solution treatment beyond 20 min at 540 °C had little
effect on the extent of spheroidisation. For a given short
solution treatment time, increasing the solution tem-
perature from 540 to 550 °C clearly increased the
number fraction of the silicon particles with low aspect
ratio. As shown in Fig. 4(b), for the same solution
treatment time of 9.5 min, the number fraction of thesilicon particles with an aspect ratio of less than 2 in-
creased by approximately 10% when the solution tem-
perature increased from 540 to 550 °C. With a longer
solution time of 19.5 min, this effect was much less sig-
nificant.
Fig. 5(a) shows the number fraction of the silicon
particles with an equivalent circle diameter of less than 1
lm and greater than 1 lm as a function of solution time.
By increasing the solution treatment time at 540 °C from
1.5 to 19.5 min, the number fraction of the silicon par-
ticles with a diameter of greater than 1 lm almost
doubled. With a further increase of the solution time,
the number fraction of the silicon particles with a dia-
meter greater than 1 lm did not increase any further.
Instead, it clearly decreased. This unusual change in the
number fraction of silicon particles with a diameter of
greater than 1 lm might be caused by the relatively large
number of small residual particles produced by the
dissolution process of many silicon particles which is an
essential process of coarsening [10]. As shown in Fig.
5(b), with a short solution treatment time of 9.5 min,
increasing the temperature from 540 to 550 °C increased
the number fraction of silicon particles with a diameter
of greater than 1 lm by more than 10%. This shows that
the rate of coarsening at the beginning of solution
treatment increases with temperature.Fig. 6(a) and (b) show the change in the average dia-
meter of the silicon particles and the average inter-
particle spacing as a function of solution time at 540 and
550 °C. The average diameter of silicon particles and the
average inter-particle spacing increased rapidly within
the first 10 min at the solution temperature. Then they
increased slowly with further increase in solution time. It
should be noted that the slight decrease in average dia-
meter and inter-particle spacing by increasing the so-
lution time from 19.5 to 29.5 min was due more to the
uncertainty of the image analysis technique and incon-
Fig. 4. The number fraction of silicon particles with an aspect ratio of
<2 and >2 (a) as a function of solution time at 540 °C and (b) cor-
responding to two solution treatment times at 540 and 550 °C.
Fig. 5. The number fraction of silicon particles with an equivalent
circle diameter of <1 lm or >1 lm (a) as a function of solution time at
540 °C and (b) corresponding to two solution treatment times at 540
and 550 °C.
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sistency of the samples than reflecting an actual trend.
After a prolonged solution treatment of 6 h at 540 °C,
the average diameter of the silicon particles increased
to 2.6 lm, while the average inter-particle spacing in-
creased to 7.8 lm.
3.1.2. Homogenisation
Fig. 7(a) and (b) show the typical distribution of
magnesium and silicon content across the width of an
a-aluminium dendrite in as-cast and solution treatedsamples. In the as-cast condition, the magnesium con-
tent is distributed fairly homogeneously with a compo-
sition of approximately 0.15 wt.% which is substantially
lower than the equilibrium level of 0.3 wt.%. After the
samples were solution treated for 9.5 min, the magne-
sium content increased to 0.3 wt.% and its distribution
became very homogeneous. In the as-cast condition, the
silicon content was significantly higher near the centre of
the dendrite arms than at their edge, as observed pre-
viously by Shivkumar et al. [4] and Closset et al. [11].
After 1.5 min at 540 °C, the distribution of silicon was
still inhomogeneous but after 9.5 min the distribution of
silicon was observed to be homogeneous. With the
higher solution temperature of 550 °C, it took a shorter
time for the distribution of silicon to become homoge-
neous (Fig. 7(b)).
Fig. 8 shows the change of average magnesium andsilicon contents across a-aluminium dendrites as a
function of solution treatment time at 540 and 550 °C.
The average magnesium content in the a-aluminium
phase increased significantly in the first 1.5 min of so-
lution treatment at 540 or 550 °C. However, the distri-
bution of magnesium was still fairly inhomogeneous as
reflected by the large variation of its value from one
dendrite to another. By increasing the solution time to
9.5 min, the average magnesium content reached the
equilibrium level of 0.3 wt.%, and the distribution be-
came very homogeneous. The trend was the same for
Fig. 6. (a) Average diameter and (b) inter-particle spacing of the silicon
particles as a function of solution treatment time at 540 and 550 °C.
Fig. 7. The distribution of magnesium and silicon across the width of
an a-aluminium dendrite arm in the as-cast condition and after 1.5 and
9.5 min of solution treatment at (a) 540 and (b) 550 °C.
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the solution time beyond 9.5 min and up to 29.5 min, the
average elongation again fluctuated around 10%. The
average elongation corresponding to a solution treat-
ment time of 6 h at 540 °C was 10.5%. This shows that
with a short solution treatment time of 30 min, 95% of
the maximum average value of the elongation to frac-
ture can be achieved.
Fig. 11 shows the impact energy obtained from the
Charpy impact tests as a function of solution treatment
time at 540 °C. The impact energy decreased with in-
creasing solution treatment time from 0 to 5.5 min, and
then gradually increased up to 11.5 min. By further in-creasing the solution treatment time beyond 11.5 min
and up to 29.5 min, the impact energy fluctuated around
4 J. The average impact energy of the samples solution
treated for 6 h at 540 °C, was 4.9 J. This shows that with
a short solution treatment time of 30 min, the alloy can
achieve approximately 80% of the maximum impact
energy. The difference in impact energy between the
samples that were solution treated for a short time and
those that were solution treated for the standard long
time of 6 h, is much more substantial than the differ-
ences measured for strength and ductility.
4. Discussion
To determine whether a short solution treatment time
is feasible, it is essential to be clear on whether the two
functions of the solution treatment can be achieved
within a short solution treatment time. As mentioned
above, the two intended functions are: (1) to raise the
magnesium and silicon solute contents to the maximum
level and to homogenize their distribution; and (2) to
sufficiently reduce the aspect ratio and increase the size
and spacing of the eutectic silicon particles. The first
function is essential for achieving the maximum level of
YS corresponding to the alloy composition and the
ageing condition used through precipitation hardening.
The second function is necessary for improving the
ductility of the alloy from the as-cast state. These two
functions are largely independent of each other. The
UTS is improved when both the yield and ductility are
improved, as has been confirmed by Taylor et al. [14]
through an empirical analysis of trends in mechanical
properties of T6 heat treated Al–Si–Mg casting alloys.
4.1. Yield strength
By consideration of both tensile properties and solute
content and distribution, this study establishes that, at
least for an Al–7wt.%Si–0.3wt.%Mg alloy, the majority
(>90%) of the potential YS is achieved after a short
solution treatment of 10 min at 540 °C. This is princi-
pally due to achieving homogenisation of the a-alu-minium dendrites with equilibrium levels of magnesium
and silicon within 10 min of solution treatment (Fig. 8).
Shivkumar et al. [4] reported that 25 min solution
treatment at 540 °C is sufficient for achieving a similar
result. However, it is important to note that the ‘‘10
min’’ referred to in this study is the net time at 540 °C,
while the ‘‘25 min’’ quoted in Shivkumar et al.’s paper
may be the total time the sample was in the furnace (not
clearly specified in their paper). In establishing our ex-
perimental method, it was found that in an air circulated
furnace, a small sample of 12 Â 12 Â 60 mm3 took 5–6
min to heat to a temperature above 520°
C. In this sense,the results of this study are probably in reasonable
agreement with those reported by Shivkumar et al. [4].
Although the majority of the YS is realised by hold-
ing the samples at 540 °C for 10 min, there still exists a
5–10% difference in strength between a short solution
treatment of 10 min and the standard long solution
treatment (6 h in this study). Yield strength is very
sensitive to the magnesium content in the matrix [15],
indicating that the actual average magnesium content
over the whole sample is possibly a little less than that
calculated from the microprobe measurements that were
taken on a relatively small number of scans.
The conclusion that only a short solution treatmenttime of less than 10 min is needed for realising sufficient
YS may only be applicable to alloys with low magne-
sium contents (e.g. 0.3–0.4 wt.%) that have been cast at
relatively high cooling rates, typical of those obtained in
the rim of low pressure die cast wheels. If the level of
magnesium is higher (up to 0.7 wt.%), a longer solution
treatment time will be required, as predicted by the so-
lution and homogenisation model for A365 and A375
alloys developed by Rometstch et al. [16]. Since solidi-
fication rate has a dramatic effect on the size and dis-
tribution of Mg2Si particles and the size of the other
Fig. 11. The impact energy of the modified Al–7wt.%Si–0.3wt.%Mg
alloy as a function of solution treatment time at 540 °C.
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possible magnesium containing phases, it is important to
be aware of the influence of the solidification rate on the
required minimum solution time for realizing the re-
quired YS. A separate study [17] indicates that a short
solution treatment of 10 min at 540 °C was also suffi-
cient for attaining greater than 90% of the strength for
the samples cut from the hub region of the wheels of Al–
7wt.%Si–0.3wt.%Mg alloy. The solidification rate in the
hub region was substantially lower than the rim region.
Therefore, a short solution treatment of 10 min may be
sufficient for most low pressure die castings or perma-
nent mould castings with a composition of magnesium
not greater than 0.4 wt.%.
Although it appears that the time required for
achieving sufficient YS is very short, it is important that
this time requirement is met at the solution treatment
temperature. This study and the study undertaken by
Shivkumar et al. [4] both show that the average mag-
nesium content in the as-cast a-aluminium dendrites is
substantially lower than the equilibrium level (i.e. 0.15compared with 0.3 wt.%), even though the distribution
of magnesium appears to be homogeneous. This would
make the strength achieved by the so-called simplified
solution treatment [18], where a casting is quenched
directly after the casting process and followed by arti-
ficial aging, to be substantially lower than that allowed
by the alloy composition. In addition, the strength ob-
tained by the simplified solution treatment may be very
sensitive to the casting condition which influences the
magnesium content and its distribution in the a-alu-
minium. This is a critical issue, since a manufacturer
generally needs consistency of strength and hardness
throughout their castings.
Increasing the solution treatment temperature to 550
°C accelerates the solution and homogenisation process,
so 10 min at 550 °C is more than sufficient to achieve
optimum YS for the Al–7wt.%Si–0.3wt.%Mg alloy. This
was confirmed by microstructural examination and
tensile property measurement.
4.2. Ductility
This study shows that a solution treatment of 30 min
at 540 °C causes substantial spheroidisation and coars-ening of the silicon particles in a fully modified Al–
7wt.%Si–0.3wt.%Mg alloy. This microstructural change
has an effect on the ductility of the alloy, resulting in a
substantial increase in the elongation to fracture. It is
interesting and important to note that the ductility of
the alloy achieved by a solution treatment of 30 min is
almost the same as that obtained by the standard solu-
tion treatment time of 6 h. Previous work [4,7,9] estab-
lished that with a fully modified as-cast microstructure,
a solution treatment of 1 h at 540 or 550 °C leads to a
significant degree of spheroidisation and coarsening of
the eutectic silicon particles. This study further confirms
the previous observations and shows that the spheroidi-
sation and coarsening achieved with a much shorter
time at 540 or 550 °C is also significant.
It is striking to observe that the average fracture
strain of the tensile test samples solution treated for
a short time of less than 8 min is substantially lower
than that corresponding to a solution treatment time of
longer than 10 min. This feature is shown for both so-
lution temperatures, indicating it is unlikely to be caused
by experimental uncertainty or microstructural incon-
sistency of the test samples. This observation suggests
that there exists a region during the early stage of so-
lution treatment (0–8 min), where the ductility of the
alloy reaches a minimum level. The cause of this region
is likely to be a timing mismatch between an increase in
strength and the improvement of the features of the
silicon particles brought about by the solution treat-
ment. With very short solution treatment time, the
strength rapidly increases to its nearly maximum leveldue to the fast solution and homogenisation kinetics.
However, over this short time the silicon particles have
only begun to spheroidise and thus the elongated silicon
fibres are more likely to fracture when the YS increases
[19]. Therefore it is possible that the ductility does not
begin to increase until most of the silicon particles begin
to approach a spheroidal morphology. Spheroidisation
and coarsening (to some extent) of the eutectic silicon
particles in the cast alloy increases the fracture strain, as
they make it more difficult for the silicon particles to
fracture [19–21]. To achieve a high confidence in the
improvement of ductility brought by solution treatment,
this region should be avoided, and therefore the solution
time used should be more than 10 min. There exists a
similar and slightly wider region of low impact energy as
a function of solution time at 540 °C (Fig. 11). Again to
avoid this region, a solution treatment time of greater
than 10 min at 540 or 550 °C should be used. It is noted
that the recently published work by Pederen and
Arnberg [22] also showed a similar phenomenon.
The results of this study show that after 30 min of
solution treatment at 540 °C, about 80% of the maxi-
mum impact energy value is achieved. The difference is
much larger than that for tensile ductility. The spheroi-
disation and coarsening of the silicon particles and theincrease in inter-particle spacing corresponding to a
short solution treatment time of 20 min appear to be
insufficient to achieve the maximum value of the impact
energy. Increasing the solution time to 30 min results in
no further improvement. The impact energy reflects the
ease of crack nucleation and growth at high strain rate.
It is likely that the inter-particle spacing plays a domi-
nant role in determining how easily cracks nucleate and
grow at a very high strain rate. The nucleation of the
cracks is likely to start with cracking of the brittle silicon
particles during impact, as has been shown by previous
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observations on the more slowly deformed tensile sam-
ples [18]. Once a large fraction of silicon particles are
cracked, cracks grow by linking microvoids formed by
the cracking of the silicon particles [23]. With a smaller
inter-particle spacing, it is easier for the microvoids to
link and grow a crack. This substantial difference in the
impact energy value between short and long solution
treatment times is a concern if the short solution treated
cast components are used in a service environment
where resistance to mechanical impact is a basic expec-
tation. The tolerance of a casting to a low impact re-
sistance and the minimum solution time required to
reduce the difference to 5–10% could be important fac-
tors in determining the practical feasibility of using a
short solution treatment time for some applications.
Another possible concern is the difference in fatigue re-
sistance between a short and a long solution treatment
time. This is a subject for further investigation.
4.3. Practical implications
The practical implications of the studies on short
solution treatment times are clear. As long as the as-cast
microstructure is appropriate, a short solution treatment
of less than 30 min at 540 or 550 °C can be used in many
cases where impact resistance is not a prime concern.
The essential precondition is that the as-cast micro-
structure of the Al–7wt.%Si–(0.25–0.45)wt.%Mg alloy
must be fully modified. This ensures fast spheroidisation
and coarsening, allowing a short solution treatment to
be implemented.
While this may not be very important for a long so-
lution treatment, it is important to know the exact time
at the intended solution temperature when a short so-
lution treatment of 30 min or less is used. To avoid the
low ductility region, longer than 10 min at temperature
must be experienced by every critical region of the cast
component. While it is not difficult to ensure this in
laboratories where small and regular samples are heat
treated, it may be a challenging task for practitioners in
manufacturing plants where castings of complex shape,
and large numbers have to be dealt with every day.
5. Conclusions
For a low pressure die cast Al–7wt.%Si–0.3wt.%Mg
Sr-modified alloy, a solution treatment of 10 min at 540
or 550 °C is sufficient to obtain the maximum level of
magnesium and silicon in the a-aluminium phase. Si-
multaneously, homogenisation of the magnesium and
silicon in the a-aluminium phase is also achieved. A
solution treatment of 30 min causes spheroidisation,
coarsening and an increase in inter-particle spacing of
the eutectic silicon particles leading to a significant im-
provement in ductility and impact resistance.
It is established that for the alloy, casting and ageing
conditions studied, a short solution treatment of 30 min
at 540 or 550 °C is sufficient to achieve more than 90%
of the maximum YS and more than 95% of the maxi-
mum UTS and the maximum average elongation to
fracture compared with a solution treatment time of 6 h.
However, only 80% of the maximum impact energy can
be attained by the short solution treatment.
There exists a region where the elongation to fracture
and the impact energy decrease to a minimum before
increasing. This region corresponds to 1.5–10 min of
solution time. The cause of this region is probably due
to a mismatch between the negative effect of solution
treatment on ductility and impact resistance associated
with a rapid increase in YS and the more slowly devel-
oping positive effect associated with the spheroidisation
and coarsening of silicon particles.
Acknowledgements
The authors thank Southern Aluminium Pty Ltd,
Comalco Aluminium and David Farnsworth for their
support for this project and Dr. Carlos Caceres for
useful discussion. The CRC for Cast Metals Manufac-
turing (CAST) was established under and is supported
by the Australian Government’s Cooperative Research
Centre Scheme.
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