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Journal of Materials Processing Technology 210 (2010) 767–775

Contents lists available at ScienceDirect

Journal of Materials Processing Technology

journa l homepage: www.e lsev ier .com/ locate / jmatprotec

ffect of conventional and rheocasting processes on microstructuralharacteristics of hypereutectic Al–Si–Cu–Mg alloy withariable Mg content

lireza Hekmat-Ardakan, Frank Ajersch ∗

cole Polytechnique de Montréal, Dép. de Génie Chimique, P.O. Box 6079, Centre-ville, Montreal, Quebec, Canada H3C 3A7

r t i c l e i n f o

rticle history:eceived 30 August 2009eceived in revised form 4 January 2010ccepted 9 January 2010

eywords:

a b s t r a c t

Hypereutectic A390 alloy (Al–17%Si–4.5%Cu–0.5%Mg (wt%)) and similar alloys with the addition of 6 wt%and 10 wt% Mg were investigated using conventional and rheocast processes. Conventional castingswere produced with two different cooling rates: −0.15 ± 0.05 ◦C s−1 and −1.0 ± 0.2 ◦C s−1. The rheo-cast samples were subjected to a rotation speed of 260 rpm during solidification with a cooling rateof −0.15 ± 0.05 ◦C s−1. The microstructure of the conventional casting for the high cooling rate sampleswas found to be finer than for the low cooling rate samples. With high Mg content, the primary sili-

luminium alloysheocastingutectic solidificationagnesium

asting

con transforms to primary Mg2Si and the nature of the eutectic reaction is also changed from a binary(Al + Si) to a ternary (Al + Si + Mg2Si) reaction. The morphology of eutectic silicon distributed in matrixmicrostructure significantly changes from large individual flakes to a fine, skeleton network with Chinesescript morphology, similar to the eutectic Mg2Si. This is caused by the decrease of the eutectic formationtemperature resulting from the addition of Mg. The matrix microstructure of the rheocast samples wasfound to contain fragmented eutectic silicon together with globular alpha aluminium phase particles.

tic sil

The morphology of eutec

. Introduction

Semisolid metal (SSM) processing can be widely used as a pre-erred fabrication route for near or net-shaped castings whereartially solidified metal with a globular structure (SSM slurry) is

njected into a mold, instead of liquid metal casting. SSM processings generally identified as either rheoprocessing or thixoprocess-ng. The rheo-route involves preparation of SSM slurry from theiquid phase by stirring the alloy during solidification and trans-erring directly into a die or mold for component shaping. Thehixo-route involves the preparation of a feedstock material with ane equiaxed or globular structure. The feedstock is then reheatedo a temperature between solidus and liquidus (mushy zone) inrder to produce a semisolid material which is injected into theold. A detailed review of SSM processing was described by Nafisi

nd Ghomashchi (2005).Alloys currently used for SSM processing are mainly conven-

ional Al–Si casting alloys. Hypereutectic Al–Si alloys with morehan 11.7 wt% Si offer the possibility of producing in situ metal

atrix composites (MMCs) with excellent wear and corrosionesistance where the silicon acts as the reinforcing phase, as indi-

∗ Corresponding author. Tel.: +1 514 340 4711x4533; fax: +1 514 340 4468.E-mail address: frank.ajersch@polymtl.ca (F. Ajersch).

924-0136/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.jmatprotec.2010.01.005

icon in rheocast samples becomes finer with increasing Mg content.© 2010 Elsevier B.V. All rights reserved.

cated by Kapranos et al. (2003). These composites are a class ofmaterials with an excellent combination of mechanical and physi-cal properties for various applications in the automotive industry.Baiqing et al. (2003) indicated that the high strength and wear resis-tance of these alloys are attributed to the presence of both hardprimary and eutectic Si particles.

Zhang et al. (2000, Material Science Engineering A) have alsoshown that hypereutectic Al–Si alloys with high Mg contentcan also be considered as in situ Al–Mg2Si MMCs with a highpotential for wear resistance because of the excellent physi-cal and mechanical properties of Mg2Si intermetallic phase thatis formed on solidification. The Mg2Si phase has a high melt-ing point (1085 ◦C), very low density (1.99 × 103 kg/m−3), highhardness (4.5 × 109 Nm−2), a low thermal expansion coefficient(7.5 × 106 K−1) and a reasonably high elastic modulus of 120 GPa(Jiang et al., 2005). Consequently, both Al–Si and Al–Mg2Si com-posite alloys can be used as wear resistant materials. In termsof properties and solidification behaviour, many similarities existbetween the Mg2Si and Si phase particles as compared by Zhang etal. (2000, Materials and Design).

In the present work, the effects of conventional and rheocastprocessing on the microstructural characteristics, particularly onthe morphology of eutectic silicon, of hypereutectic A390 alloy(Al–17%Si–4.5%Cu–0.5%Mg (wt%)) were investigated and com-pared to the microstructure of hypereutectic alloys with 6% and

768 A. Hekmat-Ardakan, F. Ajersch / Journal of Materials Processing Technology 210 (2010) 767–775

Table 1Chemical composition of test alloys and AZ91 Mg alloy (wt%).

Test alloy Al Si Cu Mg Fe Mn Zn P Ti

A390 Bal. 16.7 4.58 0.58 0.32 0.02 0.01 0.0003 0.026% Mg Bal. 16.83 4.12 6.08 0.28 0.04 0.015 0.0002 0.0210% Mg Bal. 17.14 4.32 9.74 0.25 0.05 0.02 0.0002 0.019

Zn

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Mg Al Mn

AZ91 Bal. 9.3 0.12

0% Mg. The addition of Mg and the formation of both primary andutectic Mg2Si phases can also affect the morphology of eutecticilicon. Briefly, this work investigates the effects of the chemicalomposition, the casting process and the effect of cooling rate onhe microstructure of alloy.

. Experimental procedures

The 6 wt% and 10 wt% Mg alloys were produced by adding AZ91Mg-base) alloy to A390 alloy. The chemical analysis of the testlloys and AZ91 Mg alloy is presented in Table 1.

In order to account for the oxidation loss, an additional amountf 20% Mg (as AZ91) was added to the melt. The AZ91 alloy wasrapped in aluminium foil and added to the liquid A390 alloy at

50 ◦C. Silicon and copper were also added to the melt in ordero attain the chemical composition of these elements at the samealue as for the A390 composition. The melt was degassed withrgon and stirred for 2 min before pouring into a steel mould toorm three ˚35 mm × 120 mm cylindrical samples.

.1. Conventional casting

After preparing the samples, specimens of 60 g portions were

ransferred into a graphite crucible with 28 mm inner diameter andmm wall thickness and heated to 700 ◦C using an electric resis-

ance furnace, degassed with argon and stirred for 2 min beforeampling. Samples were prepared with two cooling rates: highooling rate (HCR) and low cooling rate (LCR). For the HCR samples,

Fig. 1. (a) Schematic diagram of the apparatus used for the preparation

Si Ni Cu Fe

0.02 0.0006 0.0007 0.0046

the graphite crucible and contents were transferred to the coolingstation where a K-type thermocouple was quickly immersed intothe melt at a position about 10 mm from the bottom at the center ofthe crucible. Temperatures were recorded using a data logging sys-tem at 15 per second intervals. For the HCR cooling tests, the coolingrates averaged −1.0 ± 0.2 ◦C s−1. For the low cooling rate (LCR) tests,the samples were cooled in the furnace at a controlled cooling rateof −0.15 ± 0.05 ◦C s−1. Metallographic specimens for the conven-tionally cast samples were prepared by sectioning transversely ata level 10 mm from the bottom of crucible.

2.2. Rheocasting

Fig. 1a shows the schematic diagram of the apparatus used forrheocasting process. A spirally grooved cylindrical agitator wasused for producing the rheocast samples. Each rheocast test con-sisted of 150 g of alloy contained in a graphite crucible with aninner diameter of 41 mm, 10 mm wall thickness and a height of120 mm. Subsequently, the samples were heated to 700 ◦C and thenheld at this isothermal temperature for a period of 5 min. After thisperiod, the alloy was continuously cooled in the furnace at a rate of−0.15 ± 0.05 ◦C s−1 and sheared at an average rate of �̇ave = 52 s−1

(rotation speed = 260 rpm) as shown in Fig. 1a. Two K-type ther-

mocouples located in the crucible wall were used to measure thetemperature near the surface and near the bottom of the melt.Argon gas was introduced into the furnace chamber in order toprevent sample oxidation. When the upper limit of the viscometertorque was reached, the tests were stopped.

of rheocast samples and (b) the location of the sampled regions.

A. Hekmat-Ardakan, F. Ajersch / Journal of Materials Processing Technology 210 (2010) 767–775 769

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ig. 2. Cooling curves for A390, 6% Mg and 10% Mg samples for high cooling rateHCR) measured by a thermocouple at the center of graphite crucible during con-entional casting.

The rheocast specimens were prepared by sectioning the sam-les as shown in Fig. 1b. The specimens were polished with grindingaper up to 1200 followed by micro-polishing, using monocrystaliamond suspension (0.5 �m). All samples were etched by using.5% HF solution agent. Optical and scanning electron microscopySEM) with EDX (energy dispersive X-ray) analysis were carried outo characterize the microstructure. The Clemex software was usedo analyse the particle sizes observed in the microstructure.

. Results

.1. Conventional casting process

Fig. 2 shows the cooling curves for A390, 6% Mg and 10% Mg alloyamples for the air cooled (HCR) condition. The temperature waseasured by a thermocouple at the center of the melt during con-

entional casting. Fig. 3 compares the detail of the eutectic segmentf the cooling curves for HCR and LCR cooling of conventionally cast390, 6% Mg and 10% Mg alloys. The figure shows that for addi-

ions up to 6% Mg, the eutectic reaction temperature is reducedignificantly and becomes almost constant for additions between% and 10% Mg. Fig. 3 also indicates that the values of the eutecticemperature are very similar for both LCR and HCR conditions. Thealues of the eutectic reaction temperature obtained experimen-ally by cooling curves for both cooling conditions are presented inTable 2).

Fig. 4 shows the BSE (back scattered electron) image of A390lloy for the LCR conventional casting sample showing the formf the intermetallic compounds. The microstructure consists ofrimary polygonal silicon phase and the eutectic matrix consistsainly of binary Al + Si together with the intermetallic phases,

uAl2 (�), Cu2Mg8Si6Al5 (Q) and Al5FeSi (�). Other studies have alsobserved these microstructures with the presence of the Q phase

ssociated with the � phase (Li et al., 2006; Elmadagli and Alpas,006). However, the main intermetallic phase in these alloys is thephase.

It should be noted that tested alloys also contain small amountsf Fe, Zn, etc., which also react to form of intermetallics. The

able 2omparison of the eutectic reaction temperature determined by thermody-amic prediction using FactSage and the measurement of cooling curves at LCR−0.15 ± 0.05 ◦C s−1), HCR (−1.0 ± 0.2 ◦C s−1) for three alloys.

Predicted (◦C) Measured (◦C) LCR Measured (◦C) HCR

A390 566.2 563 ± 1 564 ± 16% Mg 549.7 545 ± 1 544 ± 110% Mg 549.2 543 ± 1 542 ± 1

Fig. 3. The eutectic segment of the cooling curves at HCR and LCR for (a) A390, (b)6% Mg and (c) 10% Mg alloys during conventional casting.

platelet Fe-containing phase, �-Al5FeSi, is the most common Feintermetallic phase and contributes to the decrease of the mechan-ical properties of Al–Si cast alloys as reported by Sreeja Kumari etal. (2007), Taghiabadi et al. (2008), and Nafisi et al. (2006). Withincreasing Mg content, the Mg2Si intermetallic phase first appearsin the matrix in the form of a eutectic morphology and then as aprimary phase. The microstructure of these three alloys for LCR and

HCR conditions are compared in Fig. 5. For both cooling conditions,the morphology of grey primary silicon particles remained polyg-onal, whereas Mg2Si crystals (dark particles) are either in the formof dendritic crystals with entrapped liquid as shown by Jiang et al.

770 A. Hekmat-Ardakan, F. Ajersch / Journal of Materials Processing Technology 210 (2010) 767–775

Fig. 4. BSE (back scattered electron) image of A390 alloy for conventionally cast low cooling rate (LCR) sample. (a) Large faceted primary silicon (1). The matrix consistsof binary Al + Si eutectic as well as the additional phases found in location A. (b) Magnified zone of location A, showing �-CuAl2 (2), Al (3), Cu2Mg8Si6Al5 (Q phase) (4) and�-Al5FeSi (5).

Fig. 5. Overall microstructure of A390 alloy (a and b), 6% Mg alloy (c and d) and 10% Mg alloy (e and f) for different cooling rates for conventionally cast samples. Images a, cand e correspond to LCR samples and b, d and f the HCR samples. The light grey phase is silicon, the dark phase is Mg2Si and the white phase is �-Al.

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A. Hekmat-Ardakan, F. Ajersch / Journal of Ma

2005) and Qin et al. (2007) or in the form of compact polygonalhapes.

The microstructure of the high cooling rate samples (Fig. 5b, d,nd f) is finer than that of the low cooling rate samples (Fig. 5a, c,nd e). Fig. 6 clearly shows that the eutectic morphology for bothhe HCR and LCR samples is significantly changed with the additionf Mg, particularly at 6% Mg when compared to A390 alloy. Theutectic silicon together with the dendritic �-Al shows a coarseake structure for A390 alloy cooled at both low and high ratesFig. 6a and b). However, the eutectic microstructure for 6% Mglloys (Fig. 6c and d) and 10% Mg alloys (Fig. 6e and f) is similar andhe eutectic silicon transforms into a fine, skeleton network withhinese script morphology, similar to the form of eutectic Mg2Si,s was also observed by Jiang et al. (2005).

The eutectic matrices of A390 and of the 6% Mg alloy at the HCRondition are compared in Fig. 7 using SEM images. The morphol-

gy of the eutectic silicon for A390 and 6% Mg alloys for the HCRondition (shown in Fig. 7) is also compared in the SEM image ofig. 8 where the surface was deep etched using 0.5% HF solution.hese figures clearly show that the eutectic silicon in the 6% Mglloy is much finer than in the A390 alloy.

ig. 6. The eutectic microstructure of A390 alloy (a and b), 6% Mg alloy (c and d) and 10%mages a, c and e correspond to LCR samples and b, d and f the HCR samples. Light grey p

s Processing Technology 210 (2010) 767–775 771

3.2. Rheocasting process

As described in the experimental procedure, the cooling condi-tion for the rheocast samples is similar to the LCR condition sincethe samples are stirred during continuous cooling solidification.

The tests were stopped at the upper limit of the viscometertorque. This maximum torque was attained at different temper-atures for each of the alloys. For A390 alloy this occurred at 541 ◦C(fliquid = 25.1%) and for the 6% and 10% Mg alloys the temperatureswere found to be 533 ◦C (fliquid = 26.3%) and 530 ◦C (fliquid = 23.3%),respectively. Therefore, the solidified structure of the remainingliquid should be similar to the conventional cooling structure with-out shearing effect at a solid fraction of about 25%.

A significant segregation of primary phases was detected inthe microstructure of the rheocast samples due to the shearforce imposed on the slurry during solidification. As a result, the

microstructure is different for two distinct regions. The region witha radial distance between 2 mm and 3 mm from the crucible wall(near crucible zone shown in Fig. 1b) is a region where the mostof the primary phase particles are concentrated, as shown in leftside images of Fig. 9 for A390, and the 6% and 10% Mg alloys. Wang

Mg alloy (e and f) alloys for different cooling rates for conventionally cast samples.hase is silicon, dark phase is Mg2Si and the white phase is �-Al grains.

772 A. Hekmat-Ardakan, F. Ajersch / Journal of Materials Processing Technology 210 (2010) 767–775

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ig. 7. Comparison of eutectic silicon phase (arrows) and microstructures of conagnification.

t al. (1994) have shown that during the solidification of hypereu-ectic Al–Si/SiC composite, the SiC particles in the composite melterve as the substrates for primary Si phase nucleation and all therimary Si particles form heterogeneously on the SiC surface. Inhe melt-stirred condition, all the Si particles become agglomeratednd are attached by the SiC particles. Sundarrajan et al. (1998) havelso explained this phenomenon based on the large undercoolingequired for the growth of polygonal Si crystals along low-energy,tomically smooth and slow-growing (1 1 1) planes. It was shownhat ceramic composites such as aluminium oxide, silicon carbidend graphite can eliminate such large undercooling and result inreferential nucleation and growth of Si on these materials. In thisase undercooling is eliminated as a result of capillary (interfa-ial) forces which bend the liquid–solid interface near the contactine of (1 1 1) facets with the composite surface. This phenomenonan explain the segregation of Si and Mg2Si particles resulting fornterfacial forces which are greater than the centrifugal forces evenhough the density of the liquid phase is higher than that of thearticles.

The region next to the agitator (near agitator zone shown inig. 1b) with a radial distance between 5 mm and 6 mm from agita-or, shows no primary phases. The globular �-Al grains observed inhis region form aggregates as shown in right side images of Fig. 9.

hen the primary phases in the region near the crucible wall areompared with non-sheared LCR samples in Fig. 5, it is clear that theegregation of primary phases caused by the shear forces is muchigher in the rheocast (sheared) samples. The primary silicon of

heared A390 alloy (Fig. 9a) also agglomerates. Fig. 10 indicateshat all three sheared samples show an increase of the size ofhe silicon particles when compared to the non-sheared sampless measured by image analysis. This implies that the shear pro-otes the coarsening of silicon particles. Lee et al. (1995) have also

ig. 8. Deep etched images showing the morphology of the eutectic Si in (a) A390 alloagnification.

nally solidified alloys at HCR condition for (a) A390 and (b) 6% Mg, at the same

reported that significant fragmentation and changes of primary Simorphology occur for an Al–15.5%Si alloy during the isothermalstirring but little change was observed when a sample was contin-uously cooled at a shear rate of �̇ = 200 s−1 and at a cooling rate of−0.03 ◦C s−1.

This shear rate however is able to fragment the dendritic mor-phology of the Mg2Si primary phase crystals with entrapped liquidas shown in Fig. 5c and e, forming smaller, compact polygonalshapes as shown in Fig. 9c and e. Fig. 11 compares the averageparticle size of the primary Mg2Si particles for the 6% and 10% Mgalloys for conventionally and rheocast samples. The reduction inparticle size is primarily due to the shear forces resulting in thefragmentation of the larger eutectic structure of the conventionallysolidified samples. In the region near the agitator, the agglomerated�-Al grains are present only in the matrix of the eutectic phaseas shown in Fig. 12a, c and e for each of these alloys. Fig. 12b,d and f represents the regions between the agglomerations con-sisting mostly of the eutectic region. For the A390 alloys (Fig. 12aand b), there is no significant difference in the microstructure of�-Al grains in matrix. However, for the high Mg-containing alloysof 6% and 10% Mg, respectively shown in Fig. 12c and d, and e and f,the microstructural difference is more pronounced. The �-Al grainsare generally surrounded by intermetallic phases, mainly by the�-phase.

It can be clearly observed that the eutectic silicon for the shearedA390 sample (Fig. 12b) becomes fragmented when compared to thenon-sheared sample shown in Fig. 6a. For the high Mg-containing

alloys, the skeleton network of eutectic silicon observed for theLCR condition, is completely destroyed and fragmented, as shownin Fig. 12d (6% Mg) and Fig. 12f (10% Mg). For the rheocast samples,the morphology of eutectic silicon becomes finer with increasingthe Mg content.

y and (b) 6% Mg alloy for conventionally HCR solidified cast samples at the same

A. Hekmat-Ardakan, F. Ajersch / Journal of Materials Processing Technology 210 (2010) 767–775 773

Fig. 9. The rheocast microstructure of A390 alloy (a and b), 6% Mg alloy (c and d) and 10% Mg alloy (e and f). Images a, c and e represent the area near the crucible wall andimages b, d and f represent the area near the agitator.

Fig. 10. Average particle size of primary silicon for conventional and rheocastingprocesses for three alloys.

Fig. 11. Average particle size of primary Mg2Si for conventional and rheocastingprocesses for alloys with 6% and 10% Mg alloys.

774 A. Hekmat-Ardakan, F. Ajersch / Journal of Materials Processing Technology 210 (2010) 767–775

F (a and b), 6% Mg alloy (c and d) and 10% Mg alloy (e and f). Images a, c and e showa cture between the agglomeration of �-Al grains.

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ig. 12. The microstructures near to the agitator for stir cast samples. A390 alloygglomeration of �-Al grains and images b, d and f illustrate the eutectic microstru

. Discussion

The results of this study have shown that Mg additions to hyper-utectic Al–Si alloy have a significant effect on the solidificationtructure of conventionally cast and rheocast alloys. As shown inig. 7, the change in morphology of eutectic silicon is due to theecrease of the eutectic temperature with the addition of Mg. Nafisind Ghomashchi (2006) also investigated the effect of the reductionf eutectic temperature on the modification of the eutectic sili-on morphology of A356 Al–Si alloy when Sr is added as modifier.n fact, they attributed this temperature reduction to the eutecticucleation temperature which has a direct effect on the number ofotential nuclei found in the melt. This produces fewer barriers forucleation with increasing nucleation temperature and generatesnumber of isolated eutectic Si particles, which results in the flake

orphology that was observed for A390 alloy.Fig. 13 compares the liquid fraction vs. temperature curves for

he three alloys according to Scheil solidification (no diffusion)sing the FACTSAGE software (Bale et al., 2002). The starting pointf the eutectic reaction, which occurs after the solidification of pri-

Fig. 13. The liquid fraction vs. temperature curves of A390, 6% Mg and 10 wt% Mgalloys calculated according to the thermodynamic predictions using the FACTSAGEsoftware for Scheil (no diffusion) solidification.

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A. Hekmat-Ardakan, F. Ajersch / Journal of Ma

ary phase, is indicated by arrows on the curves. This point is calledknee” point. For A390, the knee point represents the start of binaryeaction at 566.2 ◦C (liquid → Al + Si), whereas for 6% and 10% Mglloys, the knee point represents the start of ternary reaction at49.7 ◦C and at 549.2 ◦C, respectively, according to the followingeaction:

iquid → Al + Si + Mg2Si

The calculations show that the knee point temperatureecreases significantly for the 6% Mg alloy when compared to thealue for A390. This decrease is almost constant up to 10% Mg. TheACTSAGE calculation for Scheil solidification of A390 alloy alsohows that 68% of matrix microstructure is solid at 549 ◦C, whereasor 6% and 10% Mg alloys, this is the starting point for the solidifi-ation of eutectic matrix.

Table 2 shows the value of eutectic reaction temperature cal-ulated by thermodynamic prediction compared to the valuesbtained from the cooling curve tests for each of the three alloys.his table also indicates that the values of the eutectic tempera-ures are very similar for both LCR and HCR conditions. However,or both cases, the value of eutectic reaction temperature decreasesignificantly with the addition of 6% Mg from 566.2 ◦C to 549.7 ◦C.ncreasing the Mg content to 10% decreases the knee point temper-ture only slightly below the value for 6% Mg, to 549.2 ◦C.

For conventionally cast samples, the microstructural evolutionaused by the addition of Mg, particularly of the eutectic silicon inhe matrix (Fig. 8), should contribute to the improvement of wearehaviour of Al–Si–Cu–Mg alloys. This supposition needs moreomprehensive tests because the presence of a finer structure ofutectic phase generally results in higher wear resistance of theselloys.

. Conclusions

For conventionally cast samples, large and individual flake eutec-tic silicon phase particles were observed in A390 alloy for bothhigh cooling rate and low cooling rate conditions. The microstruc-ture of the high cooling rate samples is finer than for low coolingrate samples.With the addition of Mg, conventionally cast samples have shownthat the eutectic silicon particles form a fine, skeleton networkwith Chinese script morphology, similar to the eutectic Mg2Si.This is caused by the decrease of the eutectic formation temper-ature resulting from the addition of Mg.For the rheocast samples (rotation speed = 260 rpm, �̇ave =52 s−1), significant segregation was observed in the three sam-ples due to the shear forces imposed during the solidification.The microstructure is different for two distinct regions. Almost allprimary phases are concentrated in the region near the cruciblewall (outer layer) whereas the second region (inner layer) nextto the agitator shows no primary phases. This attributed to cap-illary forces between graphite crucible and the particles which

are greater than centrifugal forces. Therefore, in terms of wearresistance, the outer layer of samples becomes harder than theinner layer (near to the agitator).Agglomerated �-Al grains with a globular morphology arepresent in the eutectic phase matrix of the rheocast samples.

s Processing Technology 210 (2010) 767–775 775

An increase of the size of the silicon particles was observed forsheared samples confirming that shear promotes the coarseningof silicon particles. However, this shear force is also able to frag-ment the dendrite morphology of Mg2Si primary phase crystalsand allows them to form of polygonal shapes.

• Fragmentation of the eutectic silicon was also observed forthe sheared A390 sample when compared to the convention-ally solidified sample. With increasing Mg content, the skeletonnetwork of eutectic silicon was also completely destroyed andfragmented.

Acknowledgements

The authors gratefully acknowledge the financial support fromthe Fonds Quebecois de Recherche sur la Nature et les Technologies(FQRNT) and Natural Sciences and Engineering Research Council(NSERC) of Canada.

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