Solidification and Microstructural Evolution of Hypereutectic Al-15Si-4Cu-Mg Alloys with High Magnesium Contents

Solidification and Microstructural Evolution of HypereutecticAl-15Si-4Cu-Mg Alloys with High Magnesium Contents

M. TEBIB, F. AJERSCH, A.M. SAMUEL, and X.-G. CHEN

The low coefficient of thermal expansion and good wear resistance of hypereutectic Al-Si-Mgalloys with high Mg contents, together with the increasing demand for lightweight materials inengine applications have generated an increasing interest in these materials in the automotiveindustry. In the interests of pursuing the development of new wear-resistant alloys, the currentstudy was undertaken to investigate the effects of Mg additions ranging from 6 to 15 pct on thesolidification behavior of hypereutectic Al-15Si-4Cu-Mg alloy using thermodynamic calcula-tions, thermal analysis, and extensive microstructural examination. The Mg level stronglyinfluenced the microstructural evolution of the primary Mg2Si phase as well as the solidificationbehavior. Thermodynamic predictions using ThermoCalc software reported the occurrence ofsix reactions, comprising the formation of primary Mg2Si; two pre-eutectic binary reactions,forming either Mg2Si+Si or Mg2Si+ a-Al phases; the main ternary eutectic reaction formingMg2Si+Si+ a-Al; and two post-eutectic reactions resulting in the precipitation of theQ-Al5Mg8Cu2Si6 and h-Al2Cu phases, respectively. Microstructures of the four alloys studiedconfirmed the presence of these phases, in addition to that of the p-Al8Mg3FeSi6 (p-Fe) phase.The presence of the p-Fe phase was also confirmed by thermal analysis. The morphology of theprimary Mg2Si phase changed from an octahedral to a dendrite form at 12.52 pct Mg. Anyfurther Mg addition only coarsened the dendrites. Image analysis measurements revealed a closecorrelation between the measured and calculated phase fractions of the primary Mg2Si and Siphases. ThermoCalc and Scheil calculations show good agreement with the experimental resultsobtained from microstructural and thermal analyses.

DOI: 10.1007/s11661-013-1769-9� The Minerals, Metals & Materials Society and ASM International 2013

I. INTRODUCTION

HYPEREUTECTIC Al-Si alloys are widely used inautomotive and aerospace applications because of theirlow coefficient of thermal expansion, good wear resis-tance, and high hardness.[1–3] Alloys such as A390,containing 17 pct Si, 4.5 pct Cu, and 0.5 pct Mg, arewidely used in the manufacture of pistons, cylinderblocks, and AC compressors[3] (all alloy compositionsare given in weight percent (wt pct) unless otherwisestated). The presence of primary silicon in the matrixprovides good wear resistance and high hardness. In thisrespect, hypereutectic Al-Si alloys have often beenconsidered as in situ metal matrix composites (MMCs),with the primary Si acting as the reinforcement.[4–6]

However, the size and the morphology of the Si phasecan affect the mechanical properties of the alloy. Theaddition of a grain refiner and optimization of thecooling rate have been proposed to reduce the size ofsilicon particles and thus improve the mechanicalproperties of these alloys.[7]

Backerud et al.[8] studied the solidification of A390alloy. The solidification reactions reported by them arelisted in Table I. Solidification begins with the precip-itation of primary silicon. The second phase precipitatedis a small volume fraction of dendritic aluminum phase,followed by the eutectic reaction wherein Al, Si, andb-Al5FeSi phases are formed. Precipitations of Mg2Siand h-Al2Cu phases take place in later reactions, asshown in Table I.Hekmat-Ardakan et al.[9,10] and Mandal and

Makhlouf[11] observed that Mg addition changed themorphology and size of the primary silicon and theeutectic phase in hypereutectic A390 alloy, and led tothe formation of both eutectic and primary Mg2Si phaseparticles. Although the tensile strength of these alloyswas superior to that of A390 alloy, they exhibited alower hardness. Zeng et al.[12] examined the effects ofboth Si and Mg contents in Al-Si-Mg alloys, for Si andMg levels ranging from 8 to 17 pct and from 2 to 4 pct,respectively. The Mg addition ranging from 2 to 4 pctchanged the brittle b-Al5FeSi phase to thep-Al8Si6Mg3Fe phase, and improved the castabilityand mechanical properties at high Mg content. Linet al.[13] investigated the effects of Si and Mg contents onthe microstructures of hypereutectic Al-Si-Mg alloys.They reported that primary Mg2Si and primary Si areformed in the hypereutectic Al-Si alloy when thecontents of Si and Mg are equal to, or more than, 19and 4 pct, respectively. It was also found that increasing

M. TEBIB, Ph.D. Candidate, A.M. SAMUEL, and X.-G. CHEN,Professors, are with the Department of Applied Science, Universityof Quebec at Chicoutimi, 555, Boulevard de l’Universite, Sague-nay, QC G7H 2B1, Canada. Contact e-mail: [email protected]. AJERSCH, Professor, is with the Ecole Polytechnique deMontreal, Montreal, QC H3T 1J4, Canada.

Manuscript submitted December 6, 2012.Article published online April 30, 2013

4282—VOLUME 44A, SEPTEMBER 2013 METALLURGICAL AND MATERIALS TRANSACTIONS A

the amounts of Si and Mg increases the size of primarySi and primary Mg2Si particles.

More recently, Jayakumar et al.[14] evaluated the effectof Mg addition ranging from 3 to 5 pct on the solidifica-tion microstructures of homogenous and functionallygraded A390 aluminum alloys. The authors reported thatincrease in Mg content increases the concentration ofMg2Si, whereas percentage of primary silicon decreases.The size of the primaryMg2Si particles is smaller, and theyare observed to be distributed at the edges of the primarysilicon particles and also individually in the matrix.

Hekmat-Ardakan and Ajersch[10] reported that higherMg contents change the morphology of the eutecticsilicon from their long platelet-like form to a fine Chinesescript form, while the primary silicon particles aresuppressed and are replaced by primary Mg2Si phase.They also examined the wear properties of hypereutecticAl-Si-Cu-Mg alloys containing 6 and 10 pctMg in the as-cast condition and after T6 heat treatment, and comparedthe values to those of A390 alloy.[4] It was found that thewear resistance increases considerably for alloys withhigherMg content. From their investigations, the authorsconcluded that the improvement in wear resistance maybe attributed to the transformation of the coarse primarySi particles to finer Mg2Si phase particles.

Zhang et al.[15] have reported that Mg2Si has a densityof 1.99 9 103 kg m�3, a low thermal expansion coeffi-cient of 7.5 9 10�6 K�1, a high melting temperature of1358 K (1085 �C), a high hardness of 4.5 9 109 Nm�2,and a reasonably high elastic modulus of 120 GPa.These properties are important characteristics for thedevelopment of new lightweight and wear-resistantalloys containing Al and Mg2Si designed to replace thetraditional Al-Si hypereutectic alloys currently used inindustry.[16,17] Table II shows the close similarity inproperties between Mg2Si and Si particles.

According to Hekmat-Ardakan et al.,[9] the phasediagram calculations for Al-17 pct Si alloy with a min-imumMg content of 6.8 pct show that themicrostructurecontains only primary Mg2Si particles dispersed in theeutectic matrix. The lower density of Mg2Si (cf. 1.99 with2.33 g cm�3 for Si) provides an excellent opportunity forfurther weight reduction without loss in wear resistance.In view of these advantages, hypereutectic Al-Si alloyswith high Mg contents can also be considered as in situaluminum matrix composites containing a large amountof Mg2Si particles, and have garnered much interest inrecent years as potential candidates for critical applica-tions such as engine blocks and brake disks.Several studies[18–22] have reported on the methods of

modifying the Mg2Si primary phase in the Al/Mg2Siin situ composite; however, information on the solidifi-cation process of such composites produced by alloyinghypereutectic Al-Si with Mg is limited. Zhang et al.[23]

stated that microstructural changes in the Al/Mg2Sicomposite system were not explained clearly in previousinvestigations because of the very limited informationon the Al-Mg-Si ternary phase diagram, especially forhigh Mg and Si contents.Hekmat-Ardakan andAjersch[24] carried out a thermo-

dynamic investigation of hypereutectic Al-Si A390 alloyswith up to 10 pct Mg using the FactSage softwaredeveloped by the CRCT group at their institution. Theauthors reported that the addition of Mg up to 10 pctalters the formation temperature for the liquidus, and thestart of the binary and ternary eutectic reactions. TheMg2Si phaseprecipitates as an intermetallic compound forcontents up to 4.2 pct, and between 4.2 and 7.2 pct Mg.However, for compositions above 7.2 pct Mg, the Mg2Sisolidifies as the primary phase instead of Si.In the interests of pursuing the development of new

wear-resistant alloys further, and given the limitedinformation available to date, the current study wasundertaken to investigate the solidification behavior ofAl-Si-Cu-Mg quaternary alloys with Mg concentrationsup to 15 pct, beyond that investigated previously. Themicrostructures of these alloys were characterized usingoptical microscopy, scanning electron microscopy(SEM), and image analysis for quantification purposes.ThermoCalc calculations were also carried out todetermine the sequence of reactions and the phasesformed during the solidification process.

II. EXPERIMENTAL PROCEDURES

A. Thermodynamic Calculations

To better understand solidification and phase forma-tion mechanisms in the multicomponent Al-Si-Cu-Mgsystem, ThermoCalc, a software package for calculation

Table I. Solidification Reactions for A390 Alloy[8]

No. ReactionsSuggested

Temperature [K (�C)]

1 precipitation of primary Si crystal 927 (654)2 development of a-Al dendrites 901 (628)3 start of eutectic reaction

Liq. fi a-Al+Si+Al5FeSi838 (565)

4 precipitation of Mg2SiLiq. fi a-Al+Si+Mg2Si

813 (540)

5 precipitation of Al2CuL+Mg2Si fi a-Al+Si+Al2Cu+Al5Mg8Cu2Si6

773 (500)

6 precipitation of complex eutecticLiq. fi a-Al+Si+Al2Cu+Al5Mg8Cu2Si6

772 (499)

7 end of solidification 768 (494)

Table II. Physical and Mechanical Properties of Mg2Si and Si[15]

PhaseCrystalStructure

LatticeParameter (nm)

Density(g cm�3)

C.T.E.(10�6 K�1)

Elastic Modulus(GPa)

Melting Point[K (�C)]

Mg2Si cubic 0.635 1.99 7.5 120 1358 (1085)Si cubic 0.542 2.33 3.06 112 1684 (1411)

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 44A, SEPTEMBER 2013—4283

of phase equilibria based on the phenomenological orCalphad approach,[25] was employed to investigate theeffect of high Mg contents on the solidification sequenceof the hypereutectic Al-Si-Cu-Mg alloys studied. Ther-modynamic variables were computed using the AA12Brdatabase. The calculations were carried out for thequaternary system Al-Si-Cu-Mg, keeping the Si and Cucontents constant at 15 and 4 pct, respectively. Inaddition, the thermodynamic calculations were com-pared with the thermal analysis experimental results andmicrostructural observations obtained from the exper-imental data.

B. Alloy Preparation and Casting Procedures

A commercial A390 alloy and Al-50 pct Si and Al-50 pct Cumaster alloys were used as starting materials toprepare the Al-15Si-4Cu-xMg alloys. About 2.5 kg of theprepared alloy was melted in a graphite crucible, and thetemperature of the melt was maintained at 1053 ± 283 K(780 ± 10 �C). Small blocks of pure magnesium (99 pctpurity) preheated at 523 K (250 �C) and wrapped inaluminum foil were added to themelt with an extra 20 pctof the target composition to compensate for the lossthrough the oxidation of magnesium. The alloys weremelted in an electric resistance furnace under controlledargon atmosphere. Stirring was carried out after eachMgaddition. The melt was then poured into a small stainlesssteel crucible for thermal analysis. Samplings for chem-ical analysis were taken from the corresponding melts.Table III lists the chemical analyses of the four alloysstudied (coded A, B, C and D).

C. Thermal Analysis

To confirm the thermodynamic calculations, coolingcurves were obtained for the four alloys by placing twoK-type thermocouples at the center and the wall of apreheated stainless steel crucible into which the meltsample was poured, and recording the temperature–timedata during solidification using a data acquisitionsystem. The mass of the casting was about 150g andthe cooling rate was ~0.6 K s�1, calculated from thetemperatures above the liquidus.

D. Microstructural Characterization

In order to study the effect of Mg on the microstruc-tural evolution of the alloys studied, samples from thethermal analysis casting were sectioned from the centralpart, and then polished for microstructural examination

using standard metallographic procedures. Some sam-ples were also deep-etched for 8 minutes in 10 pctNaOH solution for better delineation of the 3D mor-phologies of the Mg2Si and Si particles. The micro-structures were examined using optical microscopy andthe various phases observed were identified using SEM(JEOL JSM-6480LV) equipped with energy dispersivespectroscopy (EDS). The volume fractions of Mg2Si andSi phases were quantified using an image analysis system(Clemex JS-2000) coupled with the optical microscope.

III. RESULTS AND DISCUSSION

A. Thermodynamic Predictions

The effect of Mg content on phase transformation andmicrostructural evolution during solidification of thehypereutectic Al-Si-Cu-Mg quaternary alloy system wasinvestigated by increasing the Mg level from 6.15 to20 pct, and keeping the Si and Cu contents constant at 15and 4 pct, respectively. The phase diagram was calcu-lated using ThermoCalc software and the supportingAA12Br data bank. Figure 1 summarizes the variousphase transformation reactions that occur in the hyper-eutectic Al-Si-Mg-Cu alloys during solidification overthe range of Mg concentrations studied. It should bementioned here that although the ThermoCalc can addiron to the alloy system, the resulting phase diagramobtained becomes quite complicated with the inclusion ofFe. Thus, in the interests of simplicity, only the quater-nary Al-Si-Cu-Mg system was considered in the thermo-dynamic calculations. However, while the amount of Fein the alloy is reached to 0.3 to 0.40 pct, it is sufficient toaccount for the formation of the p-Al8Mg5FeSi6 (p-Fe)phase which was observed in our study.For the range of Mg concentrations studied, the phase

diagram showed two distinct regions, I and II, separatedby a critical composition at 12.52 pct Mg, where region Irepresents Mg levels of 5.16 to 12.52 pct, and region II,from 12.52 to 20 pct Mg, corresponding to the verticaldotted lines in Figure 1. It was also found that within thisrange of Mg concentration, the solidification sequencechanged significantly. Based upon Figure 1, the sequenceof reactions and phases formed during solidificationunder equilibrium conditions is presented in Table IV.

Table III. Alloy Code and Chemical Compositions

of Experimental Alloys (Weight Percentage)

Alloy #

Elements (Wt Pct)

Si Mg Cu Fe Al

A 15.1 6.2 4.0 0.4 bal.B 15.5 8.4 4.0 0.4 bal.C 15.4 12.8 4.2 0.3 bal.D 15.5 14.9 4.0 0.3 bal.

Fig. 1—Calculated phase diagram of Al-Si-Cu-Mg system. Thedashed lines represent respectively the critical compositions ofAl-15Si-4Cu-xMg at 5.16 pct Mg and 12.52 pct Mg.


According to Figure 1, the liquidus, binary reaction,and ternary eutectic temperatures are found to varysignificantly with Mg addition. The liquidus temperatureincreases progressively from 864.4 K (591.4 �C) at5.2 pct Mg, to 963 K (690 �C) at 12.52 pct Mg, to1024 K(751 �C)at20 pct Mg;whereas thebinary reactiontemperature initially decreases from 864.4 K (591.4 �C) at5.2 pct Mg to 826 K (553 �C) at 12.52 pct Mg, and thenincreases to 845.5 K (572.5 �C) at 20 pct Mg. The tem-peratureof the ternary reaction, however, is not affectedbythe variation in Mg content. The phase diagram showsclearly that Mg2Si precipitates as the primary phase, andthen either Si andMg2Si, or a-Al andMg2Si, coprecipitatefrom the melt, depending on the Mg level in the alloy.

The liquid and solid phase fractions of each constit-uent at equilibrium conditions for the four alloys studied(i.e., four Mg levels) are shown in Figures 2(a) through(d). The solidification paths of these alloys can beclassified into three representative types (Table IV):

(i) Type I starts with the precipitation of primaryMg2Si particles, followed by a binary eutectic reac-tion where Mg2Si coprecipitates with the Si phase.

The main eutectic reaction is the ternary reactionL fi L+Si+ a-Al+Mg2Si, which produces ahigh amount of a-Al. This is followed by the qua-ternary eutectic reaction L fi Q+Si+ a-Al+Mg2Si, as shown in Figures 2(a) and (b). These twodiagrams show clearly how the phase fractions ofthe primary Mg2Si and Si particles vary with Mgcontent. Upon increasing the Mg concentration,the Mg2Si phase fraction increases linearly whilethe Si phase fraction decreases considerably.

(ii) Type II also begins with the precipitation of theMg2Si primary phase, followed by the ternaryeutectic reaction L fi L+Si+ a-Al+Mg2Si, andfinally by the quaternary eutectic L fi Q+Si+a-Al+Mg2Si reaction, as shown in Figure 2(c). Asignificant increase in the Mg2Si phase fraction canbe observed compared to alloys A and B. On theother hand, the second or binary eutectic reaction(L fi L+Mg2Si+Si or L fi L+ a-Al+Mg2Si)which occurred in alloys A, B and D is found tobe completely absent in the case of Alloy C, asshown in Figure 2(c).

Table IV. Reaction Scheme for Al-Si-Cu-Mg System According to the Present Work

*Reaction 6 was predicted only for Mg compositions up to 7 wt pct.


(iii) Type III is quite similar to the Type I solidifica-tion path corresponding to alloys A and B. Theonly difference lies in the second reaction wherethe binary eutectic reaction forming Mg2Si and Siis replaced by one forming Mg2Si and a-Al inAlloy D, as shown in Figure 2(d).

Table V lists the calculated liquid fractions, type/percentage of constituents formed at the transitiontemperatures, as well as the solid fractions correspond-ing to each reaction during the solidification of the fouralloys. The results indicate that the solid fraction at theeutectic reaction increases continually as the Mg contentrises from 7.7 to 31.2 pct for alloys A to D, as does the

Q phase fraction, which increases from 14.8 to15.76 pct; whereas those of primary Si and a-Al phasesdecrease respectively from 37.01 pct in Alloy A to7.61 pct in Alloy B.

B. Evolution of Microstructure

1. Precipitation of primary phasesFigures 3(a) through (d) display optical micrographs

of samples corresponding to alloys A through D with6.15, 8.42, 12.82, and 14.86 pct Mg, respectively. Themicrostructural analysis was supported by quantitativeimage analysis measurements of area fractions and the

Fig. 2—Calculated phase fractions for the four alloys studied, showing the amounts of the phases present under equilibrium conditions, as afunction of temperature.


average size of the black Mg2Si and gray Si primaryphase particles observed in these samples. All four alloysexhibit primary Mg2Si and Si particles, a ternaryeutectic region of Si/a-Al/Mg2Si, and intermetalliccompounds.

The microstructure of Alloy A in Figure 3(a) showsfine and uniformly distributed primary Mg2Si particlescoexisting with small polyhedral Si particles. Increasingthe Mg content increases the sizes of the primary Mg2Siand Si particles, Figure 3(b). Upon further increase inMg addition up to 12.82 pct, the primary Si particlesdisappear and are replaced by a-Al dendrites, as shownin Figure 3(c). Nami et al.[26] reported that the forma-tion of the a-Al phase on Mg2Si particles is due to thenonequilibrium solidification of the alloy, caused by thelimited diffusion rate of Mg and Si in the liquidsurrounding the Mg2Si particles.

It is of particular interest to note that (i) themorphology of the primary Mg2Si particles changesfrom a polyhedral form to a coarse dendrite crystalshape, and (ii) their size increases dramatically whencompared with the particles observed in alloys A and B,as shown in Figure 3(c). At the highest Mg content of14.86 pct (Alloy D), no further change in the morphol-ogy of the Mg2Si phase is observed; however, the Mg2Sidendrites coarsen considerably, while the primary Siparticles disappear, being replaced by the a-Al phase, asshown in Figure 3(d).

Figures 3(e) and (f) show micrographs of Alloys Aand B, respectively, taken at a higher magnification so asto render the primary particles as well as the eutecticregions much more clearly than the structures displayedin Figures 3(a) and (b) for the same alloys. The eutecticphases will be discussed in detail in Section III–B–3.

In order to better delineate the 3D morphologies ofthe primary Mg2Si and Si particles, the samples weredeep-etched using a solution of 10 pct NaOH. Examplesfor Alloy A and Alloy C are shown in the SEM

micrographs of Figures 4(a) and (b), which clearlyreveal how the relatively small, faceted polyhedralprimary Mg2Si particles in Alloy A acquire a massivedendrite form in Alloy C when the Mg level is increasedfrom 6.25 to 12.82 pct. This indicates that high Mglevels significantly influence both the nucleation andgrowth process of Mg2Si.

2. Growth of primary Mg2Si phaseThe high-magnification SEM images of Figure 5 show

three-dimensional views of particles of primary Mg2Siphase and its morphological evolution on proceedingfrom Alloy A to Alloy D, i.e., with increasing Mgcontent. At 12.82 pct Mg, the polyhedral morphology isconverted to a dendrite-like form, as exhibited by AlloyC. Further Mg addition only serves to expand thesedendrites in all directions. It should be noted here that allthese images were not taken at the same magnification,the objective being to demonstrate the morphologicaldetails as clearly as possible, without cutting off any partof the particle (or particles) shown in the image.Li et al.[27] carried out a detailed study on the

morphology and growth mechanism of primary Mg2Siphase in Al-Mg2Si composite alloys. They reported thatprimary Mg2Si particles may display different morphol-ogies, including octahedral, ‘‘hopper-like,’’ cubic, trun-cated octahedral, andmassive dendrite forms.Also, as theMgcontent is increased, the particle size increases, and themorphology changes. The images displayed in Figure 5show (a) an octahedral Mg2Si particle; (b) two Mg2Siparticles connected to each other and displaying partiallyincomplete octahedral surfaces or the so-called ‘‘hopper-like’’ structure; (c) the dendrite form of Mg2Si, with armsgrowing out in four directions; and (d) several Mg2Siarms, growing along the directions of dendrite branches.In the context of the latter, Figures 3(d) and 5(d) clearlydemonstrate the growth of several branches and second-ary arms from a primary Mg2Si dendrite center.

Table V. The Liquid and Solid Fractions Corresponding to the Reactions Occurring in the Four Alloys Studied

AlloyCode Reaction #* Reaction Temp [K (�C)]

LiquidFraction(Pct)

Constituent Percentage (At Pct) SolidFraction(Pct)Mg2Si Si a-Al Q

A 1 primary Mg2Si 882 (609) 100 — — — — 02 binary eutectic 860 (587) 97.8 100 — — — 2.24 ternary eutectic 827 (554) 92.3 62.98 37.01 — — 7.75 quaternary eutectic 812 (539) 2.3 — 8.94 76.2 14.86 97.65

B 1 primary Mg2Si 916 (643) 100 — — — — 02 binary eutectic 849 (576) 92.7 100 — — — 7.214 ternary eutectic 827 (554) 89.2 83.01 16.98 — — 10.735 quaternary eutectic 812 (539) 23.7 3.3 7.61 73.61 15.47 76.3

C 1 primary Mg2Si 966 (693) 100 — — — — 04 ternary eutectic 826 (553) 83.4 100 — — — 17.65 quaternary eutectic 812 (539) 23.9 10.06 4.15 69.03 15.75 76.1

D 1 primary Mg2Si 984 (711) 100 — — — — 03 binary eutectic 832 (559) 80.3 100 — — — 19.74 ternary eutectic 825 (552) 68.8 66.24 — 33.75 — 31.25 quaternary eutectic 812 (539) 24 13.99 3.95 66.3 15.76 76

*Reactions: (1) L fi L+Mg2Si; (2) L fi L+Si+Mg2Si; (3) L fi L+ a-Al+Mg2Si; (4) L fi L+Si+ a-Al+Mg2Si; (5) L fiQ-Al5Mg8Cu2Si6+Si+ a-Al+Mg2Si.


Wang et al.[28] have reported that Mg2Si crystals havean fcc structure, display a faceted growth, with preferredgrowth directions along h100i; the crystals exhibit a

tendency to appear in the form of faceted octahedrons,corresponding to minimum surface free energy condi-tions.[27] However, depending on the growth conditions

Fig. 3—Optical micrographs showing the typical structures obtained in the four alloys studied, containing (a) 6.15 pct, (b) 8.42 pct, (c) 12.82 pct,and (d) 14.86 pct Mg; (e), (f) High magnification micrographs of (e) Alloy A, and (f) Alloy B. Note the change in size, morphology, and areafraction of the primary Mg2Si (black) and Si (gray) particles.


and surrounding environment during the actual solidi-fication process, other morphologies may well result.The presence of certain elements, for example, may alterthe growth rate along the preferred growth directions, so

that certain planes get eliminated with the progress ofgrowth, others take precedence, ultimately altering themorphology of the Mg2Si crystals. Results of micro-structural study in the current article indicate that the

Fig. 4—SEM micrographs of deeply etched primary Mg2Si and Si particles in (a) Alloy A, and (b) Alloy C.

Fig. 5—SEM micrographs showing the morphological evolution of Mg2Si phase with increase in Mg addition: (a) Alloy A, (b) Alloy B, (c) AlloyC, and (d) Alloy D.


growth of the Mg2Si primary phase is strongly affectedby Mg levels which determine the final morphology ofthe Mg2Si primary phase.

3. Eutectic phasesThe micrographs of Figure 6 show the eutectic

network structures observed in the four alloys studied.The black Chinese script-type particles observed mostclearly in Figures 6(a) and (b) are those of the eutecticMg2Si phase—the dark gray acicular particles corre-spond to the eutectic Si phase, and the lighter dendrite-shaped areas represent the a-Al phase, respectively. Thepresence of copper and iron in these alloys also leads tothe precipitation of the h-Al2Cu and p-Al8Mg5FeSi6(p-Fe) phases, respectively.

Several morphologically distinct intermetallic con-stituents were identified with the aid of SEM coupledwith EDS. These intermetallic species, labeled 3, 4,and 5 in Figure 6, were respectively identified as the p-Fe, h-Al2Cu, and Q-Al5Mg8Cu2Si6 phases. The b-Al5FeSi phase, which appears in the form of plateletsin Al-Si alloys at low Mg contents, is completelysuppressed and is replaced by the p-Fe phase. Thepresence of all these constituents is clearly noted in allfour alloys.

A comparison of the eutectic structure in these alloysreveals a significant refinement in the size of the eutecticsilicon particles with increasing Mg content. As acomparison of Figures 6(a) and (d) reveals, the mor-phology of the eutectic Si particles is also changed fromits regular plate-like form to fine script-like-type struc-tures. The present findings also show that the additionof Mg leads to an abundant formation of the p-Feintermetallic phase and a consequent increase in the sizeof these particles with increasing Mg content from 6.15to 14.86 pct.

C. Image Analysis

A quantitative analysis of the microstructural evolu-tion of these alloys as a function of Mg addition wasalso carried out using image analysis. Figure 7(a) showsthe mean particle size of the primary Mg2Si and Siphases. It was found that the average Mg2Si particle sizeincreased from 17.64 lm in Alloy A at 6.15 pct Mg to727.66 lm at 14.86 pct Mg in Alloy D. Also, the averageSi particle size increased from 67.87 lm in Alloy A to429.68 lm in Alloy B. In Alloys C and D, no pre-eutectic Si particles were observed, which is in goodagreement with the thermodynamic predictions reported

Fig. 6—Optical micrographs obtained from the eutectic regions of the four alloys in the as-cast condition: (a) Alloy A, (b) Alloy B, (c) Alloy C,and (d) Alloy D, showing the presence of (1) eutectic Mg2Si, (2) eutectic Si, (3) p-Fe, (4) h-Al2Cu, and (5) Q-Al5Mg8Cu2Si6.


in Section III–A. Figure 7(b) compares the predicted(ThermoCalc) and the measured fractions of theprimary Mg2Si and Si phases. As may be observed,the Mg2Si phase fraction increases progressively withincreasing Mg content, while the Si phase fractiondecreases for the same Mg contents in the case of AlloysA and B. Also, while the calculated and measured phasefractions for Si show good agreement, the measuredphase fractions of the primary Mg2Si phase are largerthan the corresponding calculated values. The discrep-ancy may be explained by the fact that, unlike theprimary Si particles which have a comparatively welldefined shape, the primary Mg2Si particles are moreirregularly shaped, particularly at higher Mg levels,when they exhibit dendritic rather than polyhedralmorphology (Figure 3).

In the context of image analysis measurements,where the particles being measured are filled in with aspecific color (corresponding to a selected bit plane), it

is difficult to accurately outline the contour of suchirregularly shaped dendrite particles. Consequently, thearea or phase fraction measured would represent morethan what is actually present. Overall, the measuredand calculated Mg2Si phase fractions show the sametrend. Interestingly enough, the two values are veryclose in the case of Alloy A with the lowest Mg content,while the values for Alloy D display the largestdifference.

D. Thermal Analysis

Thermal analysis was carried out for the four alloys,using a cooling rate of 0.6 �C s�1. Figure 8 shows thecooling curves and the first derivatives obtained ineach case. The derivatives clearly reveal six peaks inthe case of Alloys A and B, but only four peaks forAlloy C, and five peaks for Alloy D, because of thevariation in the solidification sequence in accordancewith the alloy composition. Based on the phasediagram obtained using ThermoCalc and the micro-structural observations, the possible reactions corre-sponding to each individual peak were identified andare listed in Table VI.For Alloys A and B, Peak 1 corresponds to the

formation of primary Mg2Si particles from the liquidphase. Peak 2 and Peak 3 correspond to the (Mg2Si andSi) binary eutectic and the (a-Al, Si, and Mg2Si) ternaryeutectic reactions, respectively. Peak 6 represents theprecipitation of the Cu-rich intermetallic phase (h-Al2Cu). The only contentious identifications are forthose associated with Peaks 4 and 5. In order todetermine the temperatures of formation of the twophases that give rise to Peaks 4 and 5 in the coolingcurves and derivatives, an interrupted water-quenchingtechnique was employed.Figure 9 shows the microstructures of the samples

obtained from an Alloy A melt, and quenched (i) froma melt temperature that ranged between 801 K and793 K (528 �C and 520 �C), and (ii) from 783 K(510 �C), respectively. The microstructure of the samplethat was water quenched at 783 K (510 �C) showed theexistence of both Q-Al5Mg8Cu2Si6 and p-Fe phases, asseen in Figure 9(a). However, on examining the micro-structure of the samples quenched between 801 K and793 K (528 �C and 520 �C), shown in Figure 9(b), onlythe formation of p-Fe was observed. Figure 10 presentsa backscattered electron image obtained from Alloy A,showing a script-like particle within the rectangulararea, and the corresponding X-ray mappings of theelemental distributions of Si, Mg, and Fe in the area.The morphology, X-ray images, and EDS analysis ofthe Chinese script-like particle in the inset in (a)confirmed this to be the p-Fe phase. These resultsindicate that the p-Fe phase precipitates before the Q-Al5Mg8Cu2Si6 phase, so that Peaks 4 and 5 in Figure 8correspond to the formation of the p-Fe and Q phases,respectively.The first derivative of the cooling curve plotted in

Figure 8(c) for Alloy C (12.82 pct Mg) shows fourpeaks—Peaks 1, 3, 5 and 6, corresponding to theformation of primary Mg2Si, the ternary eutectic

Fig. 7—Effect of Mg content on the (a) mean particle size, and (b)phase fractions of primary Mg2Si and Si phases in hypereutectic Al-Si-Cu-Mg alloys.


containing a-Al, Si, and Mg2Si, and the precipitation ofthe Q-Al5Mg8Cu2Si6 and h-Al2Cu phases, respectively.In the case of Alloy D, Figure 8(d), the first two hightemperature peaks correspond to the formation ofprimary Mg2Si, and the (Mg2Si+ a-Al) binary eutecticreaction, as predicted by the phase diagram. Peaks 3and 5 reflect the formations of the (Mg2Si+Si+ a-Al)ternary eutectic and the Q-phases, while the last peakindicates the precipitation of the h-Al2Cu phase. In thecase of Alloys C and D, although the reaction for thep-Fe phase is not that evident, the change in the slopeof the first derivative curve between Peaks 3 and 5,nevertheless, indicates that the reaction does occur atthat point. The fact that the p-Fe phase is observed inthe microstructures of all four alloys confirms theexistence of this peak in Figures 8(c) and (d). In allprobability, the amount of the p-Fe phase producedand the corresponding exothermic energy released areinsufficient to produce a more pronounced peak as inthe cases of Alloys A and B which contain slightlyhigher Fe levels (0.4 vs 0.3 pct in Alloys C and D). It isinteresting to note that Peak 5 increases with increasing

Mg content, as more Mg is available to form the Q-phase.It can be summarized from the above that the major

reactions in all the cases can be identified by the phasediagram, from the microstructural analysis of the phasesand the thermal analysis results. The data obtained fromthermal analysis are in good agreement with thecalculated phase diagram and the observed microstruc-tures. The only exception is that although the precipi-tation of h-Al2Cu phase at 768 K (495 �C) was notpredicted in the phase diagram for alloys C and D, itwas observed in both the microstructures and in thecooling curves (Peak 6).In view of the difficulty in direct visualization of the h-

Al2Cu phase in the equilibrium phase diagram forcompositions higher than 7 wt pct Mg, an alternativecalculation using the Scheil solidification simulation wasalso carried out. This method calculates the solidifica-tion path of an alloy under nonequilibrium conditions,which are close to the real solidification conditions, i.e.,where there is no diffusion in the solid phases andthermodynamic equilibrium exists only locally, at the

Fig. 8—Temperature vs time cooling curves and their first derivatives obtained for (a) Alloy A, (b) Alloy B, (c) Alloy C, and (d) Alloy D.


liquid/solid interface.[29] Figure 11 shows the calculatedScheil solidification path for Alloys B and D, where thesolidification temperature was plotted against a functionof the solidified weight fraction. It can be seen that the h-Al2Cu phase precipitates in the last reaction, at ~773 K(500 �C) in both Alloys B and D.

Table VI presents a summary of the solidification data.It should be noted that the temperatures obtained fromthe cooling curves and listed in Table VI, are accuratewithin ±275 K (±2 �C). For comparison purposes, thepredicted temperatures (as obtained from ThermoCalc)are also listed the fourth column of the table.

Table VI. Main Reactions Observed from the Thermal Analysis of Alloys A, B, C, and D Used in the Current Study

AlloyCode

ReactionNumber

TemperatureMeasured [K (�C)]

Temperaturepredicted [K (�C)] Type of Reaction Solidification Stage

A 1 861 (588) 882 (609) L+Mg2Si primary phase2 828 (555) 860 (587) L+Si+Mg2Si binary eutectic3 811 (538) 827 (554) L+Si+ a-Al+Mg2Si ternary eutectic4 801 (528) — L+ p-Fe+Si+ a-Al+Mg2Si quaternary eutectic5 793 (520) 812 (539) L+Q+Si+ a-Al+ p-Fe+Mg2Si6 768 (495) 638 (365) Al2Cu+Q+Si+ a-Al+p-Fe+Mg2Si

B 1 899 (626) 916 (643) L+Mg2Si primary phase2* 875 (602) 849 (576) L+Si+Mg2Si binary eutectic3 806 (533) 827 (554) L+Si+ a-Al+Mg2Si ternary eutectic4 273 (525) — L+ p-Fe+Si+ a-Al+Mg2Si quaternary eutectic5 793 (520) 812 (539) L+Q+Si+ a-Al+ p-Fe+Mg2Si6 768 (495) Al2Cu+Q+Si+ a-Al+p-Fe+Mg2Si

C 1 951 (678) 966 (693) L+Mg2Si primary phase2 805 (532) 826 (553) L+Si+ a-Al+Mg2Si ternary eutectic4** — — L+ p-Fe+Si+ a-Al+Mg2Si quaternary eutectic5 792 (519) 812 (539) L+Q+Si+ a-Al+ p-Fe+Mg2Si6 767 (494) — Al2Cu+Q+Si+ a-Al+p-Fe+Mg2Si

D 1 980 (707) 984 (711) L+Mg2Si primary phase2 809 (536) 832 (559) L+ a-Al+Mg2Si binary eutectic3 804 (531) 825 (552) L+Si+ a-Al+Mg2Si ternary eutectic4** — — L+ p-Fe+Si+ a-Al+Mg2Si quaternary eutectic5 787 (514) 812 (539) L+Q+Si+ a-Al+ p-Fe+Mg2Si6 765 (492) — Al2Cu+Q+Si+ a-Al+p-Fe+Mg2Si

*The predicted temperature of Reaction 2 in Alloy B is lower than the measured temperature.**The precipitation of p-Fe phase—Peak 4 in the first derivative curves of Alloys A and B could not be clearly observed in Alloys C and D.

However, the microstructures of the four alloys showed the formation of this phase.

Fig. 9—Microstructure of Al-15Si-4Cu-6.15Mg alloy (Alloy A) water quenched at different temperatures: (a) at 783 K (510 �C), and (b) between801 K and 793 K (528 �C and 520 �C).


IV. CONCLUSIONS

The effect of the addition of Mg (6 to 15 pct) on thesolidification behavior and microstructural evolution ofhypereutectic Al-15Si-4Cu-Mg alloys was investigatedusing thermodynamic calculations and thermal analysis,accompanied by extensive microstructural examination.It was found that the presence of Mg strongly influencesthe microstructure and solidification characteristics.From the analysis of the results obtained, the followingconclusions were drawn.

1. Thermodynamic calculations using ThermoCalcsoftware revealed the occurrence of six reactionsduring the solidification of the Al-Si-Cu-Mg alloysinvestigated, comprising (i) formation of primaryMg2Si, two eutectic binary reactions which pro-duced (ii) Mg2Si+Si or (iii) Mg2Si+ a-Al; (iv) themain ternary eutectic reaction with the formation ofMg2Si+Si+ a-Al; and two quaternary eutecticreactions showing the formation of (v) the Q-Al5Mg8Cu2Si6 phase; and (vi) the h-Al2Cu phase.

2. Although the ThermoCalc equilibrium phase dia-gram predicts the formation of the h-Al2Cu phaseonly up to 7 wt pct Mg levels, the use of Scheil(nonequilibrium) conditions in calculating the solid-

ification path show the presence of the h-Al2Cuphase at higher Mg levels.

3. Depending on the Mg level, the solidification fol-lows three paths: Types I and III resemble eachother except in their binary eutectic reactions. TypeII shows the absence of any binary eutectic reactionat the critical Mg content of 12.52 pct, which sepa-rates the phase diagram into two distinct regions.The liquidus, binary, and ternary eutectic tempera-tures also vary based on the Mg content.

4. Microstructures obtained from all alloys (Fe con-tent ranging from 0.3 to 0.4 pct) showed the pres-ence of the p-Al8Mg3FeSi6 (p-Fe) phase in additionto all the other phases. The presence of the p-Fephase was also detected in the cooling curvesobtained from thermal analysis.

5. The growth of the Mg2Si primary phase is stronglyaffected by Mg addition. As the Mg content is pro-gressively increased, the morphology of the primaryphase evolves from an octahedral to a faceted den-drite form. This change in morphology occurs at12.52 wt pct Mg. Further increase in Mg contentonly coarsens the dendrite Mg2Si primary particles.

6. The results predicted from ThermoCalc and Scheil cal-culations show good agreement with the experimental

Fig. 10—(a) Backscattered electron image of Alloy A and corresponding X-ray images showing distributions of (b) Si, (c) Mg, and (d) Fe in thep-Fe phase shown in the inset in (a).


data obtained from the microstructural and thermalanalyses.

ACKNOWLEDGMENTS

The authors would like to acknowledge the financialsupport provided by the Natural Sciences and Engi-neering Research Council of Canada (NSERC) andRio Tinto Alcan through the NSERC-Rio Tinto Alcan

Industrial Research Chair in Metallurgy of AluminumTransformation at University of Quebec at Chicoutimi.

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Fig. 11—Calculated Scheil solidification paths for (a) Alloy B with8.4 pct Mg, and (b) Alloy D with 14.86 pct Mg.


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Solidification and Microstructural Evolution of Hypereutectic Al-15Si-4Cu-Mg Alloys with High Magnesium Contents