MATERIALS CHARACTERIZATION 40:145158 (1998) Elsevier Science Inc., 1998 1044-5803/98/$19.00655 Avenue of the Americas, New York, NY 10010 PII S1044-5803(98)00004-7

145

M

icrostructure, a Limiting Parameter for Determining the Engineering Range of Compositions for Light Alloys: The Al-Cu-Si System

Diego Plaza,* Juan Asensio,

Jose A. Pero-Sanz,* and Jose I. Verdeja

*Universidad Politcnica de Madrid, E.T.S.I. de Minas, 28003 Madrid, Spain; and

Universidad de Oviedo, E.T.S.I. de Minas, 33004 Oviedo, Spain

Twelve as-cast alloys of the Al-Cu-Si ternary system were investigated. In all the cases, themicrostructural phases observed were:

a

solid solution of Cu and Si in Al, CuAl

2

(

u

phase),and silicon crystals. The morphology and distribution of the

u

and Si brittle constituentslimit the percentages of Cu and Si added in the composition ranges of these commercial

alloys. Elsevier Science Inc., 1998

INTRODUCTION

The factors affecting the selection of com-positions of aluminum-based light alloysfor castings can be summarized as follows:adequate castability, specific properties inservice (mechanical, corrosion resistance,etc.), and economic factors.

It is said that an alloy has adequate casta-bility when, departing from the liquid state,it is possible to obtain sound componentsthrough solidification; that is, good repro-ducibility of the geometry of the mold (ca-stability),

compaction

(without porosities dueto entrapped gases and microshrinkagecavities), and no cracking during cooling.The fluidity of materials that solidify with aroughly planar growth morphology is sim-ply controlled by progressive solidification.These conditions are seen in pure materialsand eutectic alloys and give rise to thegreatest fluidities. These materials are knownas short-freezing-range materials [1].

Complementary considerations such asthe mechanical properties and others, usu-ally determined by the metallographic struc-ture derived from the solidification con-ditions, could require the use of othercompositions.

EXPERIMENTAL PROCEDURES

Although an enormous number of aluminumcasting alloys have been developed to date,the major ones used industrially can be sum-marized into six major alloy types: Al-Si, Al-Cu, Al-Cu-Si, Al-Mg, Al-Zn-Mg, and Al-Sn.

The present work deals with the threefirst types (Al-Si, Al-Cu, and Al-Cu-Si) and,based on the microstructure and the corre-lation with its mechanical properties, anumber of recommendations are made forthe selection of the compositional ranges inthe ternary system. Twelve alloys havebeen used, and their compositions areshown in Table 1. Figure 1 shows the loca-tion in the ternary Al-Cu-Si system of theindividual compositions in accordancewith the method proposed by Roozebum[2]. Figure 2 shows the commercial rangesof the individual Al-Cu and Al-Si systemsplotted on the ternary diagram.

Solidification of the alloys was conductedat very low cooling rates to get equilibriumstructures. The curves

T

versus

dt

/

dT

weredetermined during the solidification range,and the points of start and finish of solidifi-cation were calculated. The results are pre-sented in Fig. 3. Metallographic obser-

146

D. Plaza et al.

vation was conducted after mechanicalpolishing and etching in water-based solu-tions containing variable proportions (110vol.%) of NaOH. Some of the specimenswere etched with FeCl

3

in alcoholic solu-tion to reveal the CuAl

2

constituent in bi-nary and ternary eutectics.

RESULTS

Al-Si BINARY SYSTEM

The solubility of silicon in aluminum is al-most negligible at room temperature. The

maximum solubility (1.65 wt.% Si) is at-tained at 577

8

C. At this temperature, thereis a eutectic at 12.5 wt.% Si consisting ofsolid solution

a

(1.65% Si in solid solutionof Al) and Si, as shown in the diagram inFig. 4. This eutectic is called abnormal fortwo reasons: first, because of the great dif-ference between the melting points of eachof the individual constituents forming theeutectic (1414

8

C for Si and 660

8

C for Al)and second, because of the difference in therelative proportions of such constituents(88.91 wt.% of

a

phase and 11.09 wt.% ofSi). In consideration of a density of 2.33 and

Table 1

Composition of the SamplesExpressed in Weight Percent

Alloy samplenumber Al Cu Si

1 93 0 72 88 0 123 85 0 154 96 4 05 90 10 06 65 35 07 84 10 68 78 10 129 69.2 25.5 5.3

10 63.5 30 6.511 61 35 412 60 35 5

FIG. 1. Position of the alloys studied in the Al-Cu-Siternary diagram.

FIG. 2. Projection of the Al-Cu-Si diagram.

FIG. 3. Thermal analysis curves.

The Al-Cu-Si System

147

2.669, for Si and Al, respectively, it followsthat the volume of the aluminum in the eu-tectic is approximately seven times as bigas that of the silicon.

The effective

dominance

during the eutec-tic solidification corresponds to silicon be-cause of its higher driving force with re-spect to aluminum. The driving force isproportional to the thermal gradient be-tween the melting point of Si and that ofthe eutectic (1414577

8

C). As a result, thesolidification kinetics of silicon are fasterthan for aluminum. Therefore, Si starts tonucleate and grow in the liquid before theAl starts to solidify in the correspondingproportion to the eutectic.

The free growth of Si in the liquid tendsto form

polyhedral

forms before it becomessurrounded by eutectic Al to form the eu-tectic cells. The planar surface orientationof the polyhedrals corresponds to {111}, al-though, less frequently, the orientation ofthe faces could be parallel to less-compactplanes such as {100} and {211}. The size ofthe polyhedrals and therefore of the eutec-tic cells is obviously smaller as the temper-ature of solidification increases. For thisreason,

metallic molds

are preferred to sandmolds in Al-Si castable alloys because ofthe improved toughness achieved at highersolidification rates.

In the Si crystals, differently from the ma-jority of metals, the solidliquid interfacedoes not move isotropically or normal to it-self, but instead by a lateral displacement.This produces an acicular Si morphology,in particular for the smaller thermal gradi-ents during cooling (Figs. 5 and 6). Thisanisotropic growth is due to the formationof twins between {111} planes such that theplanes of twinning align parallel to the

,

211

.

direction, which is the preferentialdirection of growth [3].

Silicon crystals can occasionally show adendritic morphology (Fig. 7), and this isobserved in the following cases: under se-

FIG. 4. Al-Si binary diagram.

FIG. 5. Micrograph of alloy sample 2: 88%Al-12%Si (sand casting).

148

D. Plaza et al.

vere thermal gradients, for high undercool-ings, and in hypoeutectic alloys. There isevidence that the primary constituent,

a

,favors the early formation of

a

-eutecticcrystals by heterogeneous nucleation [4].

In addition, in regard to hypereutectic al-loys in Si, the primary Si crystals promotethe formation of the

a

phase by heteroge-

neous nucleation. This is thought to be con-nected to the epitaxial relations between Siand the planes {110} and {111} in Al [3].However, when hypo- and hypereutecticalloys are compared, the

a

primary crystalsare better nucleation sites for eutectic sili-con than the primary silicon crystals are forthe heterogeneous nucleation of eutectic

a

.

FIG. 6. Micrograph of alloy sample 2: 88%Al-12%Si (die casting).

FIG. 7. Micrograph of alloy sample 1: 93%Al-7%Si.

The Al-Cu-Si System

149

This explains the small sizes of silicon platesfound in the eutectic matrix observed in thehypoeutectic alloys.

At the end of the solidification process,the micrographic structure of a binary eu-tectic alloy, named Silumin or Alpax [5],consist of silicon crystalswith a polyhe-dral or acicular morphology or bothdis-persed in an aluminum matrix (Figs. 5 and6). When the whole is subjected to mechan-

ical stresses, an interfacial decohesion oc-curs between Al and the elongated Si parti-cles, followed by (or simultaneously with)the fracture of silicon because Si crystals donot deform at temperatures lower than630

8

C [6]. The fracture of silicon explainsthe low toughness of the eutectic Al-Si al-loy. Complementarily, the difference be-tween the value of the linear coefficient ofthermal expansion in pure Al in the range

FIG. 8. Micrographs of Modified Silumin: (a) general and (b) detail.

150

D. Plaza et al.

of 25450

8

C (25

m

m

/

m

/8

C) and that of Si (al-most eight times smaller) also explains itsusual fracture mode due to thermal fatigue.

It has been determined that by control-ling the solidification rate of Silumin, it ispossible to obtain smaller eutectic cells andan improvement in toughness. Figures 5and 6 show the microstructure obtained forSilumin solidified by sand casting and diecasting, respectively. In a previously pub-lished work [7], the effectiveness of theTi

3

Al and B

2

Ti additions for the refinementof the eutectic cells was studied. This re-finement takes place through heteroge-neous nucleation that provokes a modifica-tion of the microstructure, which can alsobe attained by additions of Sr, Ca, Sb, andso forth.

The possibility of improvement in themechanical properties by Na additionsmade to the melt before the start of pouringis well known. Supercooling permits theeutectic to solidify, giving a fine dispersionof small silicon particles in an aluminummatrix [Fig. 8(a, b)]. This alloy is namedModified Silumin and exhibits much bettertoughness properties than does the unmod-ified alloy. In fact, at 427

8

C, the modified al-loy exhibits superplastic behavior owing to

the spherical shape and coalescence of thesilicon crystals [3].

Al-Si binary alloys are used commer-cially when a light alloy is demanded withgood castability and corrosion resistance.Its usual composition is in the range of 5 to12.5 wt.% Si. With increasing contents of Si,the total density decreases, its castabilityimproves (shortens the freezing range), andthe toughness level decreases (because ofthe increase in the volume fraction of eutec-tic). When wear resistance is required, thehypereutectic alloys are usually preferred(Fig. 9) with additions up to 22 wt.% Si, toget the benefit from the higher volume frac-

FIG. 9. Micrograph of alloy sample 3: 85%Al-15%Si.

FIG. 10. Al-Cu binary diagram.

The Al-Cu-Si System

151

tion of primary silicon crystals. However,in this range the addition of P becomes es-sential to refine the primary silicon, thusimproving the castability without exces-sively impairing the mechanical character-istics [8].

Al-Cu BINARY SYSTEM

Copper is added to aluminum to increaseits mechanical properties by heat treatment,

solution treatment, and ageing. The addi-tion of Cu to Al also impairs its corrosionbehavior. In fact, the alloy composition Al-4%Cu was the first [9] heat-treatable single-phase alloy obtained by supersaturation andageing to provide a fine precipitation, whichhighly contributes to the total strengthen-ing. From the phase diagram observation inFig. 10, it follows that Al takes in solid solu-tion up to 5.77 wt.% Cu. When higheramounts of copper are added, between 5.77

FIG. 11. Micrographs of alloy sample 5: (a) 90%Al-10%Cu and (b) detail of the eutectic phase.

152

D. Plaza et al.

and 33.2 wt.%, two-phase structures are ob-served in the solidified structure. Figure11(a, b) illustrates the microstructure of speci-men 5 (Al-10%Cu), which is hypoeutecticand consists of primary

a

phase (solid solu-tion of Al-5.77%Cu) surrounded by a matrixof the eutectic (Al-33.22%Cu) of

a

Al andCuAl

2

[10]. The

a

phase is white (unetchedand reflective) in Fig. 11, whereas the inter-metallic CuAl

2

is dark in appearance.

Application of the lever rule to the Al-Cuphase diagram for the eutectic predictsweight fractions of 41.2%

a

phase and58.8% CuAl

2

phase. The volumetric ratio ofAl to CuAl

2

is given by:

RVAl

VCuAl2-----------------

mass Alr Al

-------------------------

mass CuAl2r CuAl2

------------------------------------

------------------------------------ 1.14= = =

FIG. 12. Micrographs of alloy sample 6: (a) 65%Al-35%Cu and (b) detail of the a-CuAl2 eutectic phase.

The Al-Cu-Si System

153

Because this ratio is close to unity, thevolume fractions of each phase should beabout the same under equilibrium condi-tions. This also applies to the eutectic mix-ture in both hypo- and hypereutectic com-positions [see Figs. 11(b) and 12(b)], wherethe proportions of

a

and CuAl

2

in the eu-tectic are approximately the same.

When the solidification takes place undernonequilibrium conditions, certain types ofsingle-phase alloys can show the presenceof the eutectic phase. Figure 13 shows anexample of the latter in an Al-4%Cu com-position. The nonequilibrium eutectic has ahigher proportion of CuAl

2 compared withthe equilibrium one [compare Fig. 13 withFig. 12(a, b)]. The nonequilibrium eutecticscan be eliminated by heat treatment; and,under industrial practices, this heat treat-ment becomes a compulsory stage beforethe ageing treatments. However, the eutec-tic in alloys with more than 5.7 wt.% Cucannot be eliminated by heat treatment.

The hypoeutectic copper alloys (Cu ,32.2 wt.%) have higher amounts of eutecticphase when the amount of added Cu in-creases, thus improving the castability butimpairing the toughness of the alloy. Theintermetallic CuAl2 is hard and brittle, justi-

fying the lack of toughness in ordinary con-ditions of these alloys, although it exhibitsa superplastic behavior at high tempera-tures [3]. As a result, the total Cu content inhypoeutectic castable alloys is limited to 10wt.% Cu.

The hypereutectic Cu alloys (Cu . 32.2wt.%) have CuAl2 as a primary constituent[Fig. 12(a)], which is believed to facilitate theheterogeneous nucleation of the eutecticCuAl2, giving mechanical properties worsethan those of hypoeutectic alloys.

Al-Cu-Si TERNARY SYSTEM

The composition limits mentioned in thebinary systems can be varied in the ternaryalloys because the third element can changeboth the solidification range of tempera-tures and the microconstituents of the al-loy. The casting fluidity of the Al-Cu alloysimproves with the addition of Si because ofthe depression in the start of solidificationobserved in the ternary system Al-Cu-Si(Fig. 14). For instance, in Al-4%Cu, the ad-dition of 5% Si increases the fluidity of theheat 20% [3], but the presence of Si is detri-mental to toughness.

FIG. 13. Micrograph of alloy sample 4: 96%Al-4%Cu.

154 D. Plaza et al.

Figure 2 shows a projection of the Al-Si-CuAl2 ternary system. The ternary constitu-ents (solid solution a, Si, and CuAl2) are notshown in the diagram. Silicon does not mod-ify the solubility of Cu in Al; between 0%and 1% Si, the maximum solubility of Cu inAl remains almost fixed. At 5728C, this valuereaches 5.7 wt.% Cu. On the other hand(Fig. 15), Si forms a quasibinary systemwith CuAl2 with the formation of an eutecticat 5718C of the following composition: 4.5wt.% Si, 51 wt.% Cu, and 44.5 wt.% Al. Theeutectic consists of Si and CuAl2 without theexistence of any intermediate ternary phases.The solubilities of Cu and Si in the eutecticare almost negligible [3]. In the eutectic ofthe CuAl2-Si system, the ratio of the weightsof CuAl2 to Si can be derived by using thelever rule. This produces a value of 21.22,which, if translated into a volumetric ratio(densities of CuAl2 and Si are 4.34 and2.33g/cc, respectively), yields a value of11.39. Because of the equality between thevolume fraction and the area fraction, CuAl2is expected to appear more abundant in theeutectic than Si. This is illustrated in Fig. 16,where the darker appearing CuAl2 appearsto be greater in amount than the silicon in

the eutectic. As a result of the brittle natureof its constituents, this eutectic and all ofthem in the valley MN (Figs. 2 and 14)are unappropriate for applications de-manding good toughness properties [11].

When a liquid of 5.5 wt.% Si, 27 wt.% Cu,and 67.5 wt.% Al cools, during solidifica-tion a ternary eutectic takes place at con-stant temperature:

with relative weights of 44.66% Al, 50.80%CuAl2, and 4.54% Si. In consideration of therespective densities, the volumetric propor-tions are: fv (Al) . fv(Al2Cu) . fv (Si) withsix times more CuAl2 than Si, and ninetimes more Al than Si (Figs. 1719). This eu-tectic has low toughness because of its mi-crostructure.

In the Al-Cu-Si ternary alloys with low Cu,Si has been verified to maintain its role as aheterogeneous nucleant in the solidificationof the binary eutectics (KN in Fig. 2) and, con-sequently, the morphology of the microcon-stituents is very similar to that of the unmodi-fied Silumin. However, the Cu added to thesealloys reduces the size of the eutectic cellsof a 1 Si [4], as observed in Figs. 9 and 20.

Those alloys that end solidification withthe formation of a ternary eutectic arethought to have different heterogeneousnucleants for solidification:

Primary Al for those alloys rich in AlPrimary Si for alloys rich in Si Primary CuAl2 for alloys with more than33.2 wt.% Cu (Fig. 16)

Liquid Al CuAl2 Si+ +fi

FIG. 14. Al-Cu-Si ternary diagram.

FIG. 15. Al2Cu-Si binary diagram.

The Al-Cu-Si System 155

However, this does not seem to modify thelow toughness of the eutectic.

CONCLUSIONS

The microstructural study of the 12 alloysconducted in the present work shows that

the only possible constituents are a, CuAl2,and Si. All the Al-Cu-Si alloys that solidifyto form binary or ternary eutectics exhibitgood molding characteristics. The 67.5%Al-27%Cu-5.5%Si alloy, which solidifies at thelowest constant temperature in the system(5258C) is particularly interesting, but it has

FIG. 17. Micrograph of alloy sample 9: 69.2%Al-25.5%Cu-5.3%Si.

FIG. 16. Micrograph of alloy sample 10 (detail).

156 D. Plaza et al.

little industrial value, because of the poormechanical properties derived from its mi-crostructure.

The same is true of the ternary alloyswhen the binary eutectic compounds areabundant (a 1 Si, a 1 CuAl2, or Si 1CuAl2) because of their low toughness

characteristics. On the other hand, becauseof the high driving force for the nucleationof silicon, derived from its high meltingtemperature (14148C), this constituent doesappear in a 1 CuAl2 (Fig. 21) binary eutec-tics, resulting in even lower toughness inthe eutectic.

FIG. 19. Micrograph of alloy sample 7 (detail of the ternary eutectic).

FIG. 18. Micrograph of alloy sample 11: 61%Al-35%Cu-4%Si (detail of the ternary eutectic).

The Al-Cu-Si System 157

As a result of the microstructure ob-served at room temperature, it follows thatthe useful compositions of ternary alloys inthe Al-Cu-Si system for industrial moldingmust be limited to the shaded areas de-picted in Fig. 2.

The compositional limits shown forcastable alloys can be expanded, providedthat semisolid forming techniques are used(e.g., rheocasting) [12], in which case therounded phases improve both the castabil-ity and the mechanical properties.

FIG. 21. Micrograph of alloy sample 11: 61%Al-35%Cu-4%Si (primary silicon).

FIG. 20. Micrograph of alloy sample 8: 78%Al-10%Cu-12%Si.

158 D. Plaza et al.

References

1. J. A. Pero-Sanz: Ciencia e Ingeniera de Materiales.Dossat, Madrid, pp. 179183 (1996).

2. J. A. Pero-Sanz: Ciencia e Ingeniera de Materiales.Dossat, Madrid, pp. 333335 (1996).

3. G. Mondolfo: Aluminium Alloys. Butterworths,Boston, pp. 368372; 513515; 253266 (1976).

4. D. Plaza, J. A. Pero-Sanz, J. I. Verdeja: La estruc-tura microgrfica determinante para los lmites encomposicin de aleaciones ligeras. Rev. Minas 11:99107 (1995).

5. Woldman: Engineering Alloys. 7th ed., AmericanSociety for Metals, Metals Park, OH, p. 71 (1990).

6. ASM Metals Handbook. 7th ed., American Societyfor Metals, Metals Park, OH, pp. 140151 (1979).

7. F. Gonzalez, J. I. Verdeja, and J. A. Pero-Sanz:

Afino de grano en aleaciones Al-Si hipoeutcticas.Rev. Minas 6:1315 (1991).

8. J. A. Pero-Sanz: Materiales Metlicos. Dossat, Ma-drid, pp. 7982 (1988).

9. The Sorby Centenial Symposium on The History of Met-allurgy. Metall. Soc. of AIME, pp. 275310 (1963).

10. Thaddeus B. Massalski: Binary Alloy Phase Dia-gram. 2nd ed., American Society for Metals, MetalsPark, OH, pp. 164165 (1986).

11. W. Hufnagel: Manual del Aluminio, 2nd ed., Edito-rial Revert S. A., pp. 4143 (1992).

12. P. Kapranos, D. H. Kirkwood, and C. M. Sellars:Semi-solid processing of aluminum and high melt-ing point alloys. Proc. Inst. Mech. Eng. 207:18 (1993).

Received May 1997; accepted December 1997.