CommunicationNucleants of Eutectic Silicon in Al-SiHypoeutectic Alloys: b-(Al, Fe, Si) orAlP Phase

SONG-MAO LIANG andRAINER SCHMID-FETZER

A thermodynamic description of the Al-Si-P-Fe qua-ternary system focused on Al-(Si)-rich alloys is devel-oped. The solidification sequence in typical Al-7Si castalloys is derived using thermodynamic calculations ofthe phase diagrams and solidification simulation underScheil and constrained Scheil conditions. The previouslyclaimed nucleation of eutectic silicon by b-(Al,Fe,Si)particles is not possible because under all conditions,b-(Al,Fe,Si) precipitates after (Si) in pertinent alloys.Variation of P in the ppm range is crucial because itchanges the solidification sequence of AlP and (Si).

DOI: 10.1007/s11661-014-2522-8� The Minerals, Metals & Materials Society and ASMInternational 2014

The impact of impurities, such as Fe or P, on thenucleation behavior of eutectic silicon in hypoeutecticAl-Si cast alloys is still a hot topic in scientific researchand also of industrial importance.[1–6] It is generallyaccepted that AlP particles are effective nuclei foreutectic (Si), a summary on the extensive experimentalwork can be found in our recent work on the thermo-dynamics of the Al-Si-P system.[7] The critical P contentfor Al-Si alloys to precipitate AlP before (Si) has beendetermined through meticulous thermodynamic calcu-lations.[7] On the other hand, Shankar et al.[8,9] proposedthe assumption that b-(Al,Si,Fe) initiates the nucleationevent of eutectic (Si), and insisted on that assump-tion[10,11] after the argument with Dahle and Hill-ert.[12,13] In the present work, we shall reexamine theassumption on b-(Al,Fe,Si) and clarify the dispute bydedicated thermodynamic calculations.

For that purpose, the thermodynamic description ofthe Al-Si-P-Fe quaternary system focused on Al-(Si)-rich alloys has been developed in this work by joiningthe well-established Al-Si-P description[7] with theAl-Fe-Si description by Eleno et al.[14] The latter ispreferred compared to earlier assessments of Al-Fe-Si[15–17] because it[14] takes all the more recent experi-mental work[18–23] into account. The results calculated

using the dataset of Reference 14 agree better with theexperimentally determined invariant reaction tempera-tures[23,24] in Al-rich alloys compared to References 15through 17. The subsystems P-Fe, Al-P-Fe, and Si-P-Feare irrelevant for applications to the Al-rich Al-Si-basedalloys because the highly stable compound AlP forms tielines to the Al-rich Al-Fe phases.[25] Therefore, none ofthe Fe-P compounds, such as Fe3P or Fe2P, can form inthese Al-rich alloys during solidification or in equilib-rium. No quaternary Al-Si-P-Fe phase is reported.It is emphasized that this description is very well

validated by the entity of experimental data of theconstituting subsystems and the consistency of thethermodynamic functions in the envisaged compositionrange of Al-Si alloys with upto a few mass percentage Feand P of the Al-Si-P-Fe system. All thermodynamiccalculations in the present work were performedusing the integrated software package Pandat (www.computherm.com).[26]

Figure 1 shows the calculated isopleth phase diagramsection of Al-Si-Fe alloys at constant 7 pct Si andincreasing Fe content. Mass pct and ppm = lg/g isused throughout this paper. Two key points (M and N)are highlighted in Figure 1(b), calculated at 0.0036 pctFe(point M), and at 0.3028 pct Fe(point N). Shankaret al.[8–10] did the same calculation with Pandat using anearly database version, PanAl2b. Their calculated phasediagrams essentially agree with our calculations, withslightly different values of point M (at 0.0038 pct Fe)and point N (at 0.42 pct Fe). According to thesediagrams, the Al-7Si-xFe alloys can be divided intothree types of alloys by the formation of b-(Al,Fe,Si)phase. Type I, with Fe content less than point M(x<M), no b-(Al,Fe,Si) phase can be formed fromequilibrium solidification; type II, Fe content locatedbetween points M and N (M £x £N), b-(Al,Fe,Si) phaseprecipitates at last stage of solidification after theeutectic (monovariant) reaction of Lfi (Al)+ (Si); typeIII, Fe content higher than at point N (x>N), someb-(Al,Fe,Si) precipitate before (Si), and subsequentlysolidification ends at Lfi (Al)+ (Si)+ b-(Al,Fe,Si) ter-nary eutectic (invariant) reaction.Shankar et al.[8–10] correctly stated that b-(Al,Fe,Si)

may occur in alloys of type II (and III), with Fe contenthigher than point M. However, they provided a wronginterpretation of the solidification sequence ‘‘…forma-tion of the primary aluminum dendrites, followed byprecipitation of the b-(Al,Fe,Si) phase, and then theeutectic phases’’ for alloys of type II. The phasediagrams in Figure 1 clearly show that (Si) crystallizesin the three-phase region L+ (Al)+ (Si) well before theoccurrence of b-(Al,Fe,Si). For the example of alloyAl-7 pct Si-0.06 pct Fe, the majority amount of 82 pctof the total (Si) crystallizes first in the monovariantLfi (Al)+ (Si) eutectic from 849.9 K to 848.8 K(576.8 �C to 575.7 �C), and only 18 pct of (Si) in thefinal Lfi (Al)+ (Si)+ b eutectic at 848.8 K (575.7 �C).These phase amounts are calculated, they cannot beread from Figure 1 because the tie lines are not in theplane of that isopleth.In order for the solidification sequence claimed in

References 8 through 10 to occur, the alloy must be of

SONG-MAO LIANG, Postdoc Research Fellow, and RAINERSCHMID-FETZER, Professor, are with the Institute of Metallurgy,Clausthal University of Technology, Robert-Koch-Str. 42, 38678Clausthal-Zellerfeld, Germany. Contact e-mail: schmid-fetzer@tu-clausthal.de

Manuscript submitted June 23, 2014.Article published online September 5, 2014

5308—VOLUME 45A, NOVEMBER 2014 METALLURGICAL AND MATERIALS TRANSACTIONS A

type III (x>N). However, all the alloys investigated inReferences 8 through 10 are of type I and II, well belowpoint N, which is calculated at 0.30 pct Fe (this work) or0.42 pct Fe (from Shankar et al.[8–10]). Even the alloywith highest Fe content, alloy A in Reference 8, containsonly 0.0637 pct Fe. Therefore, the claimed solidificationsequence Lfi (Al)fi (Al)+ b-(Al,Si,Fe)fi (Al)+b-(Al,Si,Fe)+ (Si) is a misinterpretation of the phasediagram in References 8 through 10

The phase formation in the as-cast state of alloys isusually better reflected by the Scheil-Gulliver solidifica-tion simulation,[27] assuming blocked diffusion in thesolidified phases, in contrast to infinitely fast diffusion,as discussed above for the equilibrium case. We haveperformed such Scheil simulations using the new ther-modynamic description of the Al-Si-P-Fe system. Shan-kar et al.[8,9] investigated two groups of alloys toillustrate their claim, alloy A with a specific chemicalcomposition of Al-7.45 pct Si-0.0637 pct Fe-8.57 ppmP, and alloy B with a composition range Al-(4.5~12.5) pct Si-(0.003309~0.004899) pct Fe-(3.42~5.88)ppm P. Both alloys contain ppm levels of Cu, Mn, Mgimpurities, which are not considered for the presentcalculations due to their little effect on the solidificationsequence of the alloys. Figure 2(a) shows the standardScheil simulation of the Al-7.45 pct Si-0.0637 pctFe-8 ppm P alloy (alloy A in Reference 8). The phaseprecipitation sequence is Lfi (Al)fi (Al)+AlPfi(Al)+AlP+ (Si) fi (Al)+AlP+ (Si)+ b-(Al,Si,Fe).The AlP phase precipitates before (Si), while theb-(Al,Si,Fe) is the last formed phase at the end ofsolidification at 848.8 K (575.7 �C) after most of theeutectic (Si) was formed in the third solidification stepjointly with (Al) and some AlP. Therefore, even underScheil conditions, it is impossible that b-(Al,Si,Fe)nucleates the eutectic (Si).

Finally, a more advanced calculation is performed,the constrained Scheil simulation.[7] Let us assumeb-(Al,Si,Fe) is the nucleant for (Si) in a somewhatundercooled melt as proposed, with DT upto a few K, byShankar et al.[8,9] The question is, how much underco-oling DT is required, defined as the temperature differ-ence between the regular eutectic at 848.8 K (575.7 �C)and the first occurrence of the ‘‘nucleant b’’ under theassumption that (Si) has not yet formed in the under-cooled melt. That is done by the constrained Scheilsolidification simulation, by suspending the (Si) phase asshown in Figure 2(b); the b phase precipitates at793.1 K (520 �C) and a huge undercooling ofDT = 55.7 K is required before the b phase could startnucleating eutectic (Si). That is in stark disagreementwith all the observed much smaller undercooling data.This finding is corroborated by a closer look at the

liquid composition evolution at the solidification frontfor alloy A in Figure 3. The liquid compositions understandard Scheil simulation (solid lines) and constrainedScheil simulation (dashed lines) are identical down to849.9 K (576.8 �C), at which (Si) starts precipitating inthe standard Scheil simulation with low Fe content,0.1143 pct Fe. Subsequently, the Fe content rapidlygrows to 0.61 pct Fe at 848.8 K (575.7 �C), whereb-(Al,Si,Fe) forms in the terminating eutectic reaction.For the constrained Scheil simulation with (Si) sus-pended, however, the Fe enriches only gradually in theliquid, reaching 0.1895 pct Fe at 793.1 K (520 �C) whenb-(Al,Si,Fe) precipitates. These results emphasize thatthe formation of b-(Al,Si,Fe) in the realistic standardScheil simulation is due to the prior solidification of (Si),generating strong Fe-enrichment in the rapidly shrink-ing liquid fraction until the saturation with b-(Al,Si,Fe)is reached in the final eutectic. Without the priornucleation of (Si) from other nuclei, such as AlP, the

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Fig. 1—Phase diagram section Al-7Si-xFe (a): x = 0~0.01 mass pct, (b): x = 0~0.5 mass pct. The total Fe content in the alloys (not in the pha-ses) at the two key points is 0.0036 pct Fe (point M) and 0.3028 pct Fe (point N).

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 45A, NOVEMBER 2014—5309

b-(Al,Si,Fe) precipitation is possible only with giantundercooling. Therefore, b-(Al,Si,Fe) cannot act asnuclei for (Si) in alloy A.

Finally, our recent thermodynamic analysis revealingthe impact of P content on the eutectic solidificationsequence[7] will be elaborated for the alloys A and B.[8,9]

That will provide a simple explanation for the strikinglydifferent microstructures of ‘‘Al-7Si-0.07Fe’’(A) and‘‘Al-7Si-0.003Fe’’(B) alloys (see Figure 6 in Reference10 and Figure 3 in Reference 11), which had beenclaimed by Shankar et al.[11] as the ultimate proof ‘‘thatb-(Al,Si,Fe) initiates the nucleation event of eutectic Si’’,thus ending the discussions with Dahle and Hillert.[12,13]

As detailed above, only a composition range wasprovided for alloy B in Reference 8, which containsmuch less Fe than alloy A. For alloy B, the compositionAl-8.5 pct Si-0.0032 pct Fe-4 ppm P was selected for thethermodynamic calculations, because Al-8.5 pct Si-0.0032 pct Fe is the alloy composition mentioned inFigure 4 of Reference 8 and Figure 2 of Reference 9,and 4 ppm P is in the middle of their reported range of‘‘3.42 to 5.88 ppm P’’.[8] Such reported accuracy, muchbetter than 1 ppm P, may be hardly achieved in practicalchemical analysis. The evolution of phase fractions ofthis alloy B during Scheil solidification simulation isshown in Figure 4. The phase precipitation sequenceis Lfi (Al)fi (Al)+ (Si) fi (Al)+ (Si)+AlPfi (Al)+(Si)+AlP+ b-(Al,Si,Fe). The sequence of (Si) and AlPcannot be discerned from this diagram because of theminute temperature difference of 0.003 K between theAl-Si binary eutectic and the Al-Si-AlP ternary eutectic.The sequence is, however, firmly established in thecalculation and the composition variation of the liquidphase, not shown here.

The key distinction between alloys A and B is only theprecipitation sequence of AlP and (Si). In alloy A, the

AlP phase forms before (Si); in alloy B, the AlP phaseforms after (Si). The b-(Al,Si,Fe) phase precipitatesalways at the very last stage of solidification, after the(Si) phase, in both alloys. The AlP particles are effectivenuclei for eutectic (Si), and extensive experimentalresults were compiled and compared with thermody-namic calculations in the published precipitation map.[7]

It was found that coarse eutectic (Si) is always observedif the P content is above a threshold value, resulting inAlP formation before (Si). Opposed to that, fine eutectic(Si) is always observed if the P content is below the

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Solidification sequence:L (Al)L (Al) + AlPL (Al) + AlP + (Si)L (Al) + AlP + (Si) + 575.7 °C solidus.

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Fig. 2—Solidification simulation of Al-7.45 pct Si-0.0637 pct Fe-8 ppm P alloy (alloy A in Ref. [8]): (a) standard Scheil simulation, all phases areallowed to form and (b) constrained Scheil simulation with suspended (Si) phase. In figure (a), the liquid fraction [zero at 848.8 K (575.7 �C)] isplotted with an arrow to avoid overlap with other curves. The solid lines in (b), down to 849.9 K (576.8 �C), are the same as in (a).

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Fig. 3—The evolution of liquid composition during the Scheil solidi-fication simulation of Al-7.45 pct Si-0.0637 pct Fe-8 ppm P alloy (al-loy A in Ref. [8]); the standard simulation (all phases are allowed toform, solid lines) is compared to the constrained simulation, sus-pending the (Si) phase (dashed lines).

5310—VOLUME 45A, NOVEMBER 2014 METALLURGICAL AND MATERIALS TRANSACTIONS A

threshold value, resulting in AlP formation after (Si),because then AlP cannot act as nucleant. The dramaticimpact of the P content and the increase of the Pthreshold value with Si composition were developedquantitatively in Reference 7. This approach is extendedin this work to the Al-Si-P-Fe quaternary system.

This decisive impact of P content has obviously beenoverlooked by Shankar et al.,[8–11] because in theirdiscussion with Dahle and Hillert,[12,13] they claimed aconstant P content (�8 ppm P) for both alloys A and B,e.g., in Figure 3 of Reference 11. That is in conflict withthe references displaying their original experimentalwork[8] with ‘‘8.57 ppm P’’ in alloy A and ‘‘3.42 to5.88 ppm P’’ in alloy B. Both the Fe and P content inalloys A and B were not tailored by the operator butcontrolled by the impurity of the raw ingots. Accordingto the thermodynamic simulation results presentedabove, it can be easily understood that the coarse (Si)is produced due to the nucleation of AlP precipitatebefore (Si) in alloy A rather than due to the b-(Al,Si,Fe)phase, which always precipitates after (Si).

Shankar et al.[8,9] cited various references to supporttheir claim that b-(Al,Si,Fe) initiates the nucleationevent of eutectic Si. In brief these concern the claim ofan enrichment of Fe at the solidification solid/liquidfront,[28,29] and other work suggesting ‘‘that b-(Al,Si,Fe)could be a nucleant for eutectic Si’’,[30,31] however, wefailed to detect any relevant supporting information inthe original papers.[28–31] It is noted that all theexperimental results on Al-Fe-Si collected in Reference28 have been thermodynamically assessed by Elenoet al.,[14] which was used in the present work.

The possibility that b-(Al,Si,Fe) phase might be thenucleant of eutectic (Si) was first proposed by Tayloret al.[32] because of a feasible orientation relationshipbetween b and (Si). The main contribution of Tayloret al.[32–34] is the effect of Fe content on the formation ofporosity in Al-Si-Fe alloys. The solidification sequence

of (Si) and b-(Al,Si,Fe) phase has an obvious impact onthe casting defects but little effect on the morphology ofeutectic (Si).[34] Interestingly, Shankar et al.[8–11] did notcite any work of Taylor et al.[32–34]

Later, Lu et al.[35] carried out a series of quenchexperiments to investigate the possibility of the b-(Al,Si,Fe) phase as nucleation sites of eutectic (Si). Theirresults showed that the b-(Al,Si,Fe) phase precipitatedbefore eutectic (Si) in Al-7Si-0.4Mg-0.7Fe alloy andafter eutectic (Si) in Al-7Si-0.4Mg-0.3Fe alloy, which isalso in agreement with our thermodynamic calculations.Figure 5 represents standard Scheil simulation results ofAl-7Si-0.7Fe-10 ppm P alloy, which is the simplifiedconsideration of Al-7Si-0.4Mg-0.7Fe alloy by ignoringthe effect of Mg and assuming 10 ppm contained in thealloy. The alloy composition is also representative for anAl recycling alloy, usually containing more than 0.3 pctFe. The solidification sequence of this alloy isL fi (Al)fi (Al)+AlPfi (Al)+AlP+ b-(Al,Si,Fe)fi(Al)+AlP+ b-(Al,Si,Fe)+ (Si). Clearly, b-(Al,Si,Fe)phase precipitates before (Si) phase but after AlP phase.However, the microstructure observation showed thatthe eutectic (Si) was nucleated independently of theb-(Al,Si,Fe) phase even in the Al-7Si-0.4Mg-0.7Fealloys.[35] Lu et al.[35] also discussed that some crystal-lographic matching between Si and b-(Al,Si,Fe) phasealone is not enough to demonstrate its potency as anucleant for eutectic (Si). Shankar et al.[8] have nicelyshown that b-(Al,Si,Fe) is found next to (Si) withpreferred crystallographic orientation. That is nowunderstood by its formation in the very last invarianteutectic, Lfi (Al)+ (Si)+ b-(Al,Si,Fe) with AlP form-ing before or after (Si), and b-(Al,Si,Fe) always formslast.In summary, the claim that b-(Al,Si,Fe) initiates the

nucleation event of eutectic (Si) proposed by Shankaret al.[8–11] has been disproved by dedicated thermody-namic calculations. The b-(Al,Si,Fe) phase always pre-

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β

Fig. 4—Scheil solidification simulation of Al-8.5 pct Si-0.0032 pctFe-4 ppm P alloy (alloy B in Ref. [8]); standard simulation, all pha-ses are allowed to form.

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Fig. 5—Standard Scheil solidification simulation of Al-7 pctSi-0.7 pct Fe-10 ppm P alloy, the b-(Al,Si,Fe) phase precipitates be-fore (Si) phase but after AlP phase.

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cipitates after the (Si) phase under all conditions for Fecontent below about 0.3 pctFe. The strikingly differentmicrostructures of ‘‘Al-7Si-0.07Fe’’(A) and ‘‘Al-7Si-0.003Fe’’(B) alloys[11] are clearly understood by thecritical impact of P content on the eutectic solidificationsequence.

This study is supported by the German ResearchFoundation (DFG) in the Bundled Proposal‘‘PAK499’’ under grant no. Schm 588/39.

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