effect of strontium and phosphorus on eutectic al-si nucleation...

Effect of Strontium and Phosphorus on Eutectic Al-Si Nucleationand Formation of b-Al5FeSi in Hypoeutectic Al-Si FoundryAlloys

Y.H. CHO, H.-C. LEE, K.H. OH, and A.K. DAHLE

The present investigation was carried out on hypoeutectic Al-Si alloys containing two levels ofFe, 0.5 and 1.1 wt pct, and Sr in the range of 30 to 500 ppm. The addition of Sr in excess of100 ppm significantly reduced the number of eutectic grains and also resulted in the formationof polygonal-shaped Al2Si2Sr intermetallics. Transmission electron microscopy studies revealedthat the Al2Si2Sr phase surrounded the P-rich particles. This may suggest that the otherwisepotent nuclei for the Al-Si eutectic, aluminum phosphide (AlP), become poisoned or deactivatedby the formation of the Al2Si2Sr phase around the particles. At the high-Fe level (1.1 wt pct Fe),pre-eutectic formation of b-Al5FeSi platelets further reduced the number of eutectic Al-Sinucleation events. It is proposed that both eutectic silicon and b-Al5FeSi are preferentiallynucleated on AlP particles. Nucleation of eutectic silicon, therefore, becomes more difficultwhen it is preceded by the formation of Al2Si2Sr or b-Al5FeSi, because fewer nuclei are availableto nucleate silicon. Addition of up to 60 ppm P to the alloys increased the formation temper-ature of the b-Al5FeSi platelets but did not significantly alter the size, whereas the addition of Srdecreased the b-Al5FeSi nucleation temperature by reducing the potency of the AlP particles.

DOI: 10.1007/s11661-008-9580-8� The Minerals, Metals & Materials Society and ASM International 2008

I. INTRODUCTION

THE Al-Si alloys are the most widely used aluminumfoundry alloys today, and the control of their micro-structure is one of the most important methods toimprove the mechanical properties and the castingquality. Commercial Al-Si foundry alloys usually con-tain more than 50 vol pct of Al-Si eutectic, and extensiveresearch to control their microstructure by eutecticmodification has been carried out. The addition of alkalior alkaline earth elements changes the morphology ofeutectic silicon from flakelike to branched fibres. Themechanism of eutectic modification is still not yet fullyunderstood. For a long time, alterations in the growth ofeutectic silicon by a large increase in twin density wereused to explain eutectic modification,[1] but more recentstudies have shown that modification changes thenucleation frequency and dynamics of eutectic grainswith associated effects on the growth rate.[2,3] Inunmodified commercial Al-Si alloys, a large number ofeutectic grains nucleate at or near the primary aluminumdendrite tips, and eutectic aluminum forms epitaxiallyon the primary dendrites. On the other hand, withaddition of eutectic modifiers, i.e., Sr, a dramaticdecrease in the nucleation frequency of eutectic grains

is observed, and the grains are nucleated independentlyof the primary phase at distributed centers in theinterdendritic regions.The eutectic reaction in Al-Si alloys commences with

the nucleation of the silicon phase, which is the leadingphase during growth of the Al-Si eutectic. Crossely andMondolfo[2] reported that aluminum phosphide (AlP)particles are very potent nuclei for eutectic silicon incommercial hypoeutectic Al-Si alloys, where phospho-rus is commonly present as an impurity element. Theyproposed that the addition of sodium neutralizes AlPand thus makes nucleation of eutectic silicon moredifficult. More recent studies of eutectic nucleation haveconfirmed that AlP nucleates eutectic silicon,[3,4] and thelarge reduction in nucleation frequency of eutecticgrains in Sr-modified Al-Si alloys appears to be causedby some poisoning mechanism of the potent nuclei.[5,6]

Fe is normally also present in Al-Si alloys as animpurity element, and the presence of Fe decreases theductility of the castings by the formation of Fe-richintermetallic compounds, particularly b-Al5FeSiphase.[14] With a solidification path across the liquidussurface of the equilibrium Al-Fe-Si ternary phasediagram, b-Al5FeSi can form prior to the Al-Si eutecticreaction via a binary Al-(b-Al5FeSi) reaction when theFe concentration exceeds a critical Fe content.[7] Thesepre-eutectic b-Al5FeSi phases form large needles orplates, which are brittle and, therefore, deteriorate themechanical properties of the alloys. They also aggravatethe alloy castability by decreasing the feedabilty, causingincreased porosity formation.[10–12] There has beenmuch research into controlling the size and morphologyof b-Al5FeSi intermetallics. The addition of Mn, Cr, Be,

Y.H. CHO, Ph.D. Student, H.-C. LEE, and K.H. OH, Professors,are with the Department of Materials Science and Engineering, SeoulNational University, Seoul 151-744, Korea. Contact e-mail: [email protected] A.K. DAHLE, Professor, is with the Division of MaterialsEngineering, the University of Queensland, Brisbane, Qld 4072,Australia.

Manuscript submitted December 18, 2007.Article published online July 15, 2008

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 39A, OCTOBER 2008—2435

and Ni, which are well known as ‘‘iron neutralizers,’’can change the b-Al5FeSi phase morphology frombrittle platelike to less harmful, more compact a-(Al8Fe2Si or Al15(Fe,Mn)3,Si2) phase.

[8–12] The Sr additionsabove 0.1 pct have also been reported to change themorphology of Fe-rich intermetallics from platelike(b-phase) to starlike (a-phase).[11] Furthermore, Sr levelsof approximately 300 ppm were found to effectivelyrefine the size of b-Al5FeSi along with eutectic-siliconmodification.[12–16] It was proposed that the additionof Sr poisons the nucleation sites for the b-Al5FeSiplatelets and accelerates the dissolution process of theindividual b-Al5FeSi segments by breaking them intotwo or more fragments.[20,21]

Sigworth[9] suggested that P present in the melt(probably AlP) has a specific role in nucleating b-Al5FeSiand that the formation of large brittle Fe-rich intermet-allics can be suppressed by the addition of Sr, providedthat Sr neutralizes the effect of Samuel et al.[20–21]

reported an increase in the amount of b-Al5FeSi with theaddition of P to 319.2 alloy (Al-6.5 pct Si-3.5 pct Cu)and proposed that AlP particles could act as nucleationsites for b-Al5FeSi platelets. Eutectic silicon and iron-rich b-Al5FeSi intermetallic compounds appear to have acommon nucleation site, i.e., AlP. It is, therefore,proposed that the formation of b-Al5FeSi platelets priorto the Al-Si eutectic reaction could reduce the number ofpotent nucleation sites available to nucleate the Al-Sieutectic.

The present study aimed to investigate the effect ofstrontium on the nucleation of the Al-Si eutectic, as wellas on the formation of b-Al5FeSi intermetallics in Fecontaining Al-10 wt pct Si foundry alloys and theinteraction with phosphorus. In particular, the poisoningof the AlP nucleants was of interest. The role of AlP as aheterogeneous nucleation site for theb-Al5FeSi phasewasalso investigated by adding up to 60 ppm of phosphorus.

II. EXPERIMENTAL PROCEDURE

An Al-10 wt pct Si alloy was used as a base alloy andwas melted in an induction furnace using commercial-purity aluminum (major impurities 0.08 wt pct Fe and0.03 wt pct Si) and silicon (major impurities 0.18 wt pctFe, 0.012 wt pct Ti, and 0.04 wt pct Ca). The melt wascast into ingots with an average weight of 1 kg. Theseingots were placed in a clay-graphite crucible andremelted in an electric-resistance furnace at 760 �C. Ironwas added to the melt using ALTAB (75 pct Fe, 15 pctAl, and 10 pct nonhygroscopicNa-free flux), and themeltwas held for homogenization at 760 �C for 1 hour. Twoiron levels were studied (0.5 and 1.1 wt pct), and thechemical composition of the two base alloys is given in

Table I. The addition of strontium and phosphorus to theiron-added melts was accomplished using Al-10 pct Srmaster alloy rod and Al-19 pct Cu-1.4 pct P master alloyrod, respectively. Strontium-addition levels were in therange of approximately 30 to 490 ppm, and the phos-phorus levels were in the range of approximately 10 to60 ppm,whichwere actual values analyzed by inductivelycoupled plasma.Thermal analysis was performed in tapered, stainless-

steel cups coated with a thin layer of boron nitride, usinga centrally located, stainless steel-sheathed type-Nthermocouple (Figure 1). The thermocouple was cali-brated before and after experimentation using commer-cial-purity aluminum. For each alloy composition, twointerrupted quenching experiments were carried out, oneafter the b-phase reaction and one about midway alongthe Al-Si eutectic arrest. Two samples were takensimultaneously for each quenching test by submergingthe cups into the skimmed melt. During solidification, athermocouple was placed in only one of the samples tomonitor the cooling curve and possible reactions, andthe sample without a thermocouple was quenched into awater bath at room temperature at the designated time.

Table I. Chemical Compositions of the Base Alloys (Weight Percent)

Weight Percent Si Fe Cu Mg Mn Ti Sr P

Alloy 1 9.78 0.52 0.005 <0.005 <0.005 0.007 <0.001 <0.001Alloy 2 9.77 1.14 <0.005 <0.005 <0.005 0.008 <0.001 0.0006

Fig. 1—Schematic of the experimental setup for the thermal analysis.

2436—VOLUME 39A, OCTOBER 2008 METALLURGICAL AND MATERIALS TRANSACTIONS A

The average cooling rate prior to nucleation of the firstsolid was 1.5 K/s.

For microstructural analysis, all of the samples weresectioned vertically and prepared by standard polishingprocedures with a final polishing by a 0.05-lm colloidal-silica suspension. The microstructure of the specimenswas observed using an optical and a scanning electronmicroscope (SEM). The composition of the constituentphases was analyzed by energy-dispersive spectroscopy(EDS). For inspection of the macrostructures, thesamples were etched in a solution of 60 mL water,10 g sodium hydroxide, and 5 g of potassium ferricya-nide (modified Murakami reagent). The size and area ofb-Al5FeSi in the fully solidified alloys (uninterrupted) atfour different levels of phosphorus (0, 10, 40, and60 ppm) were measured quantitatively using a LEICAQWin (Leica Imaging System Ltd., Cambridge, Eng-land). For each specimen, 20 fields of optical micro-graphs at 200 times magnification taken fromapproximately the same location on the polished surfaceof each sample were examined.

For further examination of the interaction between Srand AlP particles, the presence of P was analyzed byelectron probe X-ray microanalysis (EPMA). A focusedion beam (FIB) was used for the preparation of thin-foilsamples (with a thickness of approximately 100 nm)containing the phase of interest, i.e., P-rich particles inthis study, for the transmission electron microscopy(TEM) observation. Focused ion beam sample prepa-ration was carried out with a Nova 200 Nanolab with a30-kV Ga liquid-metal ion source, according to thefollowing procedure: (a) identifying a region or phase ofinterest; (b) deposition of Pt on the desired region of thesample to protect the top portion of the specimen; (c)excavation using a 30-kV Ga source; (d) extracting thesample from the trench, and attaching it to a TEMsample grid with a manipulator; (e) further thinning ofthe sample by 30-kV Ga; and (f) low-energy Ga milling(at 10 kV and approximately 30 to 50 pA) for the finalmilling to remove damaged layers created by the FIB.Transmission electron microscopy observations werecarried out using 200-kV TEMs equipped with an EDSsystem.

III. RESULTS

Figures 2(a) and (b) show the cooling curves obtainedduring the solidification of both unmodified and

Sr-modified Al-10 wt pct Si alloys containing low Fe(0.5 wt pct) and high Fe (1.1 wt pct), respectively. Thenucleation temperatures for a-Al (Ta), b-Al5FeSi (Tb),and Al-Si eutectic (TN) were determined by the inter-sections of the tangent of the derivative temperaturecurves in Figure 2, and the results are shown in Table II.The addition of Sr at all levels resulted in the depressionof the eutectic-nucleation temperature, minimum tem-perature prior to recalescence, and growth temperature

Fig. 2—Cooling curves for the Al-10 pct Si alloys in the unmodifiedand Sr-modified conditions containing (a) low Fe (0.5 pct) and (b)high Fe (1.1 pct).

Table II. Characteristic Temperatures of Reactions Identified in Cooling Curves as Depicted in Figure 2

Temperature

Strontium Added Alloy 1 Strontium Added Alloy 2

1 +30 ppm Sr +290 ppm Sr +490 ppm Sr 2 +40 ppm Sr +110 ppm Sr +220 ppm Sr

Ta (�C) 592.5 587.0 590.0 592.2 590.3 590.0 592.0 591.0Tb (�C) — — — — 587.0 585.2 584.0 584.0TN (�C) 577.4 572.5 570.0 571.0 577.5 576.2 572.0 572.2TG (�C) 576.5 571.9 572.0 572.0 577.0 575.0 573.0 572.6

Notes: Ta: nucleation temperature for a-Al, Tb: nucleation temperature for pre-eutectic b phase, TN: nucleation temperature for eutectic Si, andTG: growth temperature for eutectic Si.


(TG) for both iron levels. As reported in the literature,[10]

a binary Al-(b-Al5FeSi) eutectic reaction occurs priorto the eutectic Al-Si reaction when the iron levelexceeds the critical level, which is 0.7 wt pct Fe for anAl-10 wt pct Si alloy. Pre-eutectic b-Al5FeSi wasobserved to form in the 1.1 wt pct Fe containing alloys(Tb in the curve of Figure 2(b)), whereas no pre-eutecticformation of b-Al5FeSi phase was detected in the 0.5 wtpct Fe containing alloys (Figure 2(a)). Increasing thelevel of Sr in the 1.1 wt pct Fe alloys reduced theprecipitation temperature of pre-eutectic b-Al5FeSi, aswell as the Al-Si eutectic-nucleation temperature(Table II). Figure 3 shows the microstructure of thealloys, which were quenched midway along the eutecticarrest. Figures 3(a) and (c) show the unmodified alloys,and Figures 3(b) and (d) show the Sr-modified alloys.For both levels of Fe, primary aluminum den-drites, coarse flakelike silicon, and finer needle-shapedb-Al5FeSi phases were commonly observed togetherwith quenched liquid. Long b-Al5FeSi platelets, whichare likely to form prior to the eutectic Al-Si reaction,were only observed in the higher Fe (1.1 wt pct)-containing alloys (Figures 3(c) and (d)). The additionof less than 50 ppm Sr did not alter the morphology ofthe eutectic silicon significantly; however, well-refinedand fibrous eutectic silicon was observed in the alloyswith an excess of 100 ppm Sr (Figures 3(b) and (d)).Apart from the primary dendrites, eutectic silicon andb-Al5FeSi, Sr-rich intermetallics were frequentlyobserved near the quenched liquid-dendrite interface,for both Fe levels, in alloys well modified by 290 ppm Srand 220 ppm Sr (Figures 4(a) and (b), respectively).Macrographs of all of the quenched specimens areshown in Figure 5. The unmodified alloys (Figures 5(a)and (e)) and alloys containing less than 50 ppm Sr(Figures 5(b) and (f)) do not contain any noticeablefeatures. In these alloys, the eutectic grains are too smallto be resolved in the macrographs, where the term,eutectic grains, refers to the connected Al-Si eutecticgrains, which have originated from a common source.However, the Sr-modified alloys, containing in excess of100 ppm of Sr, display circular eutectic grains in theinterior of the specimens along with a layer of eutecticgrains nucleated at the container wall. The frequency ofeutectic grain nucleation is dramatically reduced withincreased Sr levels at both Fe levels. The frequency ofeutectic grain nucleation, moreover, seems to be depen-dent not only on the Sr content but also on the Fecontent of the alloys. Increasing the Fe content of thealloys also resulted in a decrease in the number ofeutectic grains (Figures 5(c) and (d)) vs Figures 5(g) and(h)). Provided that the Sr level and quenching time arealmost identical, fewer eutectic grains appear to form inthe 1.1 wt pct Fe alloys (compare Figures 5(c) and (h)).

In the well-modified alloys, the decrease in theeutectic-nucleation frequency seems to correlate withthe formation of Sr-rich intermetallics, which occasion-ally contain centrally located second-phase parti-cles (Figure 4(a)). In the Al-10 wt pct Si-1.1 wt pctFe-220 ppm Sr alloy, a compound containing phospho-rus was found to be entrapped within the Sr-richintermetallic phase by EPMA as shown in Figure 6.

Thin-foil specimens of the cross-sectional plane of theseinternal particles were prepared by FIB milling andexamined by TEM analysis.The enlarged inset in Figure 7(a) clearly shows second-

phase particles contained within the Sr-rich intermetallicphase. These internal particles were found to be aconnected single particle in three-dimensional observa-tions using the consecutive FIB milling technique,although they (arrowed) appear to be disconnected intwo-dimensional cross section. As shown in Figure 7(b),the EDS spectrum obtained from the internal particle inthe Sr-rich phase shows a strong P peak, indicating it is aP-rich phase, possibly AlP. Along with the P peak,characteristic peaks of Al, Si, and Sr, which seem to berelated to the Sr-rich intermetallic phase, Al2Si2Sr,underneath the P-rich particles, are also observed inFigure 7(b). However, no P peak was observed in theEDS spectrum obtained from the surrounding Al2Si2Srphase (Figure 7(c)). Inside the Al2Si2Sr phase, oxideparticles were often found in association with the P-richphase near the surface polished by conventional proce-dure. It is unlikely that these oxide particles formedduring cooling of the alloy and nucleated the Al2Si2Srphase. Instead, the oxides are likely to form during theconventional polishing process, because AlP has beenreported to react actively with water.A TEM micrograph of the cross-sectioned Al2Si2Sr

phase containing the P-rich particles is shown in Fig-ure 8(a). In Figure 8(b), the diffraction pattern obtainedfrom the surrounding Sr-intermetallic phase confirmsthat the phase is Al2Si2Sr (hexagonal, P�3mL, a =0.4179 nm, c = 0.7429 nm). Figure 9 shows a TEMimage (Figure 9(a)) of Al2Si2Sr phase containing P-richparticles and its EDS spectrum (Figure 9(b)) obtainedfrom the P-rich particle. The EDS map of the particle inFigure 9(c) reveals that the distribution of P correspondswell to the internal particles present in Figure 9(a), whichis consistent with the results of the EDS analysis given inFigure 7. These internal P-containing particles are mostlikely AlP phase, and it is likely that this particlenucleated the Al2Si2Sr phase.As shown in Figure 9(b), an oxygen peak was detected

along with phosphorus in large quantities in the spec-trum obtained from the AlP particle during TEManalysis, whereas the EDS analysis during FIB millingshows no significant oxygen peak. It is, therefore,considered that oxidation of the AlP phase occurredrapidly when the thin-foil sample was taken out from theFIB chamber into a nonvacuum condition. This oxida-tion event was observed to be even further acceleratedduring TEM observation, which disrupted the crystallinenature of the AlP (Figure 9(d) presents the distributionof oxygen in the P-rich particles in the EDS map).Aluminum phosphide has been suggested to play a

significant role in nucleating the b-Al5FeSi phase, as wellas eutectic silicon.[20,21,23] It is, therefore, expected thatthe addition of P may result in easy nucleation ofb-Al5FeSi phase by providing prolific nuclei, i.e., AlPparticles. To investigate the effect of P on the forma-tion of b-Al5FeSi, three levels of P (10, 40, and 60 ppm)were added to the Al-10 wt pct Si base alloy containing1.1 wt pct Fe. Characteristic reactions identified from


Fig. 3—Optical micrographs of Al-10 pct Si-X pct Fe alloys with and without Sr addition quenched midway along the eutectic arrest: (a) 0.5 pctFe, (b) 0.5 pct Fe+290 ppm Sr, (c) 1.1 pct Fe, and (d) 1.1 pct Fe+220 ppm Sr.


the cooling-curve analysis for all three alloys are listed inTable III. The increased addition of P caused nosignificant changes in the nucleation temperature ofprimary aluminum phase (Ta), minimum temperatureprior to recalescence (Tmin), and eutectic nucleation(TN), and growth temperatures (TG). However, thenucleation temperature of the pre-eutectic b-Al5FeSi,Tb, increased from 584 �C in the alloy without Paddition to 589 �C for the 60 ppm P containing alloy,steadily increasing with each P addition.

The average area and length of the b-Al5FeSi plateletswere measured using an image analyzer and are plottedagainst P content in Figures 10(a) and (b), respectively.These results do not show any significant effect ofincreasing P levels.

IV. DISCUSSION

A. Nucleation of the Al-Si Eutectic Phase

In Al-Si alloys, Sr is well accepted as a eutectic-siliconmodifier, which effectively changes the morphology ofsilicon from platelike to fibrous at addition levels of a few

hundreds parts per million.[1] Recent studies have con-firmed that Sr addition affects not only the growth of theeutectic silicon but the nucleation behavior of theeutectic phases changes significantly, as well.[2–4,7–9]

The addition of Sr in excess of the minimum amountrequired for full modification, which is suggested to beapproximately 100 ppm for the alloys and the coolingrate adopted in this study, caused a large decrease in thenucleation frequency of eutectic grains and an associateddepression of eutectic nucleation and growth tempera-tures (Table II). These observations are consistent withthe literature.[3,8,9] In unmodified hypoeutectic Al-Sialloys, prolific nucleation events of Al-Si eutectic wereobserved to occur adjacent to the dendrite tips, whereasvery few eutectic grains nucleated in the interdendriticregions in the Sr modified alloys (Figures 3 and 5). Thereduction in nucleation frequency is at least an order ofmagnitude.Hunt[10] proposed that the columnar-to-equiaxed

transition is also applicable to the formation of theeutectic during solidification, as well as to the primarydendritic growth. In unmodified alloys, eutectic Al-Sican nucleate and grow in the form of equiaxed grainswhen potent nucleants are present and sufficient thermaland constitutional undercooling is provided. As shownin Figures 3(a) and (c) and Figure 5, each eutecticcolony, which was frequently found to form ahead ofthe primary dendrites, could be considered as anequiaxed-type grain. Crossley and Mondolfo[5] sug-gested that AlP particles, which are normally presentin commercial Al-Si alloys, act as nuclei for eutecticsilicon and that the nucleation of eutectic silicon can,therefore, occur easily with little undercooling due to theextremely good efficiency and low lattice mismatchbetween AlP and Si. They suggested that modifiers, i.e.,sodium, neutralizes the AlP and prevents the easynucleation of eutectic silicon, resulting in an increaseof undercooling. Recently, Nogita et al.[6] providedconclusive evidence of the nucleation of eutectic siliconon AlP particles by TEM analysis of specimens preparedby FIB milling.Theoretically, Sr addition is expected to be rejected by

a-Al and to cause a solute buildup at the primarydendrite-liquid interface during solidification. Solutebuildup could then increase the nucleation of equiaxedeutectic colonies due to increased constitutional und-ercooling. However, Sr addition was instead found tosignificantly decrease the number of eutectic grains(Figure 5). It is, therefore, more likely that the potencyor number of effective nucleants for eutectic silicon inthe melt was reduced. In hypoeutectic Al-Si alloys, it hasbeen suggested that the decrease in the number ofeutectic-nucleation events caused by Sr additions occursdue to poisoning of the extremely prolific nuclei-AlPnucleant particles.[2,8,9,11] However, while the exactmechanism responsible for the poisoning by Sr additionhas not been confirmed, the formation of Sr-richintermetallics, Al2Si2Sr, in Sr-modified alloys has beenproposed to be a contributing factor. Polygonal-shapedintermetallic Al2Si2Sr phase with a size of less than20 lm was frequently observed in the liquid at elevatedSr levels (Figure 4). Moreover, Al2Si2Sr was found to

Fig. 4—Optical micrographs showing Sr-rich intermetallics (arrowed)frequently observed near the dendrite-quenched liquid interface (a) in0.5 pct Fe alloy+290 ppm Sr and (b) 1.1 pct Fe alloy+220 ppm Sr.


surround the AlP particles (Figures 6 and 7), which is astrong indication of the poisoning mechanism, i.e., theaddition of Sr renders AlP nuclei ineffective by forminga layer of intermetallic phase around them.

It is worth considering the early stage of the forma-tion of Sr-rich intermetallics during solidification. Dueto the small addition levels of Sr in this study, noinformation about the Al2Si2Sr reaction could beobtained from the cooling curves (Figure 2). Based onthe assumption of Scheil conditions, solidification sim-ulations were performed using the software Thermo-Calc (TCCQ, Thermo-Calc Software Inc.)[12] with theTTAL4 (Thermo Tech Ltd., Guildford, United King-dom) database. The results in Figure 11 show thatAl2Si2Sr is predicted to form prior to the formation of

primary aluminum phase with the addition of 150 ppmSr to both Al-10 wt pct Si-0.5 wt pct Fe (Figure 11(a))and Al-10 wt pct Si-1.1 wt pct Fe alloys (Figure 11(b)).Calculation of the ternary equilibrium phase diagramfor Al-Si-Sr predicts that the Al2Si2Sr precipitationcould first take place at 596 �C when the Sr level exceeds130 ppm.P-rich particles, possibly AlP phase, were generally

present despite the alloys containing less than 10 ppmphosphorus in the present work. It has been reportedpreviously that such P levels in commercial alloys aresufficient to produce some AlP.[13] When AlP formsearly, the AlP particles may be pushed to the dendrite-liquid interface as the a-Al dendrites grow, and they canlater nucleate the eutectic silicon at low undercooling in

Fig. 5—Macrostructure of samples quenched halfway along the eutectic arrest showing a decrease in the number of eutectic grains as additionsof Sr and Fe are increased: (a) 0.5 pct Fe, (b) 0.5 pct Fe+30 ppm Sr, (c) 0.5 pct Fe+290 ppm Sr, (d) 0.5 pct Fe+490 ppm Sr, (e) 1.1 pct Fe,(f) 1.1 pct Fe+40 ppm Sr, (g) 1.1 pct Fe+110 ppm Sr, and (h) 1.1 pct Fe+220 ppm Sr.


unmodified alloys. Introduction of Sr to the alloys, onthe other hand, caused formation of Al2Si2Sr intermet-allics on the pre-existing AlP particles and, thus less, orless effective, nuclei are, therefore, available to nucleatethe eutectic when the solidification path reaches theeutectic reaction. This explanation reasonably accountsfor the large decrease in eutectic-nucleation frequency inthe Sr-modified alloys.

The quenched microstructures of the Sr-modifiedalloys in Figure 4 show the presence of isolated Sr-richintermetallics near the dendrite-liquid interface. How-ever, these Al2Si2Sr precipitates were not observed tonucleate eutectic silicon when the eutectic reactioncommenced. In Sr-modified alloys, the spherical eutecticgrains, which are several orders of magnitude larger thanthose in unmodified alloys, appear to nucleate on someother unidentified nuclei. It is not fully understood whyAl2Si2Sr is ineffective for the nucleation of eutectic grains

and what nucleates the eutectic grains in Sr-modifiedalloys. However, it is likely to be related to the increase inparticle size when the intermetallics form on the AlPparticles and possibly also to the fact that the formationof the intermetallics causes a local decrease in the siliconconcentration in the melt. It is also possible that someAlP particles remain in the melt without nucleatingAl2Si2Sr, although a larger undercooling is requiredcompared to the unmodified alloys. Further studies arerequired to explain this mechanism in detail.

B. Formation of Iron-Rich b-Al5FeSi Intermetallics

The addition of Sr to Al-Si alloys has also beenreported to cause the refinement of the iron-richintermetallics, b-Al5FeSi.

[18–22] It has been suggestedthat the effect of Sr in refining b-Al5FeSi platelets islikely to involve dissolution and fragmentation.[17–22]

Fig. 6—(a) SEM image of Sr-rich intermetallics in Al-10 pct Si-1.1 pct Fe-220 ppm Sr and corresponding EDS maps showing the distribution (b)Sr and (c) P. A P-rich region (red color) is observed within the Sr-rich intermetallics.


Fig. 7—(a) SEM image of Sr-rich intermetallics containing internal particles (arrowed in the enlarged inset) in the FIB sample and correspond-ing EDS spectra obtained from (b) the internal particle (analyzed point, A) and (c) surrounding Al2Si2Sr phase (analyzed point, B).


According to previous studies, the addition of Srsuppresses the branching of b-Al5FeSi platelets bypoisoning of preferential nucleation sites[18] and alsoby partial dissolution of b-Al5FeSi platelets resulting inthe fragmentation of platelets,[22] both causing a reduc-tion in the length of the b-Al5FeSi phase. However, thefragmentation mechanism remains uncertain. In hyp-oeutectic Al-Si alloys containing Fe in excess of thecritical level, pre-eutectic b-Al5FeSi platelets are gener-ally formed by the binary Al-(b-Al5FeSi) eutecticreaction prior to the formation of the Al-Si eutectic.[12]

The microstructure of the eutectic Al-(b-Al5FeSi) isexpected to be irregular like that of the Al-Si eutectic,where the b-Al5FeSi platelet is the leading phase and islikely to be covered by the Al dendrites during eutectic

growth. Depending on the sectioning direction of thespecimen, therefore, some segments of the b-Al5FeSiplatelets were exposed to the surface, whereas othersegmentswere still found to be covered by theAl dendrites.This may make it appear as if the latter segments arefragmented or dissolved (Figure 4(b)), but it is tempting tosuggest that these fragments are connected to a singlebranched b platelet.Sigworth et al.[23] suggested that the P (possibly AlP)

present in the melt can nucleate b-Al5FeSi and that Sraddition can suppress the precipitation of the b-Al5FeSiphase. More recent studies on the formation ofb-Al5FeSi, on the other hand, proposed that Fe-richintermetallics are likely to form on the externally wettedsurface of the oxide film, which are entrained into themelt during casting.[14,15] It has been reported that thewetted surface of oxide films provides a preferentialnucleation site for b-Al5FeSi and a large oxide filmfolded to a dry side was frequently found to formcracklike defects inside b-Al5FeSi. However, a carefulinspection of the microstructure of the specimens in thepresent study showed no direct evidence of the presenceof cracks inside the b-Al5FeSi platelets. The samples inthe present work were filled by carefully submerging thestainless-steel cups into the skimmed melt followed bycooling, and thus may not facilitate the formation andentrainment of significant oxide films in the samples.Therefore, the role of the oxide films in nucleating theb-Al5FeSi phase appears to be negligible in the presentwork.The addition of up to 60 ppm P caused an increase in

the nucleation temperature of b-Al5FeSi, Tb (Table III).This suggests that the increase in P level provides alarger number of P-based nuclei, and thus b-Al5FeSinucleates more easily at a smaller undercooling. It iswell-established that the refinement of primary silicon inhypereutectic Al-Si alloys occurs with the addition ofpotent AlP nucleants to the melt,[27] and thus prolificnucleation events of b-Al5FeSi on a larger number ofnuclei could also be expected to result in a refinement ofthe size of b-Al5FeSi. The results in Figure 10 show noclear decreasing or increasing trend in the area and sizeof b-Al5FeSi with increasing P addition. It is worthconsidering this result in more detail from the viewpointof eutectic nucleation and growth. Flood and Hunt[16]

reported that the average velocity of a eutectic interfaceis inversely proportional to the total solid/liquid inter-face area of the eutectic growth front. Unlike the growthof primary Si in the melt, pre-eutectic b-Al5FeSi formsvia a binary Al-(b-Al5FeSi) eutectic reaction. Thegrowth of b-Al5FeSi may thus occur according to thetheory proposed by Flood and Hunt. Assuming thatAlP nucleates a large number of b-Al5FeSi, the totalarea of solid/liquid interface is increased by P addition,and the interface velocity of eutectic Al-(b-Al5FeSi) is,therefore, decreased compared to the case with littlenucleation. A reduction of the growth rate is expected toresult in a coarser microstructure, i.e., coarser b-Al5FeSiintermetallics.Increased addition of Sr to theAl-10 wtpct Si-1.1 wt pct

Fe alloys gradually decreased the formation temperatureof b-Al5FeSi phase, Tb (Table II). The reason for the

Fig. 8—(a) TEM image of Al2Si2Sr phase containing P-rich particles(arrowed) and (b) corresponding electron diffraction pattern from theAl2Si2Sr phase (hexagonal, P�3mL, a = 0.4179 nm, c = 0.7429 nm)zone �12�16.


decrease of Tb brought about by Sr addition is notfully understood. However, it is possible that it is causedby the poisoning of the b-Al5FeSi nucleation sites by theformation of Al2Si2Sr on AlP particles. Early formationof Al2Si2Sr on P-rich particles may make them ineffec-

tive for the nucleation of b-Al5FeSi, just as for thenucleation of eutectic silicon. The result of the Scheilsimulation given in Figure 11(b) shows that the forma-tion of b-Al5FeSi platelets commences after the forma-tion of Al2Si2Sr and the development of the primary

Fig. 9—(a) TEM image of Al2Si2Sr phase containing P-rich particles and (b) EDS spectrum obtained from P-rich particles, with EDS maps ofinterest region (squared) showing the distribution of (c) phosphorus and (d) oxygen. The distribution of oxygen is likely to correspond with thephosphorus distribution, which appears to be due to a rapid-oxidation event of the P-rich particle.


aluminum. During solidification, primary Al2Si2Srforms on AlP particles, which are also potent nucleifor b-Al5FeSi platelets. Fewer, or less effective, nucleiare, therefore, available to nucleate b-Al5FeSi and,consequently, a larger undercooling is required.

Because AlP particles are potent nucleant substratesfor both eutectic Al-Si and b-Al5FeSi, it is worthconsidering the inter-relationship of the nucleationevents between Al-Si eutectic and b-Al5FeSi. Recentwork on the interaction between Fe and Al-Si eutectic inhypoeutectic Al-Si alloys revealed that the number ofAl-Si eutectic-nucleation events decreases as the Fe

Table III. Characteristic Temperatures of the Possible Reactions Identified in Cooling Curves for Al-10 Pct Si-1.1 Pct Fe Alloys

with Phosphorus Addition

Temperature

Phosphorus Added Alloy 2

2 2+10 ppm P 2+40 ppm P 2+60 ppm P

Ta (�C) 591.5 590.7 591.0 591.6Tb (�C) 584.0 586.8 587.1 588.9Tmin (�C) 575.9 575.2 575.4 575.8TN (�C) 576.7 576.0 576.3 576.7TG (�C) 577.1 576.4 576.6 577.1

Notes: Ta: nucleation temperature for a-Al, Tb: nucleation temperature for pre-eutectic b phase, Tmin: minimum temperature prior to recalescence,TN: nucleation temperature for eutectic Si, and TG: growth temperature for eutectic Si.

Fig. 10—The effect of phosphorus addition on (a) average area and(b) average length of the b-Al5FeSi platelets in Al-10 pct Si-1.1 pctFe alloys.

Fig. 11—Scheil simulation calculated by Thermo-Calc combinedwith TTAL4 database to predict the solidification sequence in (a) Al-10 pct Si-0.5 pct Fe-150 ppm Sr alloy and (b) Al-10 pct Si-1.1 pctFe-150 ppm Sr alloy. Labeled numbers along the curve correspondto a certain range of temperature where stable phases can be present.


content increases.[17] Dinnis et al.[19] proposed that thisinteraction is due to the fact that the eutectic Al-Si andb-Al5FeSi have common nuclei, AlP. As shown inFigures 5(c) and (h), the number of eutectic grains in theSr-modified alloys containing 1.1 wt pct Fe is still lessthan that in the Sr- modified alloys containing 0.5 wt pctFe, despite a similar concentration of Sr in bothalloys. This suggests that the formation of pre-eutecticb-Al5FeSi platelets in the higher Fe-content alloysreduces the nucleation frequency of eutectic Al-Si grainsby the formation of b-Al5FeSi on AlP, in accordancewith the observations by Dinnis et al.[19]

The nucleation sequences in unmodified andSr-modified alloys containing low Fe (0.5 wt pct) andhigh Fe (1.1 wt pct) are schematically illustrated inFigure 12. If sufficient P is present in the melt, AlPparticles form early and are pushed ahead of the dendrite-liquid interface during solidification. In unmodifiedalloys with low-Fe content (Figure 12(a)), the eutecticAl-Si reaction commences as AlP nucleates the polygo-nal-shaped eutectic silicon near the dendrite tips wherethe silicon level is locally high enough. In Sr-containingalloys, on the other hand, Al2Si2Sr is likely to form onpre-existing AlP prior to the eutectic Al-Si reaction, soAlP particles would not play a significant role in thenucleation of the eutectic grains and far fewer eutecticgrains form in the interdendritic liquid (Figure 12(c)).When the Fe content is high enough to form pre-eutecticb-Al5FeSi platelets, AlP can nucleate both b-Al5FeSiplatelets and eutectic silicon. Because the pre-eutecticb-Al5FeSi platelets are formed on AlP particles prior tothe eutectic Al-Si reaction, fewer nuclei are availablefor nucleating the Al-Si eutectic, resulting in fewer Al-Sigrains (Figure 12(b)). The addition of Sr, along withhigh Fe content, reduces the nucleation of b-Al5FeSiplatelets, as well as eutectic Al-Si grains, by the formationof Al2Si2Sr phase on AlP (Figure 12(d)). Furthermore,even though the addition of Sr is not sufficient for theAl2Si2Sr phase to consume all of the AlP particles in theliquid, it is expected that the precipitation of b-Al5FeSiplatelets onto the remaining AlP results in even lessnucleation of eutectic Al-Si grains (Figure 12(d)). Thisconclusion is strongly supported by the observation thatthe decrease in the number of eutectic grains in thealloy with lower Sr and higher Fe content, the 110 ppmSr in Al-10 wt pct Si-1.1 wt pct Fe alloy, is moresignificant than that in the alloy with higher Sr and lowerFe content, the 290 ppm Sr in Al-10 wt pct Si-0.5 wt pctFe alloy.

V. CONCLUSIONS

The effect of Sr addition on eutectic Al-Si nucleationand the formation of b-Al5FeSi in hypoeutectic Al-Sifoundry alloys were investigated.

Sr additions exceeding 100 ppm dramatically reducedthe number of eutectic Al-Si grains and decreased theeutectic-nucleation temperature. Poisoning of the potentAlP nuclei for the nucleation of eutectic grains isproposed as the mechanism. The result shows that Srforms Al2Si2Sr intermetallic phase onto the AlP

Fig. 12—Schematic illustration of solidification sequence with theformation of Al2Si2Sr, primary Al dendrite, pre-eutectic b-Al5FeSi,and eutectic Al-Si: (a) AlP present at the dendrite liquid interfacenucleates eutectic silicon in unmodified alloy with low Fe content;(b) both eutectic silicon and b-Al5FeSi nucleate on AlP when the Feaddition to unmodified alloys is high enough to form pre-eutecticb-Al5FeSi platelets; (c) Sr-rich Al2Si2Sr intermetallics present inSr-modified alloys with low Fe content form on AlP and far fewereutectic grains (spherical), which do not nucleate on AlP grow in theinterdendritic liquid; and (d) both Al2Si2Sr and pre-eutectic b-Al5FeSi nucleate on AlP, and thus even fewer eutectic grains will formin Sr-modified alloys with high Fe content.


particles. The potency of AlP as nuclei for the Al-Sieutectic is thereby reduced, and, therefore, the nucle-ation of eutectic grains is reduced.

Aluminum phosphide is proposed to be a commonnucleation site for both eutectic Al-Si grains andb-Al5FeSi phase. Sr addition gradually decreased theb-Al5FeSi nucleation temperature. The formationof b-Al5FeSi appears to be suppressed by the presenceof Al2Si2Sr, which is believed to also be caused bydeactivation of the AlP nuclei.

The Al-Si eutectic-nucleation frequency was de-creased by increasing the Fe concentration in excess ofthe critical level for pre-eutectic b-Al5FeSi nucleation inthe Sr- modified alloys. The Scheil simulations showedthat the formation of pre-eutectic b-Al5FeSi occurs afterthe formation of Al2Si2Sr and primary Al dendrites. It issuggested that AlP particles present in the melt areconsumed by the formation of Al2Si2Sr, as well as thenucleation of b-Al5FeSi in the early stages of solidifica-tion. It is, therefore, concluded that the decrease innucleation frequency of eutectic Al-Si grains is causedby a lack of efficient nuclei when the solidificationprocess reaches the eutectic Al-Si reaction.

ACKNOWLEDGEMENTS

The authors thank Ms. H.K. Kang for her help inEPMA analysis and Mr. D.H. Kim for the prepara-tion of TEM using FIB. The authors also thank Drs.S. McDonald and K. Nogita, University of Queens-land, for their help in lab experiments and many valu-able discussions.

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