Effect of Strontium and Phosphorus on Eutectic Al-Si Nucleation and Formation of b-al 5 FeSi in Hypoeutectic Al-Si Foundry Alloys

Transcription

1 Effect of Strontium and Phosphorus on Eutectic Al-Si Nucleation and Formation of b-al 5 FeSi in Hypoeutectic Al-Si Foundry Alloys 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 of Fe, 0.5 and 1.1 wt pct, and Sr in the range of 30 to 500 ppm. The addition of Sr in excess of 100 ppm significantly reduced the number of eutectic grains and also resulted in the formation of polygonal-shaped Al 2 Si 2 Sr intermetallics. Transmission electron microscopy studies revealed that the Al 2 Si 2 Sr phase surrounded the P-rich particles. This may suggest that the otherwise potent nuclei for the Al-Si eutectic, aluminum phosphide (AlP), become poisoned or deactivated by the formation of the Al 2 Si 2 Sr phase around the particles. At the high-fe level (1.1 wt pct Fe), pre-eutectic formation of b-al 5 FeSi platelets further reduced the number of eutectic Al-Si nucleation events. It is proposed that both eutectic silicon and b-al 5 FeSi are preferentially nucleated on AlP particles. Nucleation of eutectic silicon, therefore, becomes more difficult when it is preceded by the formation of Al 2 Si 2 Sr or b-al 5 FeSi, because fewer nuclei are available to nucleate silicon. Addition of up to 60 ppm P to the alloys increased the formation temperature of the b-al 5 FeSi platelets but did not significantly alter the size, whereas the addition of Sr decreased the b-al 5 FeSi nucleation temperature by reducing the potency of the AlP particles. DOI: /s Ó The Minerals, Metals & Materials Society and ASM International 2008 I. INTRODUCTION THE Al-Si alloys are the most widely used aluminum foundry alloys today, and the control of their microstructure is one of the most important methods to improve the mechanical properties and the casting quality. Commercial Al-Si foundry alloys usually contain more than 50 vol pct of Al-Si eutectic, and extensive research to control their microstructure by eutectic modification has been carried out. The addition of alkali or alkaline earth elements changes the morphology of eutectic silicon from flakelike to branched fibres. The mechanism of eutectic modification is still not yet fully understood. For a long time, alterations in the growth of eutectic silicon by a large increase in twin density were used to explain eutectic modification, [1] but more recent studies have shown that modification changes the nucleation frequency and dynamics of eutectic grains with associated effects on the growth rate. [2,3] In unmodified commercial Al-Si alloys, a large number of eutectic grains nucleate at or near the primary aluminum dendrite tips, and eutectic aluminum forms epitaxially on the primary dendrites. On the other hand, with addition of eutectic modifiers, i.e., Sr, a dramatic decrease in the nucleation frequency of eutectic grains Y.H. CHO, Ph.D. Student, H.-C. LEE, and K.H. OH, Professors, are with the Department of Materials Science and Engineering, Seoul National University, Seoul , Korea. Contact huchul@ snu.ac.kr A.K. DAHLE, Professor, is with the Division of Materials Engineering, the University of Queensland, Brisbane, Qld 4072, Australia. Manuscript submitted December 18, Article published online July 15, 2008 is observed, and the grains are nucleated independently of the primary phase at distributed centers in the interdendritic regions. The eutectic reaction in Al-Si alloys commences with the nucleation of the silicon phase, which is the leading phase during growth of the Al-Si eutectic. Crossely and Mondolfo [2] reported that aluminum phosphide (AlP) particles are very potent nuclei for eutectic silicon in commercial hypoeutectic Al-Si alloys, where phosphorus is commonly present as an impurity element. They proposed that the addition of sodium neutralizes AlP and thus makes nucleation of eutectic silicon more difficult. More recent studies of eutectic nucleation have confirmed that AlP nucleates eutectic silicon, [3,4] and the large reduction in nucleation frequency of eutectic grains in Sr-modified Al-Si alloys appears to be caused by some poisoning mechanism of the potent nuclei. [5,6] Fe is normally also present in Al-Si alloys as an impurity element, and the presence of Fe decreases the ductility of the castings by the formation of Fe-rich intermetallic compounds, particularly b-al 5 FeSi phase. [14] With a solidification path across the liquidus surface of the equilibrium Al-Fe-Si ternary phase diagram, b-al 5 FeSi can form prior to the Al-Si eutectic reaction via a binary Al-(b-Al 5 FeSi) reaction when the Fe concentration exceeds a critical Fe content. [7] These pre-eutectic b-al 5 FeSi phases form large needles or plates, which are brittle and, therefore, deteriorate the mechanical properties of the alloys. They also aggravate the alloy castability by decreasing the feedabilty, causing increased porosity formation. [10 12] There has been much research into controlling the size and morphology of b-al 5 FeSi intermetallics. The addition of Mn, Cr, Be, METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 39A, OCTOBER

2 and Ni, which are well known as iron neutralizers, can change the b-al 5 FeSi phase morphology from brittle platelike to less harmful, more compact a-(al 8 Fe 2 Si or Al 15 (Fe,Mn) 3,Si 2 ) phase. [8 12] The Sr additions above 0.1 pct have also been reported to change the morphology of Fe-rich intermetallics from platelike (b-phase) to starlike (a-phase). [11] Furthermore, Sr levels of approximately 300 ppm were found to effectively refine the size of b-al 5 FeSi along with eutectic-silicon modification. [12 16] It was proposed that the addition of Sr poisons the nucleation sites for the b-al 5 FeSi platelets and accelerates the dissolution process of the individual b-al 5 FeSi segments by breaking them into two or more fragments. [20,21] Sigworth [9] suggested that P present in the melt (probably AlP) has a specific role in nucleating b-al 5 FeSi and that the formation of large brittle Fe-rich intermetallics can be suppressed by the addition of Sr, provided that Sr neutralizes the effect of Samuel et al. [20 21] reported an increase in the amount of b-al 5 FeSi with the addition of P to alloy (Al-6.5 pct Si-3.5 pct Cu) and proposed that AlP particles could act as nucleation sites for b-al 5 FeSi platelets. Eutectic silicon and ironrich b-al 5 FeSi intermetallic compounds appear to have a common nucleation site, i.e., AlP. It is, therefore, proposed that the formation of b-al 5 FeSi platelets prior to the Al-Si eutectic reaction could reduce the number of potent nucleation sites available to nucleate the Al-Si eutectic. The present study aimed to investigate the effect of strontium on the nucleation of the Al-Si eutectic, as well as on the formation of b-al 5 FeSi intermetallics in Fe containing Al-10 wt pct Si foundry alloys and the interaction with phosphorus. In particular, the poisoning of the AlP nucleants was of interest. The role of AlP as a heterogeneous nucleation site for the b-al 5 FeSi phase was also investigated by adding up to 60 ppm of phosphorus. Table I. The addition of strontium and phosphorus to the iron-added melts was accomplished using Al-10 pct Sr master alloy rod and Al-19 pct Cu-1.4 pct P master alloy rod, respectively. Strontium-addition levels were in the range of approximately 30 to 490 ppm, and the phosphorus levels were in the range of approximately 10 to 60 ppm, which were actual values analyzed by inductively coupled plasma. Thermal analysis was performed in tapered, stainlesssteel cups coated with a thin layer of boron nitride, using a centrally located, stainless steel-sheathed type-n thermocouple (Figure 1). The thermocouple was calibrated before and after experimentation using commercial-purity aluminum. For each alloy composition, two interrupted quenching experiments were carried out, one after the b-phase reaction and one about midway along the Al-Si eutectic arrest. Two samples were taken simultaneously for each quenching test by submerging the cups into the skimmed melt. During solidification, a thermocouple was placed in only one of the samples to monitor the cooling curve and possible reactions, and the sample without a thermocouple was quenched into a water bath at room temperature at the designated time. II. EXPERIMENTAL PROCEDURE An Al-10 wt pct Si alloy was used as a base alloy and was melted in an induction furnace using commercialpurity aluminum (major impurities 0.08 wt pct Fe and 0.03 wt pct Si) and silicon (major impurities 0.18 wt pct Fe, wt pct Ti, and 0.04 wt pct Ca). The melt was cast into ingots with an average weight of 1 kg. These ingots were placed in a clay-graphite crucible and remelted in an electric-resistance furnace at 760 C. Iron was added to the melt using ALTAB (75 pct Fe, 15 pct Al, and 10 pct nonhygroscopic Na-free flux), and the melt was held for homogenization at 760 C for 1 hour. Two iron levels were studied (0.5 and 1.1 wt pct), and the chemical composition of the two base alloys is given in Fig. 1 Schematic of the experimental setup for the thermal analysis. Table I. Chemical Compositions of the Base Alloys (Weight Percent) Weight Percent Si Fe Cu Mg Mn Ti Sr P Alloy <0.005 < <0.001 <0.001 Alloy <0.005 <0.005 < < VOLUME 39A, OCTOBER 2008 METALLURGICAL AND MATERIALS TRANSACTIONS A

3 The average cooling rate prior to nucleation of the first solid was 1.5 K/s. For microstructural analysis, all of the samples were sectioned vertically and prepared by standard polishing procedures with a final polishing by a 0.05-lm colloidalsilica suspension. The microstructure of the specimens was observed using an optical and a scanning electron microscope (SEM). The composition of the constituent phases was analyzed by energy-dispersive spectroscopy (EDS). For inspection of the macrostructures, the samples were etched in a solution of 60 ml water, 10 g sodium hydroxide, and 5 g of potassium ferricyanide (modified Murakami reagent). The size and area of b-al 5 FeSi in the fully solidified alloys (uninterrupted) at four different levels of phosphorus (0, 10, 40, and 60 ppm) were measured quantitatively using a LEICA QWin (Leica Imaging System Ltd., Cambridge, England). For each specimen, 20 fields of optical micrographs at 200 times magnification taken from approximately the same location on the polished surface of each sample were examined. For further examination of the interaction between Sr and AlP particles, the presence of P was analyzed by electron probe X-ray microanalysis (EPMA). A focused ion beam (FIB) was used for the preparation of thin-foil samples (with a thickness of approximately 100 nm) containing the phase of interest, i.e., P-rich particles in this study, for the transmission electron microscopy (TEM) observation. Focused ion beam sample preparation was carried out with a Nova 200 Nanolab with a 30-kV Ga liquid-metal ion source, according to the following procedure: (a) identifying a region or phase of interest; (b) deposition of Pt on the desired region of the sample to protect the top portion of the specimen; (c) excavation using a 30-kV Ga source; (d) extracting the sample from the trench, and attaching it to a TEM sample grid with a manipulator; (e) further thinning of the sample by 30-kV Ga; and (f) low-energy Ga milling (at 10 kv and approximately 30 to 50 pa) for the final milling to remove damaged layers created by the FIB. Transmission electron microscopy observations were carried out using 200-kV TEMs equipped with an EDS system. Sr-modified Al-10 wt pct Si alloys containing low Fe (0.5 wt pct) and high Fe (1.1 wt pct), respectively. The nucleation temperatures for a-al (T a ), b-al 5 FeSi (T b ), and Al-Si eutectic (T N ) were determined by the intersections of the tangent of the derivative temperature curves in Figure 2, and the results are shown in Table II. The addition of Sr at all levels resulted in the depression of the eutectic-nucleation temperature, minimum temperature prior to recalescence, and growth temperature III. RESULTS Figures 2(a) and (b) show the cooling curves obtained during the solidification of both unmodified and Fig. 2 Cooling curves for the Al-10 pct Si alloys in the unmodified and 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 Strontium Added Alloy 1 Strontium Added Alloy 2 Temperature ppm Sr +290 ppm Sr +490 ppm Sr ppm Sr +110 ppm Sr +220 ppm Sr T a ( C) T b ( C) T N ( C) T G ( C) Notes: T a : nucleation temperature for a-al, T b : nucleation temperature for pre-eutectic b phase, T N : nucleation temperature for eutectic Si, and T G : growth temperature for eutectic Si. METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 39A, OCTOBER

4 (T G ) for both iron levels. As reported in the literature, [10] a binary Al-(b-Al 5 FeSi) eutectic reaction occurs prior to the eutectic Al-Si reaction when the iron level exceeds the critical level, which is 0.7 wt pct Fe for an Al-10 wt pct Si alloy. Pre-eutectic b-al 5 FeSi was observed to form in the 1.1 wt pct Fe containing alloys (T b in the curve of Figure 2(b)), whereas no pre-eutectic formation of b-al 5 FeSi phase was detected in the 0.5 wt pct Fe containing alloys (Figure 2(a)). Increasing the level of Sr in the 1.1 wt pct Fe alloys reduced the precipitation temperature of pre-eutectic b-al 5 FeSi, as well as the Al-Si eutectic-nucleation temperature (Table II). Figure 3 shows the microstructure of the alloys, which were quenched midway along the eutectic arrest. 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 dendrites, coarse flakelike silicon, and finer needle-shaped b-al 5 FeSi phases were commonly observed together with quenched liquid. Long b-al 5 FeSi platelets, which are 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 addition of less than 50 ppm Sr did not alter the morphology of the eutectic silicon significantly; however, well-refined and fibrous eutectic silicon was observed in the alloys with an excess of 100 ppm Sr (Figures 3(b) and (d)). Apart from the primary dendrites, eutectic silicon and b-al 5 FeSi, Sr-rich intermetallics were frequently observed near the quenched liquid-dendrite interface, for both Fe levels, in alloys well modified by 290 ppm Sr and 220 ppm Sr (Figures 4(a) and (b), respectively). Macrographs of all of the quenched specimens are shown 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 noticeable features. In these alloys, the eutectic grains are too small to be resolved in the macrographs, where the term, eutectic grains, refers to the connected Al-Si eutectic grains, which have originated from a common source. However, the Sr-modified alloys, containing in excess of 100 ppm of Sr, display circular eutectic grains in the interior of the specimens along with a layer of eutectic grains nucleated at the container wall. The frequency of eutectic grain nucleation is dramatically reduced with increased Sr levels at both Fe levels. The frequency of eutectic grain nucleation, moreover, seems to be dependent not only on the Sr content but also on the Fe content of the alloys. Increasing the Fe content of the alloys also resulted in a decrease in the number of eutectic grains (Figures 5(c) and (d)) vs Figures 5(g) and (h)). Provided that the Sr level and quenching time are almost identical, fewer eutectic grains appear to form in the 1.1 wt pct Fe alloys (compare Figures 5(c) and (h)). In the well-modified alloys, the decrease in the eutectic-nucleation frequency seems to correlate with the formation of Sr-rich intermetallics, which occasionally contain centrally located second-phase particles (Figure 4(a)). In the Al-10 wt pct Si-1.1 wt pct Fe-220 ppm Sr alloy, a compound containing phosphorus was found to be entrapped within the Sr-rich intermetallic phase by EPMA as shown in Figure 6. Thin-foil specimens of the cross-sectional plane of these internal particles were prepared by FIB milling and examined by TEM analysis. The enlarged inset in Figure 7(a) clearly shows secondphase particles contained within the Sr-rich intermetallic phase. These internal particles were found to be a connected single particle in three-dimensional observations using the consecutive FIB milling technique, although they (arrowed) appear to be disconnected in two-dimensional cross section. As shown in Figure 7(b), the EDS spectrum obtained from the internal particle in the Sr-rich phase shows a strong P peak, indicating it is a P-rich phase, possibly AlP. Along with the P peak, characteristic peaks of Al, Si, and Sr, which seem to be related to the Sr-rich intermetallic phase, Al 2 Si 2 Sr, underneath the P-rich particles, are also observed in Figure 7(b). However, no P peak was observed in the EDS spectrum obtained from the surrounding Al 2 Si 2 Sr phase (Figure 7(c)). Inside the Al 2 Si 2 Sr phase, oxide particles were often found in association with the P-rich phase near the surface polished by conventional procedure. It is unlikely that these oxide particles formed during cooling of the alloy and nucleated the Al 2 Si 2 Sr phase. Instead, the oxides are likely to form during the conventional polishing process, because AlP has been reported to react actively with water. A TEM micrograph of the cross-sectioned Al 2 Si 2 Sr phase containing the P-rich particles is shown in Figure 8(a). In Figure 8(b), the diffraction pattern obtained from the surrounding Sr-intermetallic phase confirms that the phase is Al 2 Si 2 Sr (hexagonal, P 3mL, a = nm, c = nm). Figure 9 shows a TEM image (Figure 9(a)) of Al 2 Si 2 Sr phase containing P-rich particles and its EDS spectrum (Figure 9(b)) obtained from the P-rich particle. The EDS map of the particle in Figure 9(c) reveals that the distribution of P corresponds well to the internal particles present in Figure 9(a), which is consistent with the results of the EDS analysis given in Figure 7. These internal P-containing particles are most likely AlP phase, and it is likely that this particle nucleated the Al 2 Si 2 Sr phase. As shown in Figure 9(b), an oxygen peak was detected along with phosphorus in large quantities in the spectrum obtained from the AlP particle during TEM analysis, whereas the EDS analysis during FIB milling shows no significant oxygen peak. It is, therefore, considered that oxidation of the AlP phase occurred rapidly when the thin-foil sample was taken out from the FIB chamber into a nonvacuum condition. This oxidation event was observed to be even further accelerated during TEM observation, which disrupted the crystalline nature of the AlP (Figure 9(d) presents the distribution of oxygen in the P-rich particles in the EDS map). Aluminum phosphide has been suggested to play a significant role in nucleating the b-al 5 FeSi phase, as well as eutectic silicon. [20,21,23] It is, therefore, expected that the addition of P may result in easy nucleation of b-al 5 FeSi phase by providing prolific nuclei, i.e., AlP particles. To investigate the effect of P on the formation of b-al 5 FeSi, three levels of P (10, 40, and 60 ppm) were added to the Al-10 wt pct Si base alloy containing 1.1 wt pct Fe. Characteristic reactions identified from 2438 VOLUME 39A, OCTOBER 2008 METALLURGICAL AND MATERIALS TRANSACTIONS A

5 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 pct Fe, (b) 0.5 pct Fe ppm Sr, (c) 1.1 pct Fe, and (d) 1.1 pct Fe ppm Sr. METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 39A, OCTOBER

6 Fig. 4 Optical micrographs showing Sr-rich intermetallics (arrowed) frequently observed near the dendrite-quenched liquid interface (a) in 0.5 pct Fe alloy ppm Sr and (b) 1.1 pct Fe alloy ppm Sr. the cooling-curve analysis for all three alloys are listed in Table III. The increased addition of P caused no significant changes in the nucleation temperature of primary aluminum phase (T a ), minimum temperature prior to recalescence (T min ), and eutectic nucleation (T N ), and growth temperatures (T G ). However, the nucleation temperature of the pre-eutectic b-al 5 FeSi, T b, increased from 584 C in the alloy without P addition to 589 C for the 60 ppm P containing alloy, steadily increasing with each P addition. The average area and length of the b-al 5 FeSi platelets were measured using an image analyzer and are plotted against P content in Figures 10(a) and (b), respectively. These results do not show any significant effect of increasing P levels. IV. DISCUSSION A. Nucleation of the Al-Si Eutectic Phase In Al-Si alloys, Sr is well accepted as a eutectic-silicon modifier, which effectively changes the morphology of silicon from platelike to fibrous at addition levels of a few hundreds parts per million. [1] Recent studies have confirmed that Sr addition affects not only the growth of the eutectic silicon but the nucleation behavior of the eutectic phases changes significantly, as well. [2 4,7 9] The addition of Sr in excess of the minimum amount required for full modification, which is suggested to be approximately 100 ppm for the alloys and the cooling rate adopted in this study, caused a large decrease in the nucleation frequency of eutectic grains and an associated depression of eutectic nucleation and growth temperatures (Table II). These observations are consistent with the literature. [3,8,9] In unmodified hypoeutectic Al-Si alloys, prolific nucleation events of Al-Si eutectic were observed to occur adjacent to the dendrite tips, whereas very few eutectic grains nucleated in the interdendritic regions in the Sr modified alloys (Figures 3 and 5). The reduction in nucleation frequency is at least an order of magnitude. Hunt [10] proposed that the columnar-to-equiaxed transition is also applicable to the formation of the eutectic during solidification, as well as to the primary dendritic growth. In unmodified alloys, eutectic Al-Si can nucleate and grow in the form of equiaxed grains when potent nucleants are present and sufficient thermal and constitutional undercooling is provided. As shown in Figures 3(a) and (c) and Figure 5, each eutectic colony, which was frequently found to form ahead of the primary dendrites, could be considered as an equiaxed-type grain. Crossley and Mondolfo [5] suggested that AlP particles, which are normally present in commercial Al-Si alloys, act as nuclei for eutectic silicon and that the nucleation of eutectic silicon can, therefore, occur easily with little undercooling due to the extremely good efficiency and low lattice mismatch between AlP and Si. They suggested that modifiers, i.e., sodium, neutralizes the AlP and prevents the easy nucleation of eutectic silicon, resulting in an increase of undercooling. Recently, Nogita et al. [6] provided conclusive evidence of the nucleation of eutectic silicon on AlP particles by TEM analysis of specimens prepared by FIB milling. Theoretically, Sr addition is expected to be rejected by a-al and to cause a solute buildup at the primary dendrite-liquid interface during solidification. Solute buildup could then increase the nucleation of equiaxed eutectic colonies due to increased constitutional undercooling. However, Sr addition was instead found to significantly decrease the number of eutectic grains (Figure 5). It is, therefore, more likely that the potency or number of effective nucleants for eutectic silicon in the melt was reduced. In hypoeutectic Al-Si alloys, it has been suggested that the decrease in the number of eutectic-nucleation events caused by Sr additions occurs due to poisoning of the extremely prolific nuclei-alp nucleant particles. [2,8,9,11] However, while the exact mechanism responsible for the poisoning by Sr addition has not been confirmed, the formation of Sr-rich intermetallics, Al 2 Si 2 Sr, in Sr-modified alloys has been proposed to be a contributing factor. Polygonal-shaped intermetallic Al 2 Si 2 Sr phase with a size of less than 20 lm was frequently observed in the liquid at elevated Sr levels (Figure 4). Moreover, Al 2 Si 2 Sr was found to 2440 VOLUME 39A, OCTOBER 2008 METALLURGICAL AND MATERIALS TRANSACTIONS A

7 Fig. 5 Macrostructure of samples quenched halfway along the eutectic arrest showing a decrease in the number of eutectic grains as additions of Sr and Fe are increased: (a) 0.5 pct Fe, (b) 0.5 pct Fe + 30 ppm Sr, (c) 0.5 pct Fe ppm Sr, (d) 0.5 pct Fe ppm Sr, (e) 1.1 pct Fe, (f) 1.1 pct Fe + 40 ppm Sr, (g) 1.1 pct Fe ppm Sr, and (h) 1.1 pct Fe ppm Sr. surround the AlP particles (Figures 6 and 7), which is a strong indication of the poisoning mechanism, i.e., the addition of Sr renders AlP nuclei ineffective by forming a layer of intermetallic phase around them. It is worth considering the early stage of the formation of Sr-rich intermetallics during solidification. Due to the small addition levels of Sr in this study, no information about the Al 2 Si 2 Sr reaction could be obtained from the cooling curves (Figure 2). Based on the assumption of Scheil conditions, solidification simulations were performed using the software Thermo- Calc (TCCQ, Thermo-Calc Software Inc.) [12] with the TTAL4 (Thermo Tech Ltd., Guildford, United Kingdom) database. The results in Figure 11 show that Al 2 Si 2 Sr is predicted to form prior to the formation of primary aluminum phase with the addition of 150 ppm Sr 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 diagram for Al-Si-Sr predicts that the Al 2 Si 2 Sr precipitation could first take place at 596 C when the Sr level exceeds 130 ppm. P-rich particles, possibly AlP phase, were generally present despite the alloys containing less than 10 ppm phosphorus in the present work. It has been reported previously that such P levels in commercial alloys are sufficient to produce some AlP. [13] When AlP forms early, the AlP particles may be pushed to the dendriteliquid interface as the a-al dendrites grow, and they can later nucleate the eutectic silicon at low undercooling in METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 39A, OCTOBER

8 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. unmodified alloys. Introduction of Sr to the alloys, on the other hand, caused formation of Al 2 Si 2 Sr intermetallics on the pre-existing AlP particles and, thus less, or less effective, nuclei are, therefore, available to nucleate the eutectic when the solidification path reaches the eutectic reaction. This explanation reasonably accounts for the large decrease in eutectic-nucleation frequency in the Sr-modified alloys. The quenched microstructures of the Sr-modified alloys in Figure 4 show the presence of isolated Sr-rich intermetallics near the dendrite-liquid interface. However, these Al 2 Si 2 Sr precipitates were not observed to nucleate eutectic silicon when the eutectic reaction commenced. In Sr-modified alloys, the spherical eutectic grains, which are several orders of magnitude larger than those in unmodified alloys, appear to nucleate on some other unidentified nuclei. It is not fully understood why Al 2 Si 2 Sr is ineffective for the nucleation of eutectic grains and what nucleates the eutectic grains in Sr-modified alloys. However, it is likely to be related to the increase in particle size when the intermetallics form on the AlP particles and possibly also to the fact that the formation of the intermetallics causes a local decrease in the silicon concentration in the melt. It is also possible that some AlP particles remain in the melt without nucleating Al 2 Si 2 Sr, although a larger undercooling is required compared to the unmodified alloys. Further studies are required to explain this mechanism in detail. B. Formation of Iron-Rich b-al 5 FeSi Intermetallics The addition of Sr to Al-Si alloys has also been reported to cause the refinement of the iron-rich intermetallics, b-al 5 FeSi. [18 22] It has been suggested that the effect of Sr in refining b-al 5 FeSi platelets is likely to involve dissolution and fragmentation. [17 22] 2442 VOLUME 39A, OCTOBER 2008 METALLURGICAL AND MATERIALS TRANSACTIONS A

9 Fig. 7 (a) SEM image of Sr-rich intermetallics containing internal particles (arrowed in the enlarged inset) in the FIB sample and corresponding EDS spectra obtained from (b) the internal particle (analyzed point, A) and (c) surrounding Al 2 Si 2 Sr phase (analyzed point, B). METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 39A, OCTOBER

10 Fig. 8 (a) TEM image of Al 2 Si 2 Sr phase containing P-rich particles (arrowed) and (b) corresponding electron diffraction pattern from the Al 2 Si 2 Sr phase (hexagonal, P 3mL, a = nm, c = nm) zone According to previous studies, the addition of Sr suppresses the branching of b-al 5 FeSi platelets by poisoning of preferential nucleation sites [18] and also by partial dissolution of b-al 5 FeSi platelets resulting in the fragmentation of platelets, [22] both causing a reduction in the length of the b-al 5 FeSi phase. However, the fragmentation mechanism remains uncertain. In hypoeutectic Al-Si alloys containing Fe in excess of the critical level, pre-eutectic b-al 5 FeSi platelets are generally formed by the binary Al-(b-Al 5 FeSi) eutectic reaction prior to the formation of the Al-Si eutectic. [12] The microstructure of the eutectic Al-(b-Al 5 FeSi) is expected to be irregular like that of the Al-Si eutectic, where the b-al 5 FeSi platelet is the leading phase and is likely to be covered by the Al dendrites during eutectic growth. Depending on the sectioning direction of the specimen, therefore, some segments of the b-al 5 FeSi platelets were exposed to the surface, whereas other segments were still found to be covered by the Al dendrites. This may make it appear as if the latter segments are fragmented or dissolved (Figure 4(b)), but it is tempting to suggest that these fragments are connected to a single branched b platelet. Sigworth et al. [23] suggested that the P (possibly AlP) present in the melt can nucleate b-al 5 FeSi and that Sr addition can suppress the precipitation of the b-al 5 FeSi phase. More recent studies on the formation of b-al 5 FeSi, on the other hand, proposed that Fe-rich intermetallics are likely to form on the externally wetted surface of the oxide film, which are entrained into the melt during casting. [14,15] It has been reported that the wetted surface of oxide films provides a preferential nucleation site for b-al 5 FeSi and a large oxide film folded to a dry side was frequently found to form cracklike defects inside b-al 5 FeSi. However, a careful inspection of the microstructure of the specimens in the present study showed no direct evidence of the presence of cracks inside the b-al 5 FeSi platelets. The samples in the present work were filled by carefully submerging the stainless-steel cups into the skimmed melt followed by cooling, and thus may not facilitate the formation and entrainment of significant oxide films in the samples. Therefore, the role of the oxide films in nucleating the b-al 5 FeSi phase appears to be negligible in the present work. The addition of up to 60 ppm P caused an increase in the nucleation temperature of b-al 5 FeSi, T b (Table III). This suggests that the increase in P level provides a larger number of P-based nuclei, and thus b-al 5 FeSi nucleates more easily at a smaller undercooling. It is well-established that the refinement of primary silicon in hypereutectic Al-Si alloys occurs with the addition of potent AlP nucleants to the melt, [27] and thus prolific nucleation events of b-al 5 FeSi on a larger number of nuclei could also be expected to result in a refinement of the size of b-al 5 FeSi. The results in Figure 10 show no clear decreasing or increasing trend in the area and size of b-al 5 FeSi with increasing P addition. It is worth considering this result in more detail from the viewpoint of eutectic nucleation and growth. Flood and Hunt [16] reported that the average velocity of a eutectic interface is inversely proportional to the total solid/liquid interface area of the eutectic growth front. Unlike the growth of primary Si in the melt, pre-eutectic b-al 5 FeSi forms via a binary Al-(b-Al 5 FeSi) eutectic reaction. The growth of b-al 5 FeSi may thus occur according to the theory proposed by Flood and Hunt. Assuming that AlP nucleates a large number of b-al 5 FeSi, the total area of solid/liquid interface is increased by P addition, and the interface velocity of eutectic Al-(b-Al 5 FeSi) is, therefore, decreased compared to the case with little nucleation. A reduction of the growth rate is expected to result in a coarser microstructure, i.e., coarser b-al 5 FeSi intermetallics. Increased addition of Sr to the Al-10 wt pct Si-1.1 wt pct Fe alloys gradually decreased the formation temperature of b-al 5 FeSi phase, T b (Table II). The reason for the 2444 VOLUME 39A, OCTOBER 2008 METALLURGICAL AND MATERIALS TRANSACTIONS A

11 Fig. 9 (a) TEM image of Al 2 Si 2 Sr phase containing P-rich particles and (b) EDS spectrum obtained from P-rich particles, with EDS maps of interest region (squared) showing the distribution of (c) phosphorus and (d) oxygen. The distribution of oxygen is likely to correspond with the phosphorus distribution, which appears to be due to a rapid-oxidation event of the P-rich particle. decrease of T b brought about by Sr addition is not fully understood. However, it is possible that it is caused by the poisoning of the b-al 5 FeSi nucleation sites by the formation of Al 2 Si 2 Sr on AlP particles. Early formation of Al 2 Si 2 Sr on P-rich particles may make them ineffective for the nucleation of b-al 5 FeSi, just as for the nucleation of eutectic silicon. The result of the Scheil simulation given in Figure 11(b) shows that the formation of b-al 5 FeSi platelets commences after the formation of Al 2 Si 2 Sr and the development of the primary METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 39A, OCTOBER

12 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 Phosphorus Added Alloy 2 Temperature ppm P ppm P ppm P T a ( C) T b ( C) T min ( C) T N ( C) T G ( C) Notes: T a : nucleation temperature for a-al, T b : nucleation temperature for pre-eutectic b phase, T min : minimum temperature prior to recalescence, T N : nucleation temperature for eutectic Si, and T G : growth temperature for eutectic Si. Fig. 10 The effect of phosphorus addition on (a) average area and (b) average length of the b-al 5 FeSi platelets in Al-10 pct Si-1.1 pct Fe alloys. aluminum. During solidification, primary Al 2 Si 2 Sr forms on AlP particles, which are also potent nuclei for b-al 5 FeSi platelets. Fewer, or less effective, nuclei are, therefore, available to nucleate b-al 5 FeSi and, consequently, a larger undercooling is required. Because AlP particles are potent nucleant substrates for both eutectic Al-Si and b-al 5 FeSi, it is worth considering the inter-relationship of the nucleation events between Al-Si eutectic and b-al 5 FeSi. Recent work on the interaction between Fe and Al-Si eutectic in hypoeutectic Al-Si alloys revealed that the number of Al-Si eutectic-nucleation events decreases as the Fe Fig. 11 Scheil simulation calculated by Thermo-Calc combined with 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 pct Fe-150 ppm Sr alloy. Labeled numbers along the curve correspond to a certain range of temperature where stable phases can be present VOLUME 39A, OCTOBER 2008 METALLURGICAL AND MATERIALS TRANSACTIONS A

13 content increases. [17] Dinnis et al. [19] proposed that this interaction is due to the fact that the eutectic Al-Si and b-al 5 FeSi have common nuclei, AlP. As shown in Figures 5(c) and (h), the number of eutectic grains in the Sr-modified alloys containing 1.1 wt pct Fe is still less than that in the Sr- modified alloys containing 0.5 wt pct Fe, despite a similar concentration of Sr in both alloys. This suggests that the formation of pre-eutectic b-al 5 FeSi platelets in the higher Fe-content alloys reduces the nucleation frequency of eutectic Al-Si grains by the formation of b-al 5 FeSi on AlP, in accordance with the observations by Dinnis et al. [19] The nucleation sequences in unmodified and Sr-modified alloys containing low Fe (0.5 wt pct) and high Fe (1.1 wt pct) are schematically illustrated in Figure 12. If sufficient P is present in the melt, AlP particles form early and are pushed ahead of the dendriteliquid interface during solidification. In unmodified alloys with low-fe content (Figure 12(a)), the eutectic Al-Si reaction commences as AlP nucleates the polygonal-shaped eutectic silicon near the dendrite tips where the silicon level is locally high enough. In Sr-containing alloys, on the other hand, Al 2 Si 2 Sr is likely to form on pre-existing AlP prior to the eutectic Al-Si reaction, so AlP particles would not play a significant role in the nucleation of the eutectic grains and far fewer eutectic grains form in the interdendritic liquid (Figure 12(c)). When the Fe content is high enough to form pre-eutectic b-al 5 FeSi platelets, AlP can nucleate both b-al 5 FeSi platelets and eutectic silicon. Because the pre-eutectic b-al 5 FeSi platelets are formed on AlP particles prior to the eutectic Al-Si reaction, fewer nuclei are available for nucleating the Al-Si eutectic, resulting in fewer Al-Si grains (Figure 12(b)). The addition of Sr, along with high Fe content, reduces the nucleation of b-al 5 FeSi platelets, as well as eutectic Al-Si grains, by the formation of Al 2 Si 2 Sr phase on AlP (Figure 12(d)). Furthermore, even though the addition of Sr is not sufficient for the Al 2 Si 2 Sr phase to consume all of the AlP particles in the liquid, it is expected that the precipitation of b-al 5 FeSi platelets onto the remaining AlP results in even less nucleation of eutectic Al-Si grains (Figure 12(d)). This conclusion is strongly supported by the observation that the decrease in the number of eutectic grains in the alloy with lower Sr and higher Fe content, the 110 ppm Sr in Al-10 wt pct Si-1.1 wt pct Fe alloy, is more significant than that in the alloy with higher Sr and lower Fe content, the 290 ppm Sr in Al-10 wt pct Si-0.5 wt pct Fe alloy. V. CONCLUSIONS The effect of Sr addition on eutectic Al-Si nucleation and the formation of b-al 5 FeSi in hypoeutectic Al-Si foundry alloys were investigated. Sr additions exceeding 100 ppm dramatically reduced the number of eutectic Al-Si grains and decreased the eutectic-nucleation temperature. Poisoning of the potent AlP nuclei for the nucleation of eutectic grains is proposed as the mechanism. The result shows that Sr forms Al 2 Si 2 Sr intermetallic phase onto the AlP Fig. 12 Schematic illustration of solidification sequence with the formation of Al 2 Si 2 Sr, primary Al dendrite, pre-eutectic b-al 5 FeSi, and eutectic Al-Si: (a) AlP present at the dendrite liquid interface nucleates eutectic silicon in unmodified alloy with low Fe content; (b) both eutectic silicon and b-al 5 FeSi nucleate on AlP when the Fe addition to unmodified alloys is high enough to form pre-eutectic b-al 5 FeSi platelets; (c) Sr-rich Al 2 Si 2 Sr intermetallics present in Sr-modified alloys with low Fe content form on AlP and far fewer eutectic grains (spherical), which do not nucleate on AlP grow in the interdendritic liquid; and (d) both Al 2 Si 2 Sr and pre-eutectic b-al 5 FeSi nucleate on AlP, and thus even fewer eutectic grains will form in Sr-modified alloys with high Fe content. METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 39A, OCTOBER

14 particles. The potency of AlP as nuclei for the Al-Si eutectic is thereby reduced, and, therefore, the nucleation of eutectic grains is reduced. Aluminum phosphide is proposed to be a common nucleation site for both eutectic Al-Si grains and b-al 5 FeSi phase. Sr addition gradually decreased the b-al 5 FeSi nucleation temperature. The formation of b-al 5 FeSi appears to be suppressed by the presence of Al 2 Si 2 Sr, which is believed to also be caused by deactivation of the AlP nuclei. The Al-Si eutectic-nucleation frequency was decreased by increasing the Fe concentration in excess of the critical level for pre-eutectic b-al 5 FeSi nucleation in the Sr- modified alloys. The Scheil simulations showed that the formation of pre-eutectic b-al 5 FeSi occurs after the formation of Al 2 Si 2 Sr and primary Al dendrites. It is suggested that AlP particles present in the melt are consumed by the formation of Al 2 Si 2 Sr, as well as the nucleation of b-al 5 FeSi in the early stages of solidification. It is, therefore, concluded that the decrease in nucleation frequency of eutectic Al-Si grains is caused by a lack of efficient nuclei when the solidification process reaches the eutectic Al-Si reaction. ACKNOWLEDGEMENTS The authors thank Ms. H.K. Kang for her help in EPMA analysis and Mr. D.H. Kim for the preparation of TEM using FIB. The authors also thank Drs. S. McDonald and K. Nogita, University of Queensland, for their help in lab experiments and many valuable discussions. REFERENCES 1. S.Z. Lu and A. Hellawell: Metall. Trans. A, 1987, vol. 18A, pp S.D. McDonald: Ph.D. Dissertation, The University of Queensland, Brisbane, Australia, 2002, pp A.K. Dahle, K. Nogita, J.W. Zindel, S.D. McDonald, and L.M. Hogan: Metall. Mater. Trans. A, 2001, vol. 32A, pp K. Nogita, S.D. McDonald, C. Dinnis, L. Lu, and A.K. Dahle: in Solidification of Aluminum Alloys, Proc. Symp., M.G. Chu, D.A. Granger, and Q. Han, eds., Charlotte, NC, Mar , 2004, TMS, Warrendale, PA, pp P.B. Crosely and L.F. Mondolfo: AFS Trans., 1966, vol. 74, pp K. Nogita, S.D. McDonald, K. Tsujimoto, K. Yasuda, and A.K. Dahle: J. Electron. Microsc., 2004, vol. 53, pp K. Nogita, P.L. Schaffer, S.D. McDonald, L. Lu, and A.K. Dahle: Aluminium, 2005, vol. 81, pp S.D. McDonald, A.K. Dahle, J.A. Taylor, and D.H. StJohn: Metall. Mater. Trans. A, 2004, vol. 35A, pp S.D. McDonald, K. Nogita, and A.K. Dahle: Acta Mater., 2004, vol. 52, pp J.A. Taylor, G.B. Schaffer, and D.H. StJohn: Metall. Mater. Trans. A, 1999, vol. 30A, pp J.A. Taylor, G.B. Schaffer, and D.H. StJohn: Metall. Mater. Trans. A, 1999, vol. 30A, pp A.M. Samuel, F.H. Samuel, C. Villeneuve, H.W. Doty, and S. Valtierra: Int. J. Cast Met. Res., 2001, vol. 14, pp C.M. Dinnis, J.A. Taylor, and A.K. Dahle: Mater. Forum, 2004, vol. 28, pp A. Couture: Int. Cast Met. J., 1981, vol. 6, pp S.G. Shabestari and J.E. Gruzleski: Cast Met., 1994, vol. 6, pp A.N. Lakshmanan, S.G. Shabestari, and J.E. Gruzleski: Z. Metallkd., 1995, vol. 86, pp S.G. Shabestari, M. Mahmudi, M. Emamy, and J. Campbell: Int. J. Cast Met. Res., 2002, vol. 15, pp A.M. Samuel, F.H. Samuel, and H.W. Doty: J. Mater. Sci., 1996, vol. 31, pp F.H. Samuel, P. Ouellet, A.M. Samuel, and H.W. Doty: Metall. Mater. Trans. A, 1998, vol. 29A, pp A. Pennors, A.M. Samuel, F.H. Samuel, and H.W. Doty: AFS Trans., 1998, vol. 105, pp A.M. Samuel, A. Pennors, C. Villeneuve, F.H. Samuel, H.W. Doty, and S. Valtierra: Int. J. Cast Met. Res., 2000, vol. 13, pp A.M. Samuel, F.H. Samuel, C. Villeneuve, H.W. Doty, and S. Valtierra: Int. J. Cast Met. Res., 2001, vol. 14, pp G.K. Sigworth, S. Shivkumar, and D. Apelian: AFS Trans., 1989, vol. 97, pp J.D. Hunt: Mater. Sci. Eng., 1984, vol. 65, pp A.K. Dahle, K. Nogita, S.D. McDonald, C. Dinnis, and L. Lu: Mater. Sci. Eng. A-Struct., 2005, vols , pp B. Jansson, M. Schalin, M. Selleby, and B. Sundman: in Computer Software in Chemical and Extractive Metallurgy, C.W. Bale and G.A. Irins, eds., Canadian Institute of Metals, Quebec, Canada, 1993, pp L.F. Mondolfo: Aluminum Alloy: Structure and Properties, Butterworth and Co, London, 1976, pp X. Cao and J. Campbell: Metall. Mater. Trans. A, 2003, vol. 34A, pp X. Cao and J. Campbell: Int. J. Cast Met. Res., 2000, vol. 13, pp S.C. Flood and J.D. Hunt: Met. Sci., 1981, vol. 15, pp C.M. Dinnis, J.A. Taylor, and A.K. Dahl: Metall. Mater. Trans. A, 2006, vol. 37A, pp VOLUME 39A, OCTOBER 2008 METALLURGICAL AND MATERIALS TRANSACTIONS A