- Email: [email protected]
Eutectic modification and microstructure development in Al–Si Alloys
Materials Science and Engineering A 413–414 (2005) 243–248
Eutectic modification and microstructure development in Al–Si Alloys A.K. Dahle a,b,∗ , K. Nogita b , S.D. McDonald b , C. Dinnis a,b , L. Lu c a
CRC for Cast Metals Manufacturing, The University of Queensland, Brisbane, Qld 4072, Australia b Materials Engineering, The University of Queensland, Brisbane, Qld 4072, Australia c CSIRO Minerals, P.O. Box 883, Kenmore, Qld 4069, Australia Received in revised form 29 June 2005
Abstract Recent increasing applications for cast Al–Si alloys are particularly driven by the need for lightweighting components in the automotive sector. To improve mechanical properties, elements such as strontium, sodium and antimony can be added to modify the eutectic silicon from coarse and plate-like to fine and fibrous morphology. It is only recently being noticed that the morphological transformation resulting from eutectic modification is also accompanied by other, equally significant, but often unexpected changes. These changes can include a 10-fold increase in the eutectic grain size, redistribution of low-melting point phases and porosity as well as surface finish, consequently leading to variations in casting quality. This paper shows the state-of-the-art in understanding the mechanism of eutectic nucleation and growth in Al–Si alloys, inspecting samples, both quenched and uninterrupted, on the macro, micro and nano-scale. It shows that significant variations in eutectic nucleation and growth dynamics occur in Al–Si alloys as a function of the type and amount of modifier elements added. The key role of AlP particles in nucleating silicon is demonstrated. © 2005 Elsevier B.V. All rights reserved. Keywords: Al–Si eutectic; Modification; Eutectic nucleation; Eutectic grain size
1. Introduction Eutectic modification is a common process performed in aluminium–silicon based foundry alloys primarily to improve mechanical properties, particularly tensile elongation, by promoting a structural refinement of the inherently brittle eutectic silicon phase. It is well known that trace additions of strontium (a few hundred parts per million) to hypoeutectic aluminium–silicon alloys result in a transformation of the eutectic silicon morphology from a coarse plate-like structure to a well-refined fibrous structure [1]. Additions of antimony (albeit at higher concentrations of a few thousand parts per million) also give rise to a transformation of the eutectic silicon phase, however the result is not as dramatic with the coarse plates being refined rather than transformed to a fibrous morphology. Typical examples of the microstructure of unmodified, Sr-modified and Sb-modified alloys are shown in Fig. 1. Other elements including sodium and several rare earth elements are also known to result in varying degrees of modification [2].
∗
Corresponding author. Tel.: +61 7 3365 4350; fax: +61 7 3365 3888. E-mail address: [email protected] (A.K. Dahle).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.09.055
It appears to be universally accepted that the change in silicon morphology that occurs with modification is desirable. Unfortunately, modification has also been associated with negative side effects such as porosity, hot tearing and poor surface finish and is not always a recommended procedure. In recent years, it has become apparent that these complications occur because the change in silicon morphology does not occur in isolation, but is accompanied by a number of other significant changes in solidification and microstructure development. This paper compares the development of the cast structure on both macroscopic and microscopic scales in unmodified and modified alloys (both strontium and antimony). The interrelationship between nucleation and growth is demonstrated. 2. Experimental methods Several different experimental techniques have been employed throughout the research [3–5] and these are only briefly outlined in this paper. The base metal components, usually commercial purity Al and Si, were melted in an electrical resistance furnace and kept at about 720 ◦ C. A nominal composition of Al–10 wt.% Si was normally used because it contains a large volume fraction of eutectic, facilitating macroscopic inves-
244
A.K. Dahle et al. / Materials Science and Engineering A 413–414 (2005) 243–248
Fig. 1. Comparison of the silicon morphology in (a) unmodified, (b) Sr-modified (300 ppm Sr) and (c) Sb-modified (2400 ppm Sb), hypoeutectic aluminium–silicon alloys.
tigation, while still being a hypoeutectic alloy. Following stirring and surface skimming, the elements to be investigated were added. These included Sr as AlSr10 master alloy rods and Sb as AlSb10 ingots. Analysis of cooling curves was conducted according to the method described in detail elsewhere [3,4]. Additions in the range of 0–1000 ppm Sr and 0–5000 ppm Sb were examined. Samples were also quenched at different stages during the eutectic arrest. For these experiments, small tapered stainlesssteel cups coated with a thin layer of boron nitride were used to sample the melt. Two samples were taken in parallel by submerging the cups into the melt. The cups were placed on an insulated base, a thermocouple was positioned in the centre of one cup only, and insulation boards were placed on top of both samples. The samples were allowed to cool in air resulting in a cooling rate in the liquid just prior to nucleation of the primary aluminium of approximately 1.7 K/s, and a total solidification time of approximately 300 s. During solidification the cooling curve from the thermocouple was monitored on a real-time display, and the sample without a thermocouple was quenched approximately 50% of the way through the eutectic reaction. Although it is acknowledged that solidification will not proceed identically in both cups, this method prevents the thermocouple from interfering with solidification, or from damaging the microstructure during quenching. Samples were analyzed by optical and scanning electron microscopy including electron back-scattered diffraction (EBSD). To obtain three-dimensional images of the microstructures, selected quenched samples of unmodified and Sr-modified alloys were then subjected to serial sectioning. The unmodified and strontium-modified samples chosen for serial sectioning were quenched 45 and 12 s after eutectic nucleation, respectively. Serial sectioning involves the gradual removal of parallel layers of material and the imaging of each layer. Digital images of each section are “stacked” by using appropriate software to reconstruct the 3D microstructure. A full description of the procedure for producing 3D images is contained in Ref. [5]. Selected eutectic grains of unmodified samples with visible internal particles (potential nuclei) were chosen for further analysis using focused ion beam milling (FIB) followed by transmission electron microscopy (TEM). It is only recently that the
advanced techniques of FIB milling combined with in situ sample micro-manipulation have become available to selectively extract a TEM sample from any desired region within a larger sample. The FIB sample preparation was performed with a 30 kV Ga liquid metal ion source. TEM observations, micro-selected area diffractions (SAD) taken from the area of about 60 nm in diameter, small probe convergent beam electron diffraction patterns (CBED) taken from an area of ∼10 nm in diameter and EDX analysis were conducted using a 200 kV FE-TEM, see Ref. [6] for further details. 3. Results Typical cooling curves in the vicinity of the eutectic reaction are shown for the alloys used in this research in Fig. 2. Antimony additions resulted in a depression of the eutectic nucleation and growth temperatures and an increase in the amount of recalescence prior to growth. Similar but larger effects were observed with strontium modification. Fig. 3 summarises the defining micro- and macroscopic characteristics of unmodified, strontium-modified and antimonymodified alloys. Twinning in silicon fibres was examined using transmission electron microscopy for unmodified and strontiummodified alloys and results from the literature [7] for Sb have been included in Fig. 3a for comparison. In all alloys, co-zonal twinning is present in the silicon fibres with twins lying paral-
Fig. 2. Comparison of cooling curves from unmodified, Sb-modified (2600 ppm) and Sr-modified (200 ppm), Al–10% Si samples. The nucleation temperature (Tn ), minimum temperature prior to recalescence (Tmin ) and the growth temperature (Tg ) are included in the figure.
A.K. Dahle et al. / Materials Science and Engineering A 413–414 (2005) 243–248
245
Fig. 3. Comparison of unmodified, Sr-modified and Sb-modified structures. (a) TEM images of silicon crystals (note Sb-modified micrograph reproduced from a directionally solidified alloy in other research [7]), (b) EBSD maps illustrating the orientation of the eutectic aluminium relative to the primary aluminium dendrites, (c) 3D reconstruction of eutectic grains derived from serial sectioning of quenched samples (the unmodified sample is 85 m deep, the modified is 117 m deep) (d) optical micrographs of quenched samples (refined eutectic is quenched liquid) and (e) macrographs of quenched samples. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of the article.)
246
A.K. Dahle et al. / Materials Science and Engineering A 413–414 (2005) 243–248
lel to the apparent growth direction of the fibres. As shown in Fig. 3a, it was possible to find silicon crystals that were free of twinning in the unmodified alloys. Electron back-scattered diffraction was used to analyse the orientation of eutectic aluminium relative to that of the primary aluminium and typical results are shown in Fig. 3b. In unmodified alloys the vast majority of eutectic aluminium has an orientation identical to that of the surrounding dendrites, while in Sr-modified and Sb-modified alloys the eutectic aluminium has multiple orientations unrelated to the surrounding dendrites. These results indicate that the growth relation between the primary aluminium and the eutectic aluminium is epitaxial in the unmodified commercial alloys, while independent nucleation and growth occurs in the Sr- and Sb-modified alloys. There was a dramatic difference in the size of the eutectic grains. The reconstructed region of the unmodified alloy contained five separate eutectic grains (blue and red in Fig. 3c, unmodified). Of the two largest grains in the unmodified alloy; one was completely traversed by the series of sections while the other extended further into the sample. On the other hand, the Srmodified eutectic grain was very large and it appears as though only a small segment of this larger spherical grain has been reconstructed. This grain has been estimated to be approximately 500 m in diameter [5]. The modified eutectic silicon was too fine to reconstruct separately from the eutectic aluminium, so the two phases are represented as a single feature (red in Fig. 3c, Sr-modified). Fig. 3d shows optical micrographs of eutectic grains in samples that have been quenched at similar stages early during eutectic solidification. The eutectic grains in the unmodified alloy are at most a few hundred microns in diameter and consist of relatively few coarse silicon plates and a loosely coupled growth interface leading to grains that are often anisotropic in appearance. In the Sr-modified alloys, the eutectic grains are roughly circular in cross section and are typically much larger than the unmodified alloys (up to an order of magnitude). Each eutectic grain in the Sr-modified alloys contains a high density of silicon fibres. In the Sb-modified alloys, the eutectic grains are of an intermediate size and morphology, being more spherical than the unmodified grains. Each grain in the Sb-modified alloy contains numerous silicon plates, which are refined and more closely aligned than those found in the unmodified alloy. The difference in eutectic grain size between the unmodified, Sr-modified and Sb-modified alloys is further apparent in the macrographs of samples quenched early during eutectic solidification (Fig. 3e). In the unmodified sample, the eutectic grain size is so small that no grains are resolvable on the macrograph. In the Sr-modified alloy, there is a layer of coalesced eutectic grains lining the wall of the sample and several large grains can be seen independently distributed throughout the centre of the casting. In the Sb-modified alloy, the size and number of grains is intermediate to that found in the unmodified and strontiummodified alloys. In unmodified alloys, the eutectic grains were often observed to radiate from centrally located polyhedral silicon crystals, as seen in Fig. 3d. Occasionally, these polyhedral silicon crystals also have centrally located particles, which when analyzed using
Fig. 4. TEM image (top of micrograph is the exposed surface of the conventionally polished sample) and corresponding electron diffraction patterns from the Al–P–O precipitate on polished surface showing amorphous structure and small probe CBED patterns showing the AlP [0 1 1] pole and the silicon [0 1 1] pole. CBED patterns were obtained using an identical tilt angle.
conventional SEM techniques were found to contain Al, P and O [8]. In the current research, focused ion beam milling was used to selectively extract samples suitable for TEM analysis that contained both the central silicon crystal and its assumed nucleus. A typical TEM sample produced by this method along with SAD and CBED patterns is shown in Fig. 4. At some distance from the exposed conventionally prepared surface of the sample, no oxygen is present and the nucleus chemistry is consistent with AlP. Further to this, FEG-TEM analysis shows that there is no lattice mismatch between the (1 1 1) lattice planes of AlP and the Si crystal as shown in Fig. 5. Equivalent analysis on modified samples is challenging as it is difficult to locate possible nucleant particles during optical microscopy. Such difficulties are expected as statistically, the chance of sectioning through the exact center of a eutectic grain decreases dramatically as the nucleation frequency of eutectic grains decreases and their corresponding size increases.
Fig. 5. A high resolution FEG-TEM micrograph of the AlP/Si interface along with selected area diffraction patterns from regions A (Si) and B (AlP). There is no lattice mismatch between the (1 1 1) planes of AlP and Si.
A.K. Dahle et al. / Materials Science and Engineering A 413–414 (2005) 243–248
4. Discussion Contrary to expectations from the literature [9–11] it was found in this research that a higher degree of modification is not always associated with an increased density of non-cozonal twinning [12] and in fact this type of twinning was rarely observed in any of the alloys examined. This casts some doubt on the well accepted theory that the fibrous nature of the silicon in modified alloys is due solely to a high density of impurity induced twins imparting near isotropic growth properties to the silicon phase. It is possible that different growth mechanisms may operate or dominate in directionally solidified alloys, which have formed the basis for much of the pioneering TEM research into the growth mechanisms of eutectic silicon. While the growth mechanism of individual eutectic silicon fibres is likely to be important with respect to the fundamental modification mechanism, it is likely to be less critical for the processing characteristics of the solidifying melt. The difference in the nucleation frequency of eutectic grains between unmodified and modified alloys is significant. In unmodified alloys a large number of eutectic grains nucleate and these grow to only a small size prior to coalescence. In contrast, in Sr-modified alloys only a few grains nucleate and growth occurs over a significant distance prior to the coalescence of adjoining grains. In Sb-modified alloys, the grain size is intermediate between that of unmodified and Sr-modified alloys. From Fig. 3 it appears that there is a relationship between the degree of modification and the nucleation frequency of eutectic grains. The relationship appears to be that increasing the number of eutectic grains that nucleate results in a coarsening in the eutectic silicon morphology. This poses the question: Can the silicon morphology in modified alloys be simply related to the number of eutectic grains that nucleate? Certainly, as emphasised by Flood and Hunt [13] for a given rate of heat extraction the factor controlling the growth velocity of the eutectic is the solid–liquid interface area, which is directly controlled by the nucleation frequency. For a constant rate of heat extraction the growth velocity will vary inversely with the total solid–liquid surface area of the system. Therefore, the more grains that nucleate, the larger the solid–liquid surface area and the lower the growth velocity. However, although modification will cause an increase in the growth velocity during eutectic solidification, it has been shown that this increase in velocity is not sufficiently large to cause a flake-fibre transition in the eutectic silicon [8]. At this stage it must be conceded that the fundamental mechanism responsible for silicon modification is still unclear. The nucleation differences that occur between unmodified alloys and those modified with Sr or Sb are supported by the various analysis methods used in this research. Firstly, the eutectic nucleation and growth temperatures decrease when comparing unmodified, Sb-modified and Sr-modified alloys, respectively (Fig. 2). This implies an increasing degree of nucleation difficulty with Sb and Sr additions and in the case of these alloys, this seems to correspond well with the fact that fewer grains are present in the macrographs (Fig. 3e). The differences in the relative crystallographic orientations of eutectic and primary aluminium also support a difference in nucleation mechanism.
247
In the unmodified alloys it is proposed that the eutectic grains nucleate in the vicinity of aluminium dendrites and grow only a small distance, allowing the eutectic aluminium to adopt and maintain the same crystallographic orientation as the adjoining dendrite. In the modified alloys, the eutectic grains either nucleate independently of the aluminium dendrites or the large growth distance required of the larger grain, necessitates repeated nucleation during growth, giving rise to multiple crystallographic orientations of the aluminium within each eutectic grain. The difference in size of the eutectic grains is further confirmed with the results of the three dimensional reconstruction of the serial sectioning data (Fig. 3c). Despite a scarcity of research, it is likely that the size and distribution of eutectic grains in aluminium–silicon alloys will influence a variety of in-service and processing properties. These include surface appearance, corrosion resistance, mechanical properties and susceptibility to casting defects such as porosity and hot-tearing. This is because the distribution of the last few percent of liquid (important for porosity and hot-tear formation) and the location of ternary elements that precipitate during or after eutectic Al–Si solidification depends on the number and size of the eutectic grains. In hypoeutectic alloys, there is a significant proportion of aluminium dendrites present in the microstructure prior to eutectic solidification. It follows that the nucleation problem for the commencement of the eutectic reaction lies with the nucleation of the silicon phase. In the unmodified alloy, the result that AlP is the nuclei for eutectic silicon (Figs. 4 and 5) is not surprising, as both AlP and silicon have cubic structures with near identical ˚ [14] and 5.431 A ˚ [15], respeclattice parameters of 5.421 A tively. AlP was also identified as a good nucleant for eutectic silicon in the pioneering work on eutectic nucleation in Al–Si alloys by Crosley and Mondolfo [16]. Other nuclei have also been reported for unmodified silicon crystals and these include arsenic, sulphur, selenium and tellurium [17,18] and other possibilities include iron-containing phases and oxides introduced during processing. The decrease in nucleation frequency with strontium and antimony additions suggests that strontium somehow removes, or neutralizes, nuclei that are present in the unmodified alloy. There is some evidence to suggest that the poisoning mechanisms involves the nuclei (of which AlP appears to be the most prolific) being saturated by strontium containing intermetallics, most likely Al2 Si2 Sr which are frequently present in strontiummodified commercial aluminium–silicon alloys [19–21]. Similarly, Sb based intermetallics may also be important in removing phosphorus from the melt prior to eutectic solidification in Sbmodified alloys. At the extremity, modification of Al–Si alloys with Na has been shown to inhibit bulk eutectic nucleation and instead promote growth of a massive eutectic interface opposite to the thermal gradient [22]. 5. Conclusions From the current research it must be concluded that the fundamental mechanism of eutectic modification in Al–Si alloys (change in silicon morphology) is still not fully understood.
248
A.K. Dahle et al. / Materials Science and Engineering A 413–414 (2005) 243–248
Mechanisms based on the introduction of high densities of noncozonal twins may be more relevant to directionally solidified alloys than samples solidified under conditions of equiaxed growth. There are distinct differences in eutectic nucleation between unmodified and modified alloys and these are reflected in cooling curves, EBSD and micro- and macroscopic investigations. In the alloys used in this research there was increasing nucleation difficulties after individual additions of Sr or Sb to an unmodified alloy. This was reflected in the eutectic grain sizes with the eutectic grains being largest in the Sr-modified alloys and of an intermediate size in Sb-modified alloys. The differences in nucleation patterns are likely to be responsible for many of the complications commonly associated with modification. The compound AlP is a common nucleus for eutectic Si in unmodified alloys, but appears to be less active in modified alloys. The mechanism for this transition is unclear but may relate to scavenging of phosphorus by intermetallics introduced with the modification process. Although the eutectic grain size and silicon morphology are linked, they are not completely dependent. References [1] S. Lu, A. Hellawell, Metall. Trans. A 18 (1987) 1721–1733. [2] K. Nogita, S.D. McDonald, A.K. Dahle, Mater. Trans. 45 (2004) 323–326.
[3] K. Nogita, A. Knuutinen, S.D. McDonald, A.K. Dahle, J. Light Met. 1 (2001) 219–228. [4] A.K. Dahle, K. Nogita, J.W. Zindel, S.D. McDonald, L.M. Hogan, Metall. Mater. Trans. A 32A (2001) 949–959. [5] C.M. Dinnis, A.K. Dahle, J.A. Taylor, Mater. Sci. Eng. A 392 (2005) 440–448. [6] K. Nogita, S.D. McDonald, K. Tsujimoto, K. Yasuda, A.K. Dahle, J. Electron Microsc. 53 (2004) 361–369. [7] S. Khan, R. Elliott, J. Mater. Sci. 29 (1994) 736–741. [8] S.D. McDonald, A.K. Dahle, J.A. Taylor, D.H. St. John, Metall. Mater. Trans. A 35A (2004) 1829–1837. [9] M. Shamsuzzoha, L.M. Hogan, Philos. Mag. A 54 (1986) 459–477. [10] S. Lu, A. Hellawell, J. Cryst. Growth 73 (1985) 316–328. [11] M. Shamsuzzoha, L.M. Hogan, J. Mater. Sci. 24 (1989) 2849–2859. [12] K. Nogita, J. Drennan, A.K. Dahle, Mater. Trans. 44 (2003) 625– 695. [13] S.C. Flood, J.D. Hunt, Met. Sci. 15 (1981) 287–294. [14] M.E. Straumanis, E.Z. Aka, J. Appl. Phys. 23 (1952) 330. [15] C. Yeh, Z.W. Lu, S. Froyen, A. Zunger, Phys. Rev. B 46 (1992) 10086–10095. [16] P.B. Crosley, L.F. Mondolfo, AFS Trans. 74 (1966) 53–64. [17] L.F. Mondolfo, Aluminium Alloys: Structure and Properties, Butterworths, London-Boston, 1976. [18] T.B. Abbott, B.A. Parker, Cast Met. 1 (1988) 122–132. [19] G. Chai, L. B¨ackerud, AFS Trans. 100 (1992) 847–854. [20] G. Laslaz, Proceedings of the Fourth International Conference on Molten Aluminum Processing, Orlando, FL, USA, 1995, pp. 459–480. [21] L. Wang, S. Shivkumar, Z. Metallkde 86 (1995) 441–445. [22] K. Nogita, S.D. McDonald, J.W. Zindel, A.K. Dahle, Mater. Trans. JIM 42 (2001) 1981–1986.
Eutectic modification and microstructure development in Al–Si Alloys A.K. Dahle a,b,∗ , K. Nogita b , S.D. McDonald b , C. Dinnis a,b , L. Lu c a
CRC for Cast Metals Manufacturing, The University of Queensland, Brisbane, Qld 4072, Australia b Materials Engineering, The University of Queensland, Brisbane, Qld 4072, Australia c CSIRO Minerals, P.O. Box 883, Kenmore, Qld 4069, Australia Received in revised form 29 June 2005
Abstract Recent increasing applications for cast Al–Si alloys are particularly driven by the need for lightweighting components in the automotive sector. To improve mechanical properties, elements such as strontium, sodium and antimony can be added to modify the eutectic silicon from coarse and plate-like to fine and fibrous morphology. It is only recently being noticed that the morphological transformation resulting from eutectic modification is also accompanied by other, equally significant, but often unexpected changes. These changes can include a 10-fold increase in the eutectic grain size, redistribution of low-melting point phases and porosity as well as surface finish, consequently leading to variations in casting quality. This paper shows the state-of-the-art in understanding the mechanism of eutectic nucleation and growth in Al–Si alloys, inspecting samples, both quenched and uninterrupted, on the macro, micro and nano-scale. It shows that significant variations in eutectic nucleation and growth dynamics occur in Al–Si alloys as a function of the type and amount of modifier elements added. The key role of AlP particles in nucleating silicon is demonstrated. © 2005 Elsevier B.V. All rights reserved. Keywords: Al–Si eutectic; Modification; Eutectic nucleation; Eutectic grain size
1. Introduction Eutectic modification is a common process performed in aluminium–silicon based foundry alloys primarily to improve mechanical properties, particularly tensile elongation, by promoting a structural refinement of the inherently brittle eutectic silicon phase. It is well known that trace additions of strontium (a few hundred parts per million) to hypoeutectic aluminium–silicon alloys result in a transformation of the eutectic silicon morphology from a coarse plate-like structure to a well-refined fibrous structure [1]. Additions of antimony (albeit at higher concentrations of a few thousand parts per million) also give rise to a transformation of the eutectic silicon phase, however the result is not as dramatic with the coarse plates being refined rather than transformed to a fibrous morphology. Typical examples of the microstructure of unmodified, Sr-modified and Sb-modified alloys are shown in Fig. 1. Other elements including sodium and several rare earth elements are also known to result in varying degrees of modification [2].
∗
Corresponding author. Tel.: +61 7 3365 4350; fax: +61 7 3365 3888. E-mail address: [email protected] (A.K. Dahle).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.09.055
It appears to be universally accepted that the change in silicon morphology that occurs with modification is desirable. Unfortunately, modification has also been associated with negative side effects such as porosity, hot tearing and poor surface finish and is not always a recommended procedure. In recent years, it has become apparent that these complications occur because the change in silicon morphology does not occur in isolation, but is accompanied by a number of other significant changes in solidification and microstructure development. This paper compares the development of the cast structure on both macroscopic and microscopic scales in unmodified and modified alloys (both strontium and antimony). The interrelationship between nucleation and growth is demonstrated. 2. Experimental methods Several different experimental techniques have been employed throughout the research [3–5] and these are only briefly outlined in this paper. The base metal components, usually commercial purity Al and Si, were melted in an electrical resistance furnace and kept at about 720 ◦ C. A nominal composition of Al–10 wt.% Si was normally used because it contains a large volume fraction of eutectic, facilitating macroscopic inves-
244
A.K. Dahle et al. / Materials Science and Engineering A 413–414 (2005) 243–248
Fig. 1. Comparison of the silicon morphology in (a) unmodified, (b) Sr-modified (300 ppm Sr) and (c) Sb-modified (2400 ppm Sb), hypoeutectic aluminium–silicon alloys.
tigation, while still being a hypoeutectic alloy. Following stirring and surface skimming, the elements to be investigated were added. These included Sr as AlSr10 master alloy rods and Sb as AlSb10 ingots. Analysis of cooling curves was conducted according to the method described in detail elsewhere [3,4]. Additions in the range of 0–1000 ppm Sr and 0–5000 ppm Sb were examined. Samples were also quenched at different stages during the eutectic arrest. For these experiments, small tapered stainlesssteel cups coated with a thin layer of boron nitride were used to sample the melt. Two samples were taken in parallel by submerging the cups into the melt. The cups were placed on an insulated base, a thermocouple was positioned in the centre of one cup only, and insulation boards were placed on top of both samples. The samples were allowed to cool in air resulting in a cooling rate in the liquid just prior to nucleation of the primary aluminium of approximately 1.7 K/s, and a total solidification time of approximately 300 s. During solidification the cooling curve from the thermocouple was monitored on a real-time display, and the sample without a thermocouple was quenched approximately 50% of the way through the eutectic reaction. Although it is acknowledged that solidification will not proceed identically in both cups, this method prevents the thermocouple from interfering with solidification, or from damaging the microstructure during quenching. Samples were analyzed by optical and scanning electron microscopy including electron back-scattered diffraction (EBSD). To obtain three-dimensional images of the microstructures, selected quenched samples of unmodified and Sr-modified alloys were then subjected to serial sectioning. The unmodified and strontium-modified samples chosen for serial sectioning were quenched 45 and 12 s after eutectic nucleation, respectively. Serial sectioning involves the gradual removal of parallel layers of material and the imaging of each layer. Digital images of each section are “stacked” by using appropriate software to reconstruct the 3D microstructure. A full description of the procedure for producing 3D images is contained in Ref. [5]. Selected eutectic grains of unmodified samples with visible internal particles (potential nuclei) were chosen for further analysis using focused ion beam milling (FIB) followed by transmission electron microscopy (TEM). It is only recently that the
advanced techniques of FIB milling combined with in situ sample micro-manipulation have become available to selectively extract a TEM sample from any desired region within a larger sample. The FIB sample preparation was performed with a 30 kV Ga liquid metal ion source. TEM observations, micro-selected area diffractions (SAD) taken from the area of about 60 nm in diameter, small probe convergent beam electron diffraction patterns (CBED) taken from an area of ∼10 nm in diameter and EDX analysis were conducted using a 200 kV FE-TEM, see Ref. [6] for further details. 3. Results Typical cooling curves in the vicinity of the eutectic reaction are shown for the alloys used in this research in Fig. 2. Antimony additions resulted in a depression of the eutectic nucleation and growth temperatures and an increase in the amount of recalescence prior to growth. Similar but larger effects were observed with strontium modification. Fig. 3 summarises the defining micro- and macroscopic characteristics of unmodified, strontium-modified and antimonymodified alloys. Twinning in silicon fibres was examined using transmission electron microscopy for unmodified and strontiummodified alloys and results from the literature [7] for Sb have been included in Fig. 3a for comparison. In all alloys, co-zonal twinning is present in the silicon fibres with twins lying paral-
Fig. 2. Comparison of cooling curves from unmodified, Sb-modified (2600 ppm) and Sr-modified (200 ppm), Al–10% Si samples. The nucleation temperature (Tn ), minimum temperature prior to recalescence (Tmin ) and the growth temperature (Tg ) are included in the figure.
A.K. Dahle et al. / Materials Science and Engineering A 413–414 (2005) 243–248
245
Fig. 3. Comparison of unmodified, Sr-modified and Sb-modified structures. (a) TEM images of silicon crystals (note Sb-modified micrograph reproduced from a directionally solidified alloy in other research [7]), (b) EBSD maps illustrating the orientation of the eutectic aluminium relative to the primary aluminium dendrites, (c) 3D reconstruction of eutectic grains derived from serial sectioning of quenched samples (the unmodified sample is 85 m deep, the modified is 117 m deep) (d) optical micrographs of quenched samples (refined eutectic is quenched liquid) and (e) macrographs of quenched samples. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of the article.)
246
A.K. Dahle et al. / Materials Science and Engineering A 413–414 (2005) 243–248
lel to the apparent growth direction of the fibres. As shown in Fig. 3a, it was possible to find silicon crystals that were free of twinning in the unmodified alloys. Electron back-scattered diffraction was used to analyse the orientation of eutectic aluminium relative to that of the primary aluminium and typical results are shown in Fig. 3b. In unmodified alloys the vast majority of eutectic aluminium has an orientation identical to that of the surrounding dendrites, while in Sr-modified and Sb-modified alloys the eutectic aluminium has multiple orientations unrelated to the surrounding dendrites. These results indicate that the growth relation between the primary aluminium and the eutectic aluminium is epitaxial in the unmodified commercial alloys, while independent nucleation and growth occurs in the Sr- and Sb-modified alloys. There was a dramatic difference in the size of the eutectic grains. The reconstructed region of the unmodified alloy contained five separate eutectic grains (blue and red in Fig. 3c, unmodified). Of the two largest grains in the unmodified alloy; one was completely traversed by the series of sections while the other extended further into the sample. On the other hand, the Srmodified eutectic grain was very large and it appears as though only a small segment of this larger spherical grain has been reconstructed. This grain has been estimated to be approximately 500 m in diameter [5]. The modified eutectic silicon was too fine to reconstruct separately from the eutectic aluminium, so the two phases are represented as a single feature (red in Fig. 3c, Sr-modified). Fig. 3d shows optical micrographs of eutectic grains in samples that have been quenched at similar stages early during eutectic solidification. The eutectic grains in the unmodified alloy are at most a few hundred microns in diameter and consist of relatively few coarse silicon plates and a loosely coupled growth interface leading to grains that are often anisotropic in appearance. In the Sr-modified alloys, the eutectic grains are roughly circular in cross section and are typically much larger than the unmodified alloys (up to an order of magnitude). Each eutectic grain in the Sr-modified alloys contains a high density of silicon fibres. In the Sb-modified alloys, the eutectic grains are of an intermediate size and morphology, being more spherical than the unmodified grains. Each grain in the Sb-modified alloy contains numerous silicon plates, which are refined and more closely aligned than those found in the unmodified alloy. The difference in eutectic grain size between the unmodified, Sr-modified and Sb-modified alloys is further apparent in the macrographs of samples quenched early during eutectic solidification (Fig. 3e). In the unmodified sample, the eutectic grain size is so small that no grains are resolvable on the macrograph. In the Sr-modified alloy, there is a layer of coalesced eutectic grains lining the wall of the sample and several large grains can be seen independently distributed throughout the centre of the casting. In the Sb-modified alloy, the size and number of grains is intermediate to that found in the unmodified and strontiummodified alloys. In unmodified alloys, the eutectic grains were often observed to radiate from centrally located polyhedral silicon crystals, as seen in Fig. 3d. Occasionally, these polyhedral silicon crystals also have centrally located particles, which when analyzed using
Fig. 4. TEM image (top of micrograph is the exposed surface of the conventionally polished sample) and corresponding electron diffraction patterns from the Al–P–O precipitate on polished surface showing amorphous structure and small probe CBED patterns showing the AlP [0 1 1] pole and the silicon [0 1 1] pole. CBED patterns were obtained using an identical tilt angle.
conventional SEM techniques were found to contain Al, P and O [8]. In the current research, focused ion beam milling was used to selectively extract samples suitable for TEM analysis that contained both the central silicon crystal and its assumed nucleus. A typical TEM sample produced by this method along with SAD and CBED patterns is shown in Fig. 4. At some distance from the exposed conventionally prepared surface of the sample, no oxygen is present and the nucleus chemistry is consistent with AlP. Further to this, FEG-TEM analysis shows that there is no lattice mismatch between the (1 1 1) lattice planes of AlP and the Si crystal as shown in Fig. 5. Equivalent analysis on modified samples is challenging as it is difficult to locate possible nucleant particles during optical microscopy. Such difficulties are expected as statistically, the chance of sectioning through the exact center of a eutectic grain decreases dramatically as the nucleation frequency of eutectic grains decreases and their corresponding size increases.
Fig. 5. A high resolution FEG-TEM micrograph of the AlP/Si interface along with selected area diffraction patterns from regions A (Si) and B (AlP). There is no lattice mismatch between the (1 1 1) planes of AlP and Si.
A.K. Dahle et al. / Materials Science and Engineering A 413–414 (2005) 243–248
4. Discussion Contrary to expectations from the literature [9–11] it was found in this research that a higher degree of modification is not always associated with an increased density of non-cozonal twinning [12] and in fact this type of twinning was rarely observed in any of the alloys examined. This casts some doubt on the well accepted theory that the fibrous nature of the silicon in modified alloys is due solely to a high density of impurity induced twins imparting near isotropic growth properties to the silicon phase. It is possible that different growth mechanisms may operate or dominate in directionally solidified alloys, which have formed the basis for much of the pioneering TEM research into the growth mechanisms of eutectic silicon. While the growth mechanism of individual eutectic silicon fibres is likely to be important with respect to the fundamental modification mechanism, it is likely to be less critical for the processing characteristics of the solidifying melt. The difference in the nucleation frequency of eutectic grains between unmodified and modified alloys is significant. In unmodified alloys a large number of eutectic grains nucleate and these grow to only a small size prior to coalescence. In contrast, in Sr-modified alloys only a few grains nucleate and growth occurs over a significant distance prior to the coalescence of adjoining grains. In Sb-modified alloys, the grain size is intermediate between that of unmodified and Sr-modified alloys. From Fig. 3 it appears that there is a relationship between the degree of modification and the nucleation frequency of eutectic grains. The relationship appears to be that increasing the number of eutectic grains that nucleate results in a coarsening in the eutectic silicon morphology. This poses the question: Can the silicon morphology in modified alloys be simply related to the number of eutectic grains that nucleate? Certainly, as emphasised by Flood and Hunt [13] for a given rate of heat extraction the factor controlling the growth velocity of the eutectic is the solid–liquid interface area, which is directly controlled by the nucleation frequency. For a constant rate of heat extraction the growth velocity will vary inversely with the total solid–liquid surface area of the system. Therefore, the more grains that nucleate, the larger the solid–liquid surface area and the lower the growth velocity. However, although modification will cause an increase in the growth velocity during eutectic solidification, it has been shown that this increase in velocity is not sufficiently large to cause a flake-fibre transition in the eutectic silicon [8]. At this stage it must be conceded that the fundamental mechanism responsible for silicon modification is still unclear. The nucleation differences that occur between unmodified alloys and those modified with Sr or Sb are supported by the various analysis methods used in this research. Firstly, the eutectic nucleation and growth temperatures decrease when comparing unmodified, Sb-modified and Sr-modified alloys, respectively (Fig. 2). This implies an increasing degree of nucleation difficulty with Sb and Sr additions and in the case of these alloys, this seems to correspond well with the fact that fewer grains are present in the macrographs (Fig. 3e). The differences in the relative crystallographic orientations of eutectic and primary aluminium also support a difference in nucleation mechanism.
247
In the unmodified alloys it is proposed that the eutectic grains nucleate in the vicinity of aluminium dendrites and grow only a small distance, allowing the eutectic aluminium to adopt and maintain the same crystallographic orientation as the adjoining dendrite. In the modified alloys, the eutectic grains either nucleate independently of the aluminium dendrites or the large growth distance required of the larger grain, necessitates repeated nucleation during growth, giving rise to multiple crystallographic orientations of the aluminium within each eutectic grain. The difference in size of the eutectic grains is further confirmed with the results of the three dimensional reconstruction of the serial sectioning data (Fig. 3c). Despite a scarcity of research, it is likely that the size and distribution of eutectic grains in aluminium–silicon alloys will influence a variety of in-service and processing properties. These include surface appearance, corrosion resistance, mechanical properties and susceptibility to casting defects such as porosity and hot-tearing. This is because the distribution of the last few percent of liquid (important for porosity and hot-tear formation) and the location of ternary elements that precipitate during or after eutectic Al–Si solidification depends on the number and size of the eutectic grains. In hypoeutectic alloys, there is a significant proportion of aluminium dendrites present in the microstructure prior to eutectic solidification. It follows that the nucleation problem for the commencement of the eutectic reaction lies with the nucleation of the silicon phase. In the unmodified alloy, the result that AlP is the nuclei for eutectic silicon (Figs. 4 and 5) is not surprising, as both AlP and silicon have cubic structures with near identical ˚ [14] and 5.431 A ˚ [15], respeclattice parameters of 5.421 A tively. AlP was also identified as a good nucleant for eutectic silicon in the pioneering work on eutectic nucleation in Al–Si alloys by Crosley and Mondolfo [16]. Other nuclei have also been reported for unmodified silicon crystals and these include arsenic, sulphur, selenium and tellurium [17,18] and other possibilities include iron-containing phases and oxides introduced during processing. The decrease in nucleation frequency with strontium and antimony additions suggests that strontium somehow removes, or neutralizes, nuclei that are present in the unmodified alloy. There is some evidence to suggest that the poisoning mechanisms involves the nuclei (of which AlP appears to be the most prolific) being saturated by strontium containing intermetallics, most likely Al2 Si2 Sr which are frequently present in strontiummodified commercial aluminium–silicon alloys [19–21]. Similarly, Sb based intermetallics may also be important in removing phosphorus from the melt prior to eutectic solidification in Sbmodified alloys. At the extremity, modification of Al–Si alloys with Na has been shown to inhibit bulk eutectic nucleation and instead promote growth of a massive eutectic interface opposite to the thermal gradient [22]. 5. Conclusions From the current research it must be concluded that the fundamental mechanism of eutectic modification in Al–Si alloys (change in silicon morphology) is still not fully understood.
248
A.K. Dahle et al. / Materials Science and Engineering A 413–414 (2005) 243–248
Mechanisms based on the introduction of high densities of noncozonal twins may be more relevant to directionally solidified alloys than samples solidified under conditions of equiaxed growth. There are distinct differences in eutectic nucleation between unmodified and modified alloys and these are reflected in cooling curves, EBSD and micro- and macroscopic investigations. In the alloys used in this research there was increasing nucleation difficulties after individual additions of Sr or Sb to an unmodified alloy. This was reflected in the eutectic grain sizes with the eutectic grains being largest in the Sr-modified alloys and of an intermediate size in Sb-modified alloys. The differences in nucleation patterns are likely to be responsible for many of the complications commonly associated with modification. The compound AlP is a common nucleus for eutectic Si in unmodified alloys, but appears to be less active in modified alloys. The mechanism for this transition is unclear but may relate to scavenging of phosphorus by intermetallics introduced with the modification process. Although the eutectic grain size and silicon morphology are linked, they are not completely dependent. References [1] S. Lu, A. Hellawell, Metall. Trans. A 18 (1987) 1721–1733. [2] K. Nogita, S.D. McDonald, A.K. Dahle, Mater. Trans. 45 (2004) 323–326.
[3] K. Nogita, A. Knuutinen, S.D. McDonald, A.K. Dahle, J. Light Met. 1 (2001) 219–228. [4] A.K. Dahle, K. Nogita, J.W. Zindel, S.D. McDonald, L.M. Hogan, Metall. Mater. Trans. A 32A (2001) 949–959. [5] C.M. Dinnis, A.K. Dahle, J.A. Taylor, Mater. Sci. Eng. A 392 (2005) 440–448. [6] K. Nogita, S.D. McDonald, K. Tsujimoto, K. Yasuda, A.K. Dahle, J. Electron Microsc. 53 (2004) 361–369. [7] S. Khan, R. Elliott, J. Mater. Sci. 29 (1994) 736–741. [8] S.D. McDonald, A.K. Dahle, J.A. Taylor, D.H. St. John, Metall. Mater. Trans. A 35A (2004) 1829–1837. [9] M. Shamsuzzoha, L.M. Hogan, Philos. Mag. A 54 (1986) 459–477. [10] S. Lu, A. Hellawell, J. Cryst. Growth 73 (1985) 316–328. [11] M. Shamsuzzoha, L.M. Hogan, J. Mater. Sci. 24 (1989) 2849–2859. [12] K. Nogita, J. Drennan, A.K. Dahle, Mater. Trans. 44 (2003) 625– 695. [13] S.C. Flood, J.D. Hunt, Met. Sci. 15 (1981) 287–294. [14] M.E. Straumanis, E.Z. Aka, J. Appl. Phys. 23 (1952) 330. [15] C. Yeh, Z.W. Lu, S. Froyen, A. Zunger, Phys. Rev. B 46 (1992) 10086–10095. [16] P.B. Crosley, L.F. Mondolfo, AFS Trans. 74 (1966) 53–64. [17] L.F. Mondolfo, Aluminium Alloys: Structure and Properties, Butterworths, London-Boston, 1976. [18] T.B. Abbott, B.A. Parker, Cast Met. 1 (1988) 122–132. [19] G. Chai, L. B¨ackerud, AFS Trans. 100 (1992) 847–854. [20] G. Laslaz, Proceedings of the Fourth International Conference on Molten Aluminum Processing, Orlando, FL, USA, 1995, pp. 459–480. [21] L. Wang, S. Shivkumar, Z. Metallkde 86 (1995) 441–445. [22] K. Nogita, S.D. McDonald, J.W. Zindel, A.K. Dahle, Mater. Trans. JIM 42 (2001) 1981–1986.