Solidification of Cast Magnesium Alloys

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2Monash University, School of Physics & Materials Engineering, Melbourne ... Each year magnesium alloys are finding new applications in the automotive industry despite the fact that there is continual pressure to reduce the cost of producing automotive components ..... Dynamics (Academic Press, New York, 1973), 407.
Magnesium Technology 2003 Edited by Howard I. Kaplan TMS (The Minerals, Metals & Materials Society), 2003

SOLIDIFICATION OF CAST MAGNESIUM ALLOYS D.H. StJohn1, A.K. Dahle1, T. Abbott2, M.D. Nave3, and Ma Qian1 CRC for Cast Metals Manufacturing (CAST), UDP No. 55, University of Queensland, 4072, Australia 1 University of Queensland, Division of Materials Engineering, Brisbane Qld 4072, Australia 2 Monash University, School of Physics & Materials Engineering, Melbourne Vic 3168, Australia 3 Deakin University, School of Engineering & Technology, Geelong Vic 3217, Australia The key solidification processes that control the final microstructure are the initial nucleation events, the growth of these nuclei into primary dendrites and finally eutectic solidification. The slow cooling rates encountered in sand casting tend to produce coarse microstructures that have low ductility. To improve ductility the grain size needs to be decreased by the addition of grain refiner. Alternatively, the component can be heat treated which increases ductility but also lowers the yield strength. In contrast, HPDC creates a very dynamic environment for solidification and solidification begins in the shot sleeve even before the melt enters the cavity. The mush enters the die cavity at high speed and solidification progresses to high solid fraction while the mush continues to fill the cavity. This can result in banded defects that may be detrimental to mechanical properties. Thus solidification processes need to be considered in the light of the casting process being used. This paper briefly describes the response of magnesium alloys to the key elements of the solidification process and also considers solidification in the dynamic environment encountered in HPDC and to some extent in Thixoforming and squeeze casting.

Abstract A description of the key solidification steps in the formation of the as-cast microstructure of magnesium alloys is presented. The focus is on the two common magnesium alloy groups: Mg-Al alloys and Mg-Zn-rare earth alloys. The key elements described are: nucleation (including grain refinement), growth of the primary phase and the formation of the eutectic phases. In addition the effect of casting process (e.g. high-pressure diecasting and sand casting) on the outcomes from solidification are discussed. This includes consideration of the formation of banded defects during solidification in the dynamic environment of highpressure die casting. Introduction Each year magnesium alloys are finding new applications in the automotive industry despite the fact that there is continual pressure to reduce the cost of producing automotive components [1]. Magnesium has a number of advantages particularly in allowing parts consolidation which results in a lower overall cost. However, the need to achieve consistent properties is becoming paramount as applications move to components requiring a higher level of structural integrity. Therefore, a better understanding of the solidification process during the casting of components is required to ensure high quality and low cost [2]. Casting is still the dominant process for producing magnesium components. By far the most popular process is high pressure die casting (HPDC) which is used for large thin section castings such as instrument panels and seat frames. Thixoforming is used for small components such as mobile phone and computer cases and development work is underway to extend this process to larger automotive components. Sand casting has been the main casting method for large aerospace castings such as helicopter gear box housings.

Nucleation and Grain Refinement In HPDC nucleation is not usually an issue as the high chill of the mould materials ensures many small grains of primary magnesium phase are formed and these are carried into the melt to generate a fine microstructure. In sand casting, on the other hand, grain refinement is required to provide uniform properties and a high yield strength. Mg-Al alloys are difficult to adequately grain refine as there is not a simple and reliable grain refining method. Attempts to find a suitable refining technology are the subject of ongoing research and a review of these methods was recently published by Lee et al. [3]. This paper will focus on the nonaluminium containing alloys where zirconium is a very effective grain refiner. Zirconium (Zr) is a potent grain refiner for magnesium alloys that contain little (impurity level) or no Al, Mn, Si, Fe, Ni, Co, Sn and Sb (zirconium forms stable compounds with these elements) [4]. When added to these magnesium alloys, where the maximum solubility of Zr in molten magnesium is approximately 0.6%, the use of Zr can readily cause the average grain size to decrease by 80% or more for normal cooling rates. This exceptional grainrefining ability makes Zr an important alloying element for magnesium alloys containing zinc, rare earths, thorium, calcium, or a combination of these elements.

The properties achieved depend to a large extent on the alloy used and the resulting microstructure achieved. The scale of the microstructure can vary by a large magnitude due to this very wide range of casting methods. There are two main types of magnesium alloys. Magnesium – aluminium alloys (e.g. AZ91, AM60 and AM50) are predominantly used for HPDC and thixoforming of thin section castings of electronic components and structural automotive castings such as instrument panels and steering wheels. The other group of alloys is magnesium - zinc alloys (e.g. ZE41, ZK60) that can be alloyed with zirconium to provide a fine grain size and rare earth elements to provide creep resistance. Applications include engine components and extrusion billet.

The most characteristic feature of the microstructure of a magnesium alloy containing more than a few tenths per cent soluble zirconium is the Zr-rich cores that exist in most magnesium grains [4], particularly when observed using a scanning electron microscope (SEM) in the backscattered (BES) 95

mode (See Fig. 1a). It is generally accepted that solidification of Mg-Zr alloys containing Zr close to 0.6% occurs by a peritectic mechanism [4,5] and the Zr-rich cores are products of this peritectic reaction. Observations of the Zr-rich cores on polished sections cut at different angles from a Mg-0.56Zr alloy cone sample (including transverse, longitudinal and a 45q diagonal) demonstrated that these Zr-rich cores generally appear with a nearly spherical or elliptical form and many contain a tiny particle in their central regions. Electron microprobe analyses of a total of 50 such tiny particles observed in the Mg-0.56Zr alloy showed that the majority were almost pure Zr particles. On the other hand, examination of the distribution of Zr in areas that surround the Zrrich cores showed little presence of Zr in most traversed areas, indicating that almost all of the Zr in solid solution is concentrated in the Zr-rich cores. The distribution of Zr in these cores was found to be inhomogeneous and vary in a wide range, e.g., from 0.5% to 3% (the maximum solubility of Zr in solid magnesium is ~ 3.8%). The Zr-rich cores observed in different Mg-Zr alloys are usually less than 20 Pm in diameter for normal cooling rates and the tiny particles inside are usually less than 2 Pm.

rich cores can exist anywhere inside a grain rather than just in the centre of the grain. These are an important microstructural feature that must be understood in order to fully understand the solidification of Mg-Zr alloys. a

b

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Figure 2: Grain refinement with zirconium: (a) pure magnesium and (b) after the addition of 1% Zr at 730qC. Micrographs are of the same magnification.

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Grain refinement of magnesium alloys by Zr is noticeable at very low levels of soluble Zr (~0.04%) [3] and, in general, improves with increasing soluble zirconium towards the maximum soluble Zr content (~0.6%). According to the early work by Sauerwald [4], only Zr that is dissolved in the liquid magnesium at the time of pouring is effective in grain refinement. Insoluble Zr such as undissolved Zr particles was believed to be irrelevant to grain refinement, despite both D-zirconium and magnesium having the same type of crystal structure and nearly identical lattice parameters. This view appeared to be widely accepted until it was recently shown by Tamura et al. [6] that undissolved Zr particles actually play an important role in the grain refinement of magnesium alloys. More recent work [7] has confirmed the beneficial role of undissolved Zr particles in grain refinement and showed that grain refinement of magnesium alloys by Zr is dictated by both soluble and insoluble zirconium. As such, an ideal Zr alloying process should end up with both high soluble and high total Zr at the time of pouring in order to achieve the best grain refinement. Figure 2 shows an example of grain refinement of pure magnesium with 1% Zr addition

Figure 1: a) A typical view of the zirconium-rich cores in a Mg0.56Zr alloy (backscatter SEM), and b) SEM backscatter image showing that each magnesium grain may contain a number of Zrrich cores (Mg-0.56Zr alloy). Despite the general agreement on a peritectic mechanism for the formation of the characteristic Zr-rich cores, it is still far from clear exactly how the subsequent grain growth occurs once a Zrrich core has formed. It has been found that many magnesium grains in a Zr-containing Mg alloy contain a number of Zr-rich cores rather than a single core, as shown in Fig. 1b, and these Zr-

The above findings confirm that the mechanism of grain refinement in magnesium alloys is identical to that in aluminium alloys where both potent nucleant particles and a strongly segregating solute element must be present [8]. Considering the 96

work on grain refinement of aluminium alloys it is clear that the nucleant particles present in Mg-Al alloys are being poisoned and thus it is difficult to obtain sufficient refinement unless this poisoning effect is overcome. Understanding the poisoning mechanism is a key to developing a suitable grain refiner for MgAl alloys.

Eutectic Growth Magnesium forms eutectic systems with a wide range of alloying elements [9] and the microstructures of the common casting alloys, AZ91, AM50/60, AS21/41, AE42, as well as recently developed alloys such as ZA85, contain small amounts of eutectic [10, 11]. While alloys containing any proportion of eutectic may be produced [12], commercial casting alloys contain very low volume fractions of eutectic compared with cast irons and aluminium-silicon alloys. Pure magnesium is very weak and low to moderate levels of alloying provide solid-solution strengthening and strengthening from the presence of massive second-phase particles formed during eutectic solidification. Addition of some elements also allows precipitation hardening. While higher levels of alloying addition should improve castability by increasing the volume fraction of eutectic, the dendrite-eutectic interface provides an increasingly convenient crack path around the more ductile magnesium dendrites [13] resulting in a sharp decrease in the ductility of the alloys [9]. Large additions of heavy and/or expensive elements also detract from the economic viability of the alloys by significantly increasing their density and cost. For these reasons, addition levels are generally below equilibrium maximum solid-solubilities and it is only the non-equilibrium cooling in practical casting processes that causes eutectic formation in commercial alloys.

Dendritic Growth Figure 3 shows a typical microstructure of Mg-Al alloys. Because magnesium has a hexagonal close-packed structure, the dendrite arms branch with a six-fold symmetry. Well-developed primary D-Mg dendrites with secondary arms (A) showing sixfold symmetry are clearly visible in Figure 3. The eutectic of Mg17Al12 (B) and D-Mg solid solution (C) is located in the interdendritic regions. According to the Mg-Al equilibrium phase diagram, the eutectic phase (Mg17Al12) is expected to appear when the aluminium content reaches ~13 wt%. However, the eutectic phase appears in alloys containing as little as 2 wt% Al for nonequilibrium cooling conditions normally encountered in castings.

Figure 3: Micrograph of fully developed dendrites in a Mg-15 wt% Al alloy permanent mould casting. The magnesium dendrites have a characteristic sixfold symmetric shape (A). The white phase between the dendrites is secondary eutectic phase Mg17Al12 (B) and the dark regions between the dendrites are Al-rich Mg solid solution (C). The growth of magnesium dendrites occurs according to the usual principles of primary phase solidification. Solute elements with a distribution coefficient less than unity are rejected ahead of the dendrites and this can cause changes in the microstructure. For example, a small addition of aluminium to pure magnesium leads to a morphological change of the primary phase from a cellular to a dendritic structure. Rosette-like globular equiaxed grains form with Al-rich solid solution between the dendrite arms. As the aluminium content is increased further to 5 wt%, dendrites with pools of eutectic phase between the dendrite arms start to develop and, when the aluminium content is further increased, a fully developed dendritic structure with sharp tips is observed. The concept of growth restriction factor (GRF) seems to be applicable for magnesium alloys and the elements with the largest segregation potential, in decreasing order, are Zr, Ca and Si.

a

b

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Figure 4: (a) Lamellar, (b) fibrous, (c) partially divorced and (d) fully divorced morphologies in Mg-Al alloys of various compositions [12,14]. Several eutectic morphologies that may be produced in magnesium alloys are shown in Fig. 4. For alloy compositions and casting conditions used commercially, the eutectic morphology is generally fully or partially divorced. In the fully divorced morphology (Figure 4d), the “eutectic” Mg coats the primary dendrites and the intermetallic phase exists as massive particles or a thin layer between the coated dendrites. The partially divorced eutectic (Figure 4c) is similar, however not all the “eutectic” Mg coats the primary dendrites – a small amount solidifies as islands inside the intermetallic phase. Fully divorced morphologies are usually formed in die-casting, while partially divorced morphologies are more likely at lower cooling rates. In permanent mould cast Mg-Al [12] and Mg-Al-Zn [14] alloys, a 97

reduction in aluminium content from 15 wt% Al to 9 wt% Al also favours a more divorced morphology. This trend is expected to continue to lower aluminium contents, but is not always evident among AZ, AM and AS alloys due to the different ternary additions. It is worth noting that as the morphology of the eutectic is quite sensitive to small changes in alloy content and cooling conditions, different morphologies are sometimes observed in the same alloy cast under similar conditions and even in nearby regions of the same casting.

only small amounts of Mg17Al12 and larger amounts of the more stable Mg2Si in the grain boundary regions. Creep strength is further improved in AE42 [17], where the Mg-Mg17Al12 eutectic reaction is avoided completely and the Mg17Al12 is replaced with Al4RE and Mg12RE [18], respectively. Minor additions of Bi and Sb to AZ91 provide remarkable improvements in creep strength, which are partly due to the suppression of discontinuous precipitation and the presence of rod-shaped Mg3Bi2 and Mg3Sb2 particles that are stable at high temperatures and often straddle grain boundaries [16]. It should be noted, however, that grain boundary sliding is only one of a number of creep mechanisms active in magnesium alloys and the dominant mechanism varies with alloy composition and service conditions (e.g. Table 3 in [18]).

The general similarities in the eutectic morphologies of magnesium casting alloys and the variation of these morphologies with cooling rate and alloy content can be explained in the following manner [12,14]: The predominance of fully and partially divorced morphologies is primarily a consequence of the high volume fraction of primary dendrites, which restricts eutectic solidification to small regions between dendrite arms and adjacent dendrites. When the interdendritic regions are small, a relatively small undercooling for nucleation and growth of the intermetallic phase is sufficient to cause solidification of practically all of the interdendritic liquid to occur outside the coupled zone, producing fully or partially divorced morphologies. Furthermore, if a large number of regions of interdendritic liquid become isolated from one another, nucleation of the intermetallic phase is required in each of these regions and these regions are more likely to solidify with a fully divorced morphology, since the undercooling required for nucleation is larger than the undercooling required for growth. Consequently, fully divorced morphologies are more likely in alloys with low levels of alloying additions, since the interdendritic regions tend to be smaller and more isolated when “eutectic” solidification commences.

Dynamic Solidification and Defect Formation High-pressure die cast magnesium alloys display several characteristic macro- and microstructural features including: x Distinctive bi-modal grain size distribution x Skin x Defect bands In cold chamber machines it is common for some solidification to occur in the shot sleeve prior to the commencement of the shot. The solid phase formed in this way is able to grow considerably larger than the material solidified during die filling, giving rise to a bimodal distribution of primary magnesium grains. The grains formed in the shot sleeve are referred to as externally solidified grains. The amount of these grains varies depending upon the metal temperature and other factors, and can range from zero up to about 20% [19].

Generally, high cooling rates produce more highly branched dendrites and trap the eutectic liquid into smaller, more isolated spaces. High cooling rates also increase the average undercooling required for nucleation of the intermetallic phase and for both of these reasons, the eutectic is frequently more divorced at higher cooling rates [10,12,14]. Although higher cooling rates are expected to produce more coring during dendrite growth and therefore a larger volume fraction of eutectic, which should produce less divorced eutectic morphologies, in most cases this effect appears to be counteracted by the mechanism just described. Since eutectic morphology is highly sensitive to the morphology of the primary dendrites [12,14], grain refinement levels and columnar-to-equiaxed transitions in the primary phase may also affect the morphology of the eutectic.

Another common feature of pressure die cast magnesium is the presence of a surface layer or skin. The skin is characterised by a higher volume fraction of second phases, Mg17Al12 in the case of Mg-Al alloys and, often, a finer microstructure [20]. The formation of the skin suggests the segregation of solute elements to the surface that may occur due to the semi-solid state of the metal as it flows in the cavity. The depletion of solid phase adjacent to surfaces is a common phenomenon for flowing suspensions [21]. Alternatively the higher Mg17Al12 content may be a consequence of a faster cooling rate near the surface without enrichment in solute. Possibly the most visually apparent microstructural feature of magnesium pressure die castings are bands of defects which, in most cases, form parallel to the casting surface. The make up of these bands seems to vary considerably and may consist of cracks, porosity (shrinkage and/or gas porosity) and/or segregation. Examples of defect bands are shown in Fig. 5. The bands may be single or multiple and can sometimes intersect with the surface. According to a mechanism proposed by Dahle and StJohn [22] the solid phase, within the semi-solid mixture flowing in the die cavity, migrates away from the walls and towards the centre. While this effect raises the solid fraction towards the centre, heat extraction through the die results in increased solid fraction near the surface. As a consequence there is an intermediate zone with relatively high liquid fraction. As solidification and filling proceeds, the centre and surface regions begin to develop some mechanical strength resulting in a localisation of shearing in the high liquid fraction region. This mechanism should result in the central region having a lower solute content than the outer region,

The effect of ternary elements on eutectic morphology is quite complex. Different elements may affect eutectic growth in very different ways. It has been established that the addition of small amounts of zinc to Mg-Al alloys increases the probability of forming a fully divorced eutectic [14]. There are several possible reasons for this effect, but the most likely reasons are the promotion of a more highly-branched dendrite structure (primarily due to the low partition coefficient of Zn in Mg) and changes in the equilibrium and non-equilibrium phase boundaries [14]. Besides the detrimental effect of the divorced Mg-Mg17Al12 eutectic on the ductility of Mg-Al alloys, the presence of Mg17Al12 particles along the grain boundaries and discontinuous precipitation of Mg17Al12 during high temperature service may contribute to their low creep strength by assisting intergranular cavitation and grain boundary sliding [15,16]. Compared with AZ and AM alloys, creep strength is higher in AS alloys, which have 98

which has been reported to be the case in at least one instance [19].

microstructural examination, that in the runner and in sections of the casting where flow is uni-directional, flow occurs through a cylindrical section much smaller than the physical cross-section. The size of the cylindrical section appeared to be determined by the flow velocity, which typically fell in the range of 140 165m/s. It was claimed that solidification initially occurred in the runner, reducing the effective cross section and increasing the velocity. Once a velocity of about 150m/s was attained, viscous heating prevented further solidification. It was claimed that because viscosity rises rapidly for solid percentages above 50%, the viscous heating process would stabilise the solid fraction of material entering the cavity. As this phenomenon was also observed in sections of the cavity where flow was uni-directional, this suggests that the bands of defects occur where solidification was temporarily arrested by viscous heating. Rodrigo and Ahuja [24] also claim that defect bands are more defined in regions where flow is largely unidirectional.

a

The observation of banded defects is far more common in magnesium alloys than in aluminium alloys. Also, HPDC of magnesium alloys is able to fill long thin sections while aluminium is not. The reason for this is not clear. However, the two behaviours may be linked. A possible explanation can be proposed by considering the differences in solidification behaviour between aluminium and magnesium alloys. AZ91 has a freezing range of 160oC which is five times larger than Al7wt%Si (e.g. A356) at about 30oC. Also, the volume fraction of eutectic in AZ91 is about 25% compared with about 50% for A356. By considering this and assuming thermal properties are similar, it can be envisaged that AZ91 will flow significantly longer distances before the semisolid microstructure locks up with eutectic solidification. This extra time allows for rearrangement of the primary magnesium grains and the formation of a liquid band to lubricate the flow of the semisolid mush into the die cavity. This segregated solute enriched band allows flow to continue until the temperature drops to near the eutectic temperature. Hence the bands are clearly visible in the microstructure.

b

Figure 5: Example of defect bands (a) in a 2mm thick cast tensile bar, (b) in a commercial casting (3mm thick).

Defect bands occur in a range of forms, such as cracks, bands of porosity (gas and/or shrinkage) and segregation. These all have superficial similarity, however it is not clear whether they form by the same process or whether multiple mechanisms are involved. Further investigations are necessary before their formation mechanism is fully understood.

Mao [23] observed bands of pores that appeared spherical suggesting they were due to gas porosity. The porosity was thought to have formed due to a sudden drop in pressure, possibly as a result of a small back movement of the plunger tip during intensification. This mechanism would appear only to explain instances of gas porosity and not shrinkage porosity or segregation.

Summary

Rodrigo and Ahuja [24] investigated the effects of casting parameters on the formation of defect bands. Although they did not propose a formation mechanism they did assess the validity of the above mechanisms in light of their experimental findings. The model proposed by Dahle and StJohn [22] was disputed by Rodrigo and Ahuja [24] on the basis that the presence or absence of externally solidified grains did not influence the appearance of the pore bands and that when externally solidified grains were present they often appeared on either side of the defect bands. However, Dahle and StJohn’s theory [19,22,25] also applies to the situation where externally solidified grains are not present so it is not clear whether this is in fact contradictory evidence. The theory proposed by Mao [20] was also disputed by Rodrigo and Ahuja [24] as the observed porosity was due to shrinkage rather than gas.

The solidification of magnesium alloys differs in a number of essential ways from the solidification of aluminium foundry alloys. Magnesium alloys have a much lower volume fraction of eutectic and the composition of most alloys is within the range of the maximum solid solubility of the major alloying elements such as aluminium and zinc. This means that commercial alloy microstructures usually have divorced eutectics that form due to non-equilibrium solidification. This can be a problem for structural alloys as it decreases ductility, but the high solidification rates in HPDC or heat treatment result in improved ductility. On the other hand, divorced eutectic formation is likely to be an advantage for creep resistant alloys as the stable intermetallic phase present in the grain boundaries increases the microstructure’s resistance to deformation at elevated temperatures. Grain refinement of sand cast alloys also assists in improving the distribution of the intermetallic phase and this is

A patent application by Murray and Cope [26] suggests an alternative mechanism for the formation of these bands. Although they do not explicitly refer to bands, they claim, from 99

used to good effect in non-aluminium containing magnesium alloys by adding zirconium. Zirconium provides very potent nucleant particles and high segregating solute, both of which ensure very good grain refinement. Unfortunately a similarly effective grain refiner for aluminium containing alloys has not been developed and is the subject of ongoing research.

12. M.D. Nave, A.K. Dahle, and D.H. StJohn, "Eutectic Growth Morphologies in Magnesium-Aluminium Alloys," Magnesium Technology 2000, eds. H.I. Kaplan, J. Hryn and B. Clow, (TMS, 2000), 233-242. 13. S. Lee, S.H. Lee, and D.H. Kim, "Effect of Y, Sr and Nd Additions on the Microstructure and Microfracture Mechanism of Squeeze-Cast AZ91-X Magnesium Alloys," Met. Mat. Trans., 29A (1998), 1221-1235.

Another difference compared to aluminium alloys is the large freezing range. This may enhance the formation of banded defects in HPDC and thixocast alloys. These defects may reduce the ductility of the cast components but may also be responsible for the ability of magnesium alloys to fill long thin sections. A better understanding of the rheological behaviour of partially solidified alloys is required in order to reduce the negative effects of banded defects on mechanical performance while optimising the filling behaviour of magnesium alloys during HPDC.

14. M.D. Nave, A.K. Dahle, and D.H. StJohn, "The Role of Zinc in the Eutectic Solidification of Magnesium-AluminiumZinc Alloys," Magnesium Technology 2000, eds. H.I. Kaplan, J. Hryn and B. Clow, (TMS, 2000), 243-250. 15. M. Regev et al., "Creep Studies of Coarse-grained AZ91D Magnesium Castings," Mat. Sci. Eng., A252 (1998), 6-16.

Acknowledgements

16. Y. Guangyin, S. Yangshan, and D. Wenjiang, "Effects of Bismuth and Antimony Additions on the Microstructure and Mechanical Properties of AZ91 Magnesium Alloy," Mat. Sci. Eng., A308 (2001), 38-44.

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