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Agglomeration and bending of equiaxed crystals during solidification of hypoeutectic Al and Mg alloys
Available online at www.sciencedirect.com
Acta Materialia 58 (2010) 261–271 www.elsevier.com/locate/actamat
Agglomeration and bending of equiaxed crystals during solidification of hypoeutectic Al and Mg alloys S. Otarawanna a,*, C.M. Gourlay b, H.I. Laukli c, A.K. Dahle d a
CAST CRC, Materials Engineering, The University of Queensland, Brisbane, Qld 4072, Australia b Department of Materials, Imperial College London, London SW7 2AZ, UK c Hydro Aluminium, Research & Technology Development, N-6601 Sunndalsøra, Norway d ARC CoE for Design in Light Metals, Materials Engineering, The University of Queensland, Brisbane, Qld 4072, Australia Received 11 August 2009; received in revised form 31 August 2009; accepted 2 September 2009 Available online 1 October 2009
Abstract Agglomeration and bending of equiaxed crystals were studied by microstructural characterization of specimens produced by near-static cooling, high-pressure die casting and ThixomoldingÒ, where the solidifying crystals experience different levels of mechanical stresses. EBSD was used to acquire statistical grain misorientation data which is linked to crystal agglomeration and bending behavior during solidification. An aluminum alloy and two magnesium alloys were used to compare grain misorientations for different crystal structures. The length fraction of low-energy grain boundaries in HPDC and Thixomolded samples was substantially higher than in “statically cooled” samples. This is attributed to the high shear stresses and pressure applied on the solidifying alloy, which promote crystal collisions and agglomeration. In-grain misorientations were found to be significant only in branched dendritic crystals which were subjected to significant shear stresses. This is related to the increased bending moment acting on long, protruding dendrite arms compared to more compact crystal morphologies. Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Solidification microstructure; Electron backscattering diffraction (EBSD); Aluminum alloys; Grain boundary energy; Misorientation
1. Introduction Agglomeration and bending of primary crystals in the mushy zone during solidification are important phenomena influencing many aspects of casting processes and often affect the microstructure of the as-cast component. Both phenomena strongly influence the rheology of partially solid alloys. For example, agglomeration (and disagglomeration) of crystals is the origin of thixotropy in semi-solid metallic suspensions [1–4]. Agglomeration of crystals often entraps liquid, increasing the effective solid fraction and the apparent viscosity [1–4]. Bending of solidifying dendrite arms creates misorientation within crystals by dynamic recovery and can subsequently cause arm fragmentation *
Corresponding author. Tel.: +61 7 3365 1387; fax: +61 7 3365 3888. E-mail addresses: [email protected], [email protected] (S. Otarawanna).
[5–9], which reduces the grain size and the mush strength [2,4,10,11]. Agglomeration and bending are both strongly influenced by the application of external mechanical stresses during solidification [2,4] and these phenomena are therefore particularly important in casting processes where the alloy is forced into a mold during solidification. Two types of process where the solidifying alloy is severely deformed are high-pressure die casting (HPDC) and the semi-solid metal processing (SSM or SSP) variants, ranging from thixocasting to thixoforging. In the HPDC process, fully liquid alloy is transferred into the shot chamber where some uncontrolled solidification occurs, resulting in socalled externally solidified crystals (ESCs), which occupy up to 30 vol.% in the as-cast microstructure [12]. ESCs undergo intense shear during die filling where shear rates at the gate are typically 104 s1 [13]. Some further solidification occurs during die filling under strong forced
1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2009.09.002
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convection. Once the cavity is full, intensification pressure of up to 120 MPa [14] is applied to the casting and the solidifying alloy is forced deeper into the cavity to feed the casting. Since the shot-sleeve cooling rate (on the order of 10 °C s1 [15]) is significantly lower than the in-cavity cooling rate (in the range 102–103 °C s1 [16,17]), ESCs are larger than in-cavity solidified grains, often by an order of magnitude [18]. In SSM, a solid–liquid mixture with globular solid morphology and controlled solid fraction (fs) is produced prior to die filling. This mixture is often generated by strong forced convection at a known temperature in the freezing range [2,4]. The SSM variant examined in this paper is Thixomolding, also known as metal injection molding [19]. In Thixomolding, alloy chips are fed into a screw mixer and heated to a temperature in the freezing range while being intensely sheared. The primary globule fs is typically 10% or below [20], which is lower than is typical for SSM variants such as thixocasting and rheocasting (fs 50% [21]). The slurry is then injected rapidly into the die cavity, after which an after-pressure, analogous to intensification pressure in HPDC, is applied to assist feeding. In this work, electron backscatter diffraction (EBSD) is used to study agglomeration and bending of equiaxed crystals in HPDC, Thixomolding and on samples solidified without forced convection or an applied pressure. Additionally, an AlSi4MgMn aluminum alloy, and AM50 and AZ91 magnesium alloys, are used to compare grain misorientations for different crystal structures, i.e. face-centered cubic (fcc) in Al and hexagonal close-packed (hcp) in Mg. 2. Experimental Al alloy AlSi4MgMn and Mg alloys AM50 and AZ91, with the compositions shown in Table 1, were used in this work. Alloys were cast by HPDC, Thixomolding and a “static cooling” technique. HPDC samples were produced with 80 °C superheat. The AlSi4MgMn alloy was cast by the cold-chamber HPDC process into a tensile bar with 125-mm total length, 30-mm gage length and 6-mm gage diameter. The AM50 alloy was cast by the hot-chamber HPDC process into a steering wheel. Further details of the HPDC process parameters and die geometries are available in Ref. [22]. Thixomolded samples were manufactured using Mg alloy AZ91 by Neue Materialien Fu¨rth GmbH (NMF) in a JLM-220MG casting machine from
Table 1 Alloy compositions for AlSi4MgMn, AM50 and AZ91 (wt.%). Alloy
Al
Si
Mn
Mg
Zn
Fe
AlSi4MgMn AM50 AZ91
Balance 4.79 8.3–9.7
4.30 0.011 60.10
0.67 0.23 0.15
0.19 Balance Balance
– – 0.35–1
0.15 60.002 60.005
Japan Steel Works. The present study examined 6-mm thick plates cast with an initial primary globule fs of 10– 15%. Further details about the Thixomolded specimens are reported by Lohmu¨ller et al. [19]. In order to recreate the relatively rapid in-cavity cooling rates of HPDC and SSM, but without forced convection or applied pressure, a “static cooling” technique was used. Small chips of alloy were completely melted in a 4-mm inner-diameter cylindrical mold to a height of 40 mm and the liquid alloy was held at 80 °C superheat. The cylinder was then immersed into a mixture of ice and water. For Mg alloys, a steel tube was used as a cylindrical vessel. For Al alloys, a quartz-glass tube was used to avoid excessive dissolution of Fe into the melt. Al–5Ti–1B master alloy was added to the AlSi4MgMn alloy to a composition of 0.05 wt.% Ti and 0.01 wt.% B for grain refinement. With this technique, the mean cooling rates in the whole freezing range of AM50 and AlSi4MgMn were approximately 80 and 60 °C s1 respectively, measured by a K-type thermocouple immersed in the alloys. For metallographic examination, specimens were sectioned transverse to the bulk melt flow direction for the HPDC and Thixomolded samples, and in the transverse cross-section of the cylindrical castings produced by the “static cooling” technique. The HPDC specimens were sectioned from the center of the tensile-bar gage length (the same location as sectioned in Ref. [22]). The “statically cooled” samples were sectioned from the middle of the cylinder length. All samples were mounted in cold-setting epoxy resin and prepared to a 1-lm finish by standard metallographic techniques. For optical microscopy, the grain structure was revealed in AlSi4MgMn polished cross-sections by anodizing for 90 s in a 2.6% HBF4 solution, whereas the technique developed by Maltais et al. [23] for AZ91 was applied to the AM50 samples by modifying the etching solution to 130 ml C2H5OH + 70 ml H2O + 1 ml CH3COOH. EBSD samples less than 1 mm thick were sectioned from the etched cross-sections. The samples were reground to remove the etched surface layer, polished to a 1-lm finish and then argon-ion-beam milled in a GatanTM 600TMPA dual ion mill (Gatan Inc., Pleasanton, CA) under conditions of 13° tilt, 4.5 kV, 1 mA for 3.5 h with liquid N2 stage cooling. To acquire grain-misorientation information, EBSD mapping was performed in a JEOLTM JSM6460LA scanning electron microscope (SEM) (Japan Electron Optics Ltd., Tokyo, Japan) equipped with an HKL NordlysTM EBSD detector and Channel 5TM EBSD software (Oxford Instruments, Oxfordshire, UK). For each specimen, mapping results were acquired from at least 200 grains for statistical validity. The mapping step size was selected to have at least five mapping points across each grain width on the cross-section. Fig. 1a shows a typical raw orientation map after EBSD mapping. Typically, some regions are unsuccessfully indexed by the EBSD software, particularly in the eutectic region. In order to obtain statistical data of grain orienta-
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tions, extrapolation to the unindexed regions in the raw map, e.g. Fig. 1a, was performed. Fig. 1b is the extrapolated map from the raw map in Fig. 1a. Three types of GBs—low-angle GBs (LAGBs), high-angle GBs (HAGBs) and coincidence-site-lattice GBs (CSL-GBs)—were determined after result extrapolation in each EBSD map. Similar to previous research [24], LAGBs and HAGBs are defined here as boundaries with a misorientation between 5° and 15°, and more than 15°, respectively. The Brandon criterion [25] was used to determine the angular deviation thresholds from the exact CSL-GBs. For illustration purposes, Fig. 1c shows the three GB types superimposed on the backscattered electron (BSE) image. Extrapolated EBSD maps, such as in Fig. 1b, were used for grain size measurements. Special care was taken to clean up artificial grains resulting from result extrapolation from inaccurately indexed points [18,26]. Regions with misorientations exceeding 15° were defined as different grains in accordance with past studies [18,22,26–28]. Grain size was determined using the equivalent circle diameter (ECD) method. The values plotted are the average of the values obtained from three EBSD maps and the error bars show the range of raw data values. The primary grains measured for HPDC and Thixomolding are in-cavity solidified grains. 3. Results 3.1. Morphology and size of primary grains Typical as-cast microstructures produced by the three casting techniques are shown in Fig. 2a–e. The primary grains have equiaxed morphology in all cases. The HPDC samples contain two populations of primary grains (Fig. 2c): (i) small, globular-rosette in-cavity solidified grains (10 lm) and (ii) large dendritic grains (30– 300 lm) originating from the shot sleeve [15,29–31], socalled ESCs. Fig. 2e shows that there are also two populations of primary grains in the AZ91 Thixomolded specimen: (i) small, globular-rosette in-cavity solidified grains similar to those in HPDC (10 lm) and (ii) large globular grains (50–100 lm) that were created during isothermal stirring of the semi-solid slurry prior to die filling. For the same alloy, the “static cooling” technique produced a significantly larger grain size (100 lm in AlSi4MgMn and 50 lm in AM50) than the in-cavity solidified grains in HPDC (10 lm). The average grain sizes for all samples are shown in Fig. 2f. 3.2. Misorientations between grains Fig. 1. Different types of images can be plotted after EBSD mapping: (a) raw EBSD map, (b) EBSD map in (a) after result extrapolation, and (c) backscattered electron (BSE) image superimposed by GB lines. Unindexed points in (a) are dark green and other colors represent the orientations of grains. In (c), HAGBs, LAGBs and CSL-GBs are marked by white, yellow and red lines, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 3a–e shows extrapolated EBSD maps superimposed on BSE micrographs for each sample. The LAGBs, CSLGBs and HAGBs are shown in different colors. The fraction of ESCs in HPDC samples is relatively low such that there are only few ESCs in each EBSD map (Fig. 3c and d). The AZ91 Thixomolded sample, Fig. 3e, reveals that
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Fig. 2. Typical microstructures of the casting specimens in this study: (a) and (b) etched optical micrographs under polarized light, (c) and (e) etched optical micrographs under the normal bright-field mode, and (d) BSE micrograph. (f) Mean grain size for each sample type measured by the equivalent circle diameter method. The primary grains measured for HPDC and Thixomolding are in-cavity solidified grains.
all GBs between primary globules are low-energy GBs (LAGBs and CSL-GBs). Fig. 4 quantifies the fraction of LAGBs and CSL-GBs in each of the samples investigated. Note that only GBs between in-cavity solidified grains are shown for HPDC and Thixomolded samples because the number of GBs between presolidified crystals is not sufficient for statistical validity. For each alloy, the fraction of low-energy GBs (LAGBs and CSL-GBs) is higher in the HPDC sample than in the “statically cooled” sample. Compared to the “statically cooled” sample, the increased fraction of low-energy GBs in AM50 is from both the LAGB and CSL-GB fractions while the increase in low-energy GBs in AlSi4MgMn is from the LAGB fraction only. The fractions of LAGBs and CSLGBs in AM50 HPDC and AZ91 Thixomolded specimens appear to be similar. In “statically cooled” specimens, both the fractions of LAGBs and CSL-GBs in AlSi4MgMn are much higher than in AM50. In HPDC samples, the fraction of LAGBs in AlSi4MgMn is much higher than in AM50 whereas the fractions of CSL-GBs are similar.
The misorientation-angle distributions for all samples are shown in Fig. 5a–e. Similar to Fig. 4, the distributions for HPDC and Thixomolding only include for in-cavity solidified grains. The theoretical distribution expected from a randomly oriented grain structure, the so-called Mackenzie curve [32], is plotted along with the misorientation-angle distribution for each casting microstructure in Fig. 5a–e. Fig. 5a–e shows that the misorientation-angle distributions for HPDC and Thixomolded specimens deviate substantially from the random distributions while those for “statically cooled” specimens fit much better with the Mackenzie curves. In HPDC and Thixomolded samples (Fig. 5c–e), the fraction of LAGBs is substantially higher than the theoretical prediction and there are local peaks at the misorientation angles corresponding with some CSL-GBs. Furthermore, at these misorientation angles, the proportion of CSL-GBs is relatively high as indicated by the misorientation-axis distributions in Fig. 5c–e. In “statically cooled” samples (Fig. 5a and b), although the
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Fig. 3. BSE micrographs superimposed by GB lines in color, as in Fig. 1c, for all types of samples.
local peaks in the misorientation plots are not as pronounced as in HPDC, the misorientation-axis distributions in Fig. 5a and b still show a tendency for GBs to correspond with some CSL-GBs. In addition, the fraction of LAGBs in “statically cooled” samples is slightly higher than random (Fig. 5a and b). For the same casting technique, the fraction of LAGBs in AlSi4MgMn (Fig. 5a and c) deviates more from the Mackenzie prediction than in AM50 (Fig. 5b and d).
3.3. In-grain misorientations In-grain misorientations (misorientations within the same grain) were measured in most of the grains shown in Fig. 3a–e. In-grain misorientations are relatively small (less than 1°) for most grains measured, with ESCs in HPDC samples being the only exception. For example, the misorientation profile in Fig. 6a shows that in-grain misorientations are less than 1° even in large primary glob-
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orientations of primary grains in any of the samples. This suggests that the agglomeration occurred on the scale of several crystals and therefore no long-range network of agglomerated crystals existed in the samples.
Fig. 4. Length fractions of LAGBs and CSL-GBs in all types of specimens. In HPDC and Thixomolded samples, only GBs between incavity solidified grains were included in the measurements. The values plotted are the average of the values obtained from two EBSD maps and the error bars show the range of raw data values.
ules in the Thixomolded AZ91 sample. In contrast, Fig. 6b shows significant in-grain misorientations of 3° for an ESC in the AlSi4MgMn HPDC sample. 4. Discussion 4.1. Crystal agglomeration By producing samples with different casting methods where the solidifying alloy experiences different levels of external mechanical stresses, the effects of mechanical stresses applied during solidification can be assessed. The length fraction of low-energy GBs shown in Fig. 4 is associated with crystal agglomeration during solidification. When two growing primary crystals impinge with one another, GB formation depends on the interfacial energy of the potential new GB. Coalescence or bridging is the transformation of two impinging solidification fronts into a solid bridge [33] and in this case a new agglomerate can form. Therefore, coalescence can be considered as the disappearance of two solid–liquid interfaces, each with interfacial energy cs/l, and the formation of a GB with interfacial energy cgb [34]. When two solidifying crystals impinge on one another, bridging occurs readily if cgb < 2cs/l [10,33]. On the other hand, some energy is required to form a new boundary if cgb > 2cs/l cgb is a function of the misorientation between the two impinging crystals [34]. cgb < 2cs/l occurs if the misorientation is less than 15° (LAGBs) or corresponds to a CSL-GB [34]. In this case, there is an attractive force to bring the two crystals together and coalescence occurs as soon as the two interfaces are close enough (at a distance on the order of ten atomic distances [33]). Only GBs with cgb < 2cs/l are thought to form after a collision of two crystals [35]. If not, a liquid film is stable and the colliding crystals bounce back. It should be pointed out that pole figures of the EBSD data in Fig. 3a–e do not show any texture/preferred
4.1.1. Primary crystals in “statically cooled” samples In “statically cooled” samples, although external stresses acting on the alloy during solidification are insignificant, at fs less than the dendrite coherency solid fraction (fsCoh ) [36– 39], primary equiaxed crystals can still move due to gravity, buoyancy and the pressure differential resulting from solidification shrinkage. This movement of solidifying equiaxed crystals promotes collisions between them. It has been estimated for a range of alloys that solid–solid interfaces are stable at only 1–7% of boundaries [40–42] and Martin et al. [43] estimate that only 2% of collisions between crystals will lead to new GBs, assuming that only LAGBs are stable. The number of low-energy GBs formed by crystal collisions therefore depends on the number of collisions. Increasing the number of collisions leads to an increase in the number of low-energy GBs. As the formation of lowenergy GBs is energetically favorable, their fraction in the “statically cooled” samples in this study is slightly higher than the prediction by the Mackenzie curve which does not take into account GB energy (Fig. 5a and b). The difference in the fractions of LAGBs and CSL-GBs between the aluminum alloy AlSi4MgMn and the magnesium alloys AM50/AZ91 in Fig. 4 could arise largely from the difference in crystal structure: the cubic structure of Al is more symmetrical than the hcp structure of Mg. Therefore, it is more likely for two randomly colliding crystals to have a small misorientation angle or a misorientation corresponding to a CSL-GB in Al than in Mg. However, it is also possible that the difference in the LAGBs and CSL-GBs fractions partly resulted from the difference in surface tension of liquid in AlSi4MgMn and AM50/AZ91. 4.1.2. In-cavity solidified crystals in HPDC and Thixomolding The relatively high proportion of low-energy GBs among in-cavity solidified grains in HPDC and Thixomolded samples (Fig. 4) indicates that agglomeration of in-cavity solidified grains occurs. The high shear stresses and pressure applied on the solidifying alloy during HPDC and Thixomolding are likely to promote crystal collisions and result in a strong degree of crystal agglomeration. Agglomeration of in-cavity solidified grains can occur during (i) the die-filling stage, where there is both solidification and forced convection, and/or (ii) the intensification stage, when mushy material is pushed deeper into the die cavity. External mechanical stresses acting on the solidifying alloy in HPDC and Thixomolding are much higher than in the “statically cooled” casting. During die filling, the mixture of molten alloy and solid crystals is injected into the die cavity. The solid crystals in HPDC and Thixomolding are ESCs and primary globules, respectively. In-cavity
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Fig. 5. Misorientation-angle distributions for all types of specimens. The misorientation angles less than 5° are not plotted in order to represent only the boundaries between regions with significantly different orientations. The fractions are displayed in boundary length. Misorientation-axis distributions are shown for some angular ranges. The Mackenzie curve for each case is shown by a line.
solidified crystals start to form in the early stages of die filling under conditions of strong forced convection and relatively high pressure. The driving force for agglomeration among in-cavity solidified crystals is simply turbulent flow, similar to the shearing of primary globules, but the shearing time for in-cavity solidified crystals here is very short
(less than a tenth of a second) and not all in-cavity solidified crystals have nucleated. For this reason, it is not likely that shear during die filling alone leads to agglomeration among in-cavity solidified grain in this work. In-cavity solidified crystals in both processes are subjected to further shear and compressive stresses during
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Fig. 6. Misorientation profiles relative to the first point along the lines drawn across regions within the same grains: (a) within a primary globule in an AZ91 Thixomolded sample and (b) within an ESC in an AlSi4MgMn HPDC sample.
the intensification stage. A pressure on the order of 10– 100 MPa is applied to the biscuit to feed the solidification shrinkage and compress entrapped air in the cavity. The intensification pressure is transmitted through the runner system to the casting, and subsequently transports mushy material deeper into the die cavity. The deformation of the partially solid material with fs < fsCoh is in the form of suspension flow which promotes collisions between solidifying crystals. In the deforming mush with fs P fsCoh , the intensification pressure pushes crystals into
one another (collisions) and the mechanism of mush deformation is grain rearrangement [22] which causes more collisions (the misorientations between neighboring crystals change during deformation due to crystal rotation and translation within the crystal assembly). This could explain the stronger degree of crystal agglomeration due to compressive loading created by intensification pressure, which has also been observed previously in a compression experiment on partially remelted Al alloy by Verrier et al. [44].
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The very high fraction of LAGBs in the AlSi4MgMn HPDC sample (Figs. 4 and 5c) could be explained by the relatively symmetrical geometry of the fcc crystal structure and therefore two equiaxed crystals have a relatively high probability of colliding with a small misorientation angle. However, other reasons, such as the values of GB energy, would also be expected to have a significant effect. 4.1.3. ESCs in HPDC and primary globules in Thixomolding ESCs in HPDC microstructures can be compared with primary globules in SSM microstructures, such as Thixomolding, as they are both crystals that form before die filling. Since ESCs and primary globules are present in the alloy before die filling, they are exposed to significant mechanical stresses during both die filling and pressure intensification. They experience large shear stresses when they are squeezed through the relatively thin gate during die filling. The velocity at the gate in HPDC and Thixomolding is typically in the range 10–50 m s1. In Thixomolding and other SSM techniques, partially solidified alloy is isothermally sheared, mechanically or electromagnetically, and fragmentation of dendritic crystals occurs due to strong forced convection. Fragmented dendrites subsequently become primary globules [2]. Further shearing has been found to cause agglomeration of these primary globules by the evidence that most GBs between large globules are low-energy GBs (Fig. 3e) [35,45]. The effect of stirring on the degree of crystal agglomeration has been studied previously: Arnberg et al. [46] found that the fraction of low-energy GBs in castings solidifying with electromagnetic stirring is significantly higher than in castings solidifying without stirring. They suggested that shear stresses applied to the semi-solid alloy, e.g. by electromagnetic stirring, increase the crystal collision frequency and therefore promote agglomeration of crystals. Even though primary globules in Thixomolding and ESCs in HPDC are similar in some aspects, there are a number of important differences. First, primary globules are intentionally made and the process conditions, e.g. temperature and stirring parameters, are carefully monitored, while the formation of ESCs is poorly controlled. Second, the shearing time of primary globules in Thixomolding (typically several minutes [19]) is significantly longer than that for ESCs in HPDC (less than tenth of a second during die filling and just several seconds during intensification). Thus, the number of collisions is expected to be significantly higher between globules than ESCs. These differences make the formation of GBs between ESCs significantly different from the formation of GBs between primary globules: agglomeration of primary globules, evidenced in the EBSD mapping results such as in Fig. 3e, was not observed for ESCs in this study. This suggests that external mechanical stresses acting on ESCs during die filling and intensification do not cause agglomeration of ESCs to occur frequently. The combination of a low fraction of ESCs and short period of shear
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must have led to fewer collisions and therefore a low number of ESC–ESC GBs in this study. 4.2. Crystal bending Another possible reason for the presence of a high fraction of LAGBs in the microstructures in this study is crystal bending during solidification. Doherty [47] has suggested that primary crystals can be bent by either mechanical stresses or changes in the growth directions. When mechanical stresses act on solidifying crystals, arising from forced convection [48], solidification contraction [47], gravitational force [7], etc., the crystals can plastically deform leading to bending of their arms. The preferred growth directions of dendrites in certain alloys are known to vary either abruptly or continuously, e.g. primary dendrite growth directions in Al–Zn alloys exhibit continuous changes of direction from h1 0 0i to h1 1 0i over a range of alloy composition [49]. In this case, a change in alloy composition ahead of the solidification front, e.g. due to convection, may lead a change in growth direction leading to crystal bending. If a misorientation between adjacent regions in a bent dendrite arm is within 0–15° (LAGBs), the bent arm still exists. On the other hand, the arm is broken if misorientations exceeds 15° (HAGBs) because cgb > 2cs/l, causing wetting of the liquid at the GB and separation of the crystals. In columnar crystals, crystal bending can be evidenced by in-grain misorientations when sectioned through the longitudinal section of the columnar dendrite. If sectioned through the plane transverse to the primary trunk, dendrite bending can be evidenced by LAGBs between primary and secondary arms. In equiaxed crystals, the growth direction is almost isotropic. If dendrite bending exists in equiaxed structures, it can be observed as either in-grain misorientations or LAGBs. In this study, the primary crystals in the “statically cooled” samples and in-cavity solidified crystals in HPDC and Thixomolding contain only relatively low (<1°) in-grain misorientations (e.g. Fig. 6a). This suggests that the effect of dendrite bending is insignificant or too small to be detected by EBSD. The reason is most likely the compact morphology of the equiaxed crystals in this study. When shear stresses, originating from solidification shrinkage, gravity, external mechanical stresses, etc., act on compact-shaped equiaxed crystals, the crystals may not be able to deform as significantly as in highly branched crystals. ESCs in HPDC and primary globules in Thixomolding are both subjected to more external mechanical stresses than in-cavity solidified grains, particularly at the gate during die filling. The significant in-grain misorientations for the ESCs shown in Fig. 6b suggest that shearing has caused bending of the ESCs. This could be a result of the higher bending moment acting on the longer dendrite arms of ESCs compared to those acting on the more compact globules in the Thixomolded sample and the small, compact incavity solidifying crystals.
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5. Conclusions Microstructural characterization of AlSi4MgMn, AM50 and AZ91 specimens produced by HPDC, Thixomolding and a “static cooling” technique has been performed using EBSD. The study has focused on the influence of external mechanical stresses on agglomeration and bending of equiaxed crystals during solidification. The following conclusions can be drawn. 1. In all samples, the fraction of low-energy GBs is higher than the theoretical prediction for a randomly oriented microstructure. This is attributed to the attractive force between solidifying crystals which have a misorientation relationship favoring the formation of a low-energy GB. 2. In HPDC and Thixomolded samples, the fraction of low-energy GBs among in-cavity solidified grains is significantly higher than in “statically cooled” specimens. This is attributed to the increased number of crystal collisions during HPDC and Thixomolding, which promotes agglomeration of favorably oriented crystals. 3. The distribution of grain misorientations among in-cavity solidified grains in HPDC and Thixomolded specimens are comparable. 4. In-grain misorientations are not significant in primary crystals in all types of specimens, except ESCs in HPDC. This suggests that bending of solidifying crystals is not significant in equiaxed crystals with compact morphology. In contrast, bending was observed in ESCs, which is attributed to the increased bending moment acting on dendrite arms in crystals with a more branched morphology.
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Acknowledgments
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The authors acknowledge the support from the CAST Cooperative Research Centre and the Australian Research Council’s Centre of Excellence for Design in Light Metals. The technical, scientific and financial assistance from the Australian Microscopy and Microanalysis Research Facility (AMMRF) are also acknowledged. We thank Dr. A. Lohmu¨ller (NMF GmbH, Nu¨rnberg, Germany) for supplying the ThixomoldedÒ castings. SO gratefully acknowledges the financial support provided by the Royal Thai Government and a UQ Graduate School Research Travel Grant (GSRTG). SO thanks Drs. Peter Hines and Gwe´nae¨lle Proust (EMU, The University of Sydney) for hospitality and support, and Drs. Tim Williams and Peter Miller (MCEM, Monash University) for the help on preliminary EBSD results.
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Acta Materialia 58 (2010) 261–271 www.elsevier.com/locate/actamat
Agglomeration and bending of equiaxed crystals during solidification of hypoeutectic Al and Mg alloys S. Otarawanna a,*, C.M. Gourlay b, H.I. Laukli c, A.K. Dahle d a
CAST CRC, Materials Engineering, The University of Queensland, Brisbane, Qld 4072, Australia b Department of Materials, Imperial College London, London SW7 2AZ, UK c Hydro Aluminium, Research & Technology Development, N-6601 Sunndalsøra, Norway d ARC CoE for Design in Light Metals, Materials Engineering, The University of Queensland, Brisbane, Qld 4072, Australia Received 11 August 2009; received in revised form 31 August 2009; accepted 2 September 2009 Available online 1 October 2009
Abstract Agglomeration and bending of equiaxed crystals were studied by microstructural characterization of specimens produced by near-static cooling, high-pressure die casting and ThixomoldingÒ, where the solidifying crystals experience different levels of mechanical stresses. EBSD was used to acquire statistical grain misorientation data which is linked to crystal agglomeration and bending behavior during solidification. An aluminum alloy and two magnesium alloys were used to compare grain misorientations for different crystal structures. The length fraction of low-energy grain boundaries in HPDC and Thixomolded samples was substantially higher than in “statically cooled” samples. This is attributed to the high shear stresses and pressure applied on the solidifying alloy, which promote crystal collisions and agglomeration. In-grain misorientations were found to be significant only in branched dendritic crystals which were subjected to significant shear stresses. This is related to the increased bending moment acting on long, protruding dendrite arms compared to more compact crystal morphologies. Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Solidification microstructure; Electron backscattering diffraction (EBSD); Aluminum alloys; Grain boundary energy; Misorientation
1. Introduction Agglomeration and bending of primary crystals in the mushy zone during solidification are important phenomena influencing many aspects of casting processes and often affect the microstructure of the as-cast component. Both phenomena strongly influence the rheology of partially solid alloys. For example, agglomeration (and disagglomeration) of crystals is the origin of thixotropy in semi-solid metallic suspensions [1–4]. Agglomeration of crystals often entraps liquid, increasing the effective solid fraction and the apparent viscosity [1–4]. Bending of solidifying dendrite arms creates misorientation within crystals by dynamic recovery and can subsequently cause arm fragmentation *
Corresponding author. Tel.: +61 7 3365 1387; fax: +61 7 3365 3888. E-mail addresses: [email protected], [email protected] (S. Otarawanna).
[5–9], which reduces the grain size and the mush strength [2,4,10,11]. Agglomeration and bending are both strongly influenced by the application of external mechanical stresses during solidification [2,4] and these phenomena are therefore particularly important in casting processes where the alloy is forced into a mold during solidification. Two types of process where the solidifying alloy is severely deformed are high-pressure die casting (HPDC) and the semi-solid metal processing (SSM or SSP) variants, ranging from thixocasting to thixoforging. In the HPDC process, fully liquid alloy is transferred into the shot chamber where some uncontrolled solidification occurs, resulting in socalled externally solidified crystals (ESCs), which occupy up to 30 vol.% in the as-cast microstructure [12]. ESCs undergo intense shear during die filling where shear rates at the gate are typically 104 s1 [13]. Some further solidification occurs during die filling under strong forced
1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2009.09.002
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convection. Once the cavity is full, intensification pressure of up to 120 MPa [14] is applied to the casting and the solidifying alloy is forced deeper into the cavity to feed the casting. Since the shot-sleeve cooling rate (on the order of 10 °C s1 [15]) is significantly lower than the in-cavity cooling rate (in the range 102–103 °C s1 [16,17]), ESCs are larger than in-cavity solidified grains, often by an order of magnitude [18]. In SSM, a solid–liquid mixture with globular solid morphology and controlled solid fraction (fs) is produced prior to die filling. This mixture is often generated by strong forced convection at a known temperature in the freezing range [2,4]. The SSM variant examined in this paper is Thixomolding, also known as metal injection molding [19]. In Thixomolding, alloy chips are fed into a screw mixer and heated to a temperature in the freezing range while being intensely sheared. The primary globule fs is typically 10% or below [20], which is lower than is typical for SSM variants such as thixocasting and rheocasting (fs 50% [21]). The slurry is then injected rapidly into the die cavity, after which an after-pressure, analogous to intensification pressure in HPDC, is applied to assist feeding. In this work, electron backscatter diffraction (EBSD) is used to study agglomeration and bending of equiaxed crystals in HPDC, Thixomolding and on samples solidified without forced convection or an applied pressure. Additionally, an AlSi4MgMn aluminum alloy, and AM50 and AZ91 magnesium alloys, are used to compare grain misorientations for different crystal structures, i.e. face-centered cubic (fcc) in Al and hexagonal close-packed (hcp) in Mg. 2. Experimental Al alloy AlSi4MgMn and Mg alloys AM50 and AZ91, with the compositions shown in Table 1, were used in this work. Alloys were cast by HPDC, Thixomolding and a “static cooling” technique. HPDC samples were produced with 80 °C superheat. The AlSi4MgMn alloy was cast by the cold-chamber HPDC process into a tensile bar with 125-mm total length, 30-mm gage length and 6-mm gage diameter. The AM50 alloy was cast by the hot-chamber HPDC process into a steering wheel. Further details of the HPDC process parameters and die geometries are available in Ref. [22]. Thixomolded samples were manufactured using Mg alloy AZ91 by Neue Materialien Fu¨rth GmbH (NMF) in a JLM-220MG casting machine from
Table 1 Alloy compositions for AlSi4MgMn, AM50 and AZ91 (wt.%). Alloy
Al
Si
Mn
Mg
Zn
Fe
AlSi4MgMn AM50 AZ91
Balance 4.79 8.3–9.7
4.30 0.011 60.10
0.67 0.23 0.15
0.19 Balance Balance
– – 0.35–1
0.15 60.002 60.005
Japan Steel Works. The present study examined 6-mm thick plates cast with an initial primary globule fs of 10– 15%. Further details about the Thixomolded specimens are reported by Lohmu¨ller et al. [19]. In order to recreate the relatively rapid in-cavity cooling rates of HPDC and SSM, but without forced convection or applied pressure, a “static cooling” technique was used. Small chips of alloy were completely melted in a 4-mm inner-diameter cylindrical mold to a height of 40 mm and the liquid alloy was held at 80 °C superheat. The cylinder was then immersed into a mixture of ice and water. For Mg alloys, a steel tube was used as a cylindrical vessel. For Al alloys, a quartz-glass tube was used to avoid excessive dissolution of Fe into the melt. Al–5Ti–1B master alloy was added to the AlSi4MgMn alloy to a composition of 0.05 wt.% Ti and 0.01 wt.% B for grain refinement. With this technique, the mean cooling rates in the whole freezing range of AM50 and AlSi4MgMn were approximately 80 and 60 °C s1 respectively, measured by a K-type thermocouple immersed in the alloys. For metallographic examination, specimens were sectioned transverse to the bulk melt flow direction for the HPDC and Thixomolded samples, and in the transverse cross-section of the cylindrical castings produced by the “static cooling” technique. The HPDC specimens were sectioned from the center of the tensile-bar gage length (the same location as sectioned in Ref. [22]). The “statically cooled” samples were sectioned from the middle of the cylinder length. All samples were mounted in cold-setting epoxy resin and prepared to a 1-lm finish by standard metallographic techniques. For optical microscopy, the grain structure was revealed in AlSi4MgMn polished cross-sections by anodizing for 90 s in a 2.6% HBF4 solution, whereas the technique developed by Maltais et al. [23] for AZ91 was applied to the AM50 samples by modifying the etching solution to 130 ml C2H5OH + 70 ml H2O + 1 ml CH3COOH. EBSD samples less than 1 mm thick were sectioned from the etched cross-sections. The samples were reground to remove the etched surface layer, polished to a 1-lm finish and then argon-ion-beam milled in a GatanTM 600TMPA dual ion mill (Gatan Inc., Pleasanton, CA) under conditions of 13° tilt, 4.5 kV, 1 mA for 3.5 h with liquid N2 stage cooling. To acquire grain-misorientation information, EBSD mapping was performed in a JEOLTM JSM6460LA scanning electron microscope (SEM) (Japan Electron Optics Ltd., Tokyo, Japan) equipped with an HKL NordlysTM EBSD detector and Channel 5TM EBSD software (Oxford Instruments, Oxfordshire, UK). For each specimen, mapping results were acquired from at least 200 grains for statistical validity. The mapping step size was selected to have at least five mapping points across each grain width on the cross-section. Fig. 1a shows a typical raw orientation map after EBSD mapping. Typically, some regions are unsuccessfully indexed by the EBSD software, particularly in the eutectic region. In order to obtain statistical data of grain orienta-
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tions, extrapolation to the unindexed regions in the raw map, e.g. Fig. 1a, was performed. Fig. 1b is the extrapolated map from the raw map in Fig. 1a. Three types of GBs—low-angle GBs (LAGBs), high-angle GBs (HAGBs) and coincidence-site-lattice GBs (CSL-GBs)—were determined after result extrapolation in each EBSD map. Similar to previous research [24], LAGBs and HAGBs are defined here as boundaries with a misorientation between 5° and 15°, and more than 15°, respectively. The Brandon criterion [25] was used to determine the angular deviation thresholds from the exact CSL-GBs. For illustration purposes, Fig. 1c shows the three GB types superimposed on the backscattered electron (BSE) image. Extrapolated EBSD maps, such as in Fig. 1b, were used for grain size measurements. Special care was taken to clean up artificial grains resulting from result extrapolation from inaccurately indexed points [18,26]. Regions with misorientations exceeding 15° were defined as different grains in accordance with past studies [18,22,26–28]. Grain size was determined using the equivalent circle diameter (ECD) method. The values plotted are the average of the values obtained from three EBSD maps and the error bars show the range of raw data values. The primary grains measured for HPDC and Thixomolding are in-cavity solidified grains. 3. Results 3.1. Morphology and size of primary grains Typical as-cast microstructures produced by the three casting techniques are shown in Fig. 2a–e. The primary grains have equiaxed morphology in all cases. The HPDC samples contain two populations of primary grains (Fig. 2c): (i) small, globular-rosette in-cavity solidified grains (10 lm) and (ii) large dendritic grains (30– 300 lm) originating from the shot sleeve [15,29–31], socalled ESCs. Fig. 2e shows that there are also two populations of primary grains in the AZ91 Thixomolded specimen: (i) small, globular-rosette in-cavity solidified grains similar to those in HPDC (10 lm) and (ii) large globular grains (50–100 lm) that were created during isothermal stirring of the semi-solid slurry prior to die filling. For the same alloy, the “static cooling” technique produced a significantly larger grain size (100 lm in AlSi4MgMn and 50 lm in AM50) than the in-cavity solidified grains in HPDC (10 lm). The average grain sizes for all samples are shown in Fig. 2f. 3.2. Misorientations between grains Fig. 1. Different types of images can be plotted after EBSD mapping: (a) raw EBSD map, (b) EBSD map in (a) after result extrapolation, and (c) backscattered electron (BSE) image superimposed by GB lines. Unindexed points in (a) are dark green and other colors represent the orientations of grains. In (c), HAGBs, LAGBs and CSL-GBs are marked by white, yellow and red lines, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 3a–e shows extrapolated EBSD maps superimposed on BSE micrographs for each sample. The LAGBs, CSLGBs and HAGBs are shown in different colors. The fraction of ESCs in HPDC samples is relatively low such that there are only few ESCs in each EBSD map (Fig. 3c and d). The AZ91 Thixomolded sample, Fig. 3e, reveals that
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Fig. 2. Typical microstructures of the casting specimens in this study: (a) and (b) etched optical micrographs under polarized light, (c) and (e) etched optical micrographs under the normal bright-field mode, and (d) BSE micrograph. (f) Mean grain size for each sample type measured by the equivalent circle diameter method. The primary grains measured for HPDC and Thixomolding are in-cavity solidified grains.
all GBs between primary globules are low-energy GBs (LAGBs and CSL-GBs). Fig. 4 quantifies the fraction of LAGBs and CSL-GBs in each of the samples investigated. Note that only GBs between in-cavity solidified grains are shown for HPDC and Thixomolded samples because the number of GBs between presolidified crystals is not sufficient for statistical validity. For each alloy, the fraction of low-energy GBs (LAGBs and CSL-GBs) is higher in the HPDC sample than in the “statically cooled” sample. Compared to the “statically cooled” sample, the increased fraction of low-energy GBs in AM50 is from both the LAGB and CSL-GB fractions while the increase in low-energy GBs in AlSi4MgMn is from the LAGB fraction only. The fractions of LAGBs and CSLGBs in AM50 HPDC and AZ91 Thixomolded specimens appear to be similar. In “statically cooled” specimens, both the fractions of LAGBs and CSL-GBs in AlSi4MgMn are much higher than in AM50. In HPDC samples, the fraction of LAGBs in AlSi4MgMn is much higher than in AM50 whereas the fractions of CSL-GBs are similar.
The misorientation-angle distributions for all samples are shown in Fig. 5a–e. Similar to Fig. 4, the distributions for HPDC and Thixomolding only include for in-cavity solidified grains. The theoretical distribution expected from a randomly oriented grain structure, the so-called Mackenzie curve [32], is plotted along with the misorientation-angle distribution for each casting microstructure in Fig. 5a–e. Fig. 5a–e shows that the misorientation-angle distributions for HPDC and Thixomolded specimens deviate substantially from the random distributions while those for “statically cooled” specimens fit much better with the Mackenzie curves. In HPDC and Thixomolded samples (Fig. 5c–e), the fraction of LAGBs is substantially higher than the theoretical prediction and there are local peaks at the misorientation angles corresponding with some CSL-GBs. Furthermore, at these misorientation angles, the proportion of CSL-GBs is relatively high as indicated by the misorientation-axis distributions in Fig. 5c–e. In “statically cooled” samples (Fig. 5a and b), although the
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Fig. 3. BSE micrographs superimposed by GB lines in color, as in Fig. 1c, for all types of samples.
local peaks in the misorientation plots are not as pronounced as in HPDC, the misorientation-axis distributions in Fig. 5a and b still show a tendency for GBs to correspond with some CSL-GBs. In addition, the fraction of LAGBs in “statically cooled” samples is slightly higher than random (Fig. 5a and b). For the same casting technique, the fraction of LAGBs in AlSi4MgMn (Fig. 5a and c) deviates more from the Mackenzie prediction than in AM50 (Fig. 5b and d).
3.3. In-grain misorientations In-grain misorientations (misorientations within the same grain) were measured in most of the grains shown in Fig. 3a–e. In-grain misorientations are relatively small (less than 1°) for most grains measured, with ESCs in HPDC samples being the only exception. For example, the misorientation profile in Fig. 6a shows that in-grain misorientations are less than 1° even in large primary glob-
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orientations of primary grains in any of the samples. This suggests that the agglomeration occurred on the scale of several crystals and therefore no long-range network of agglomerated crystals existed in the samples.
Fig. 4. Length fractions of LAGBs and CSL-GBs in all types of specimens. In HPDC and Thixomolded samples, only GBs between incavity solidified grains were included in the measurements. The values plotted are the average of the values obtained from two EBSD maps and the error bars show the range of raw data values.
ules in the Thixomolded AZ91 sample. In contrast, Fig. 6b shows significant in-grain misorientations of 3° for an ESC in the AlSi4MgMn HPDC sample. 4. Discussion 4.1. Crystal agglomeration By producing samples with different casting methods where the solidifying alloy experiences different levels of external mechanical stresses, the effects of mechanical stresses applied during solidification can be assessed. The length fraction of low-energy GBs shown in Fig. 4 is associated with crystal agglomeration during solidification. When two growing primary crystals impinge with one another, GB formation depends on the interfacial energy of the potential new GB. Coalescence or bridging is the transformation of two impinging solidification fronts into a solid bridge [33] and in this case a new agglomerate can form. Therefore, coalescence can be considered as the disappearance of two solid–liquid interfaces, each with interfacial energy cs/l, and the formation of a GB with interfacial energy cgb [34]. When two solidifying crystals impinge on one another, bridging occurs readily if cgb < 2cs/l [10,33]. On the other hand, some energy is required to form a new boundary if cgb > 2cs/l cgb is a function of the misorientation between the two impinging crystals [34]. cgb < 2cs/l occurs if the misorientation is less than 15° (LAGBs) or corresponds to a CSL-GB [34]. In this case, there is an attractive force to bring the two crystals together and coalescence occurs as soon as the two interfaces are close enough (at a distance on the order of ten atomic distances [33]). Only GBs with cgb < 2cs/l are thought to form after a collision of two crystals [35]. If not, a liquid film is stable and the colliding crystals bounce back. It should be pointed out that pole figures of the EBSD data in Fig. 3a–e do not show any texture/preferred
4.1.1. Primary crystals in “statically cooled” samples In “statically cooled” samples, although external stresses acting on the alloy during solidification are insignificant, at fs less than the dendrite coherency solid fraction (fsCoh ) [36– 39], primary equiaxed crystals can still move due to gravity, buoyancy and the pressure differential resulting from solidification shrinkage. This movement of solidifying equiaxed crystals promotes collisions between them. It has been estimated for a range of alloys that solid–solid interfaces are stable at only 1–7% of boundaries [40–42] and Martin et al. [43] estimate that only 2% of collisions between crystals will lead to new GBs, assuming that only LAGBs are stable. The number of low-energy GBs formed by crystal collisions therefore depends on the number of collisions. Increasing the number of collisions leads to an increase in the number of low-energy GBs. As the formation of lowenergy GBs is energetically favorable, their fraction in the “statically cooled” samples in this study is slightly higher than the prediction by the Mackenzie curve which does not take into account GB energy (Fig. 5a and b). The difference in the fractions of LAGBs and CSL-GBs between the aluminum alloy AlSi4MgMn and the magnesium alloys AM50/AZ91 in Fig. 4 could arise largely from the difference in crystal structure: the cubic structure of Al is more symmetrical than the hcp structure of Mg. Therefore, it is more likely for two randomly colliding crystals to have a small misorientation angle or a misorientation corresponding to a CSL-GB in Al than in Mg. However, it is also possible that the difference in the LAGBs and CSL-GBs fractions partly resulted from the difference in surface tension of liquid in AlSi4MgMn and AM50/AZ91. 4.1.2. In-cavity solidified crystals in HPDC and Thixomolding The relatively high proportion of low-energy GBs among in-cavity solidified grains in HPDC and Thixomolded samples (Fig. 4) indicates that agglomeration of in-cavity solidified grains occurs. The high shear stresses and pressure applied on the solidifying alloy during HPDC and Thixomolding are likely to promote crystal collisions and result in a strong degree of crystal agglomeration. Agglomeration of in-cavity solidified grains can occur during (i) the die-filling stage, where there is both solidification and forced convection, and/or (ii) the intensification stage, when mushy material is pushed deeper into the die cavity. External mechanical stresses acting on the solidifying alloy in HPDC and Thixomolding are much higher than in the “statically cooled” casting. During die filling, the mixture of molten alloy and solid crystals is injected into the die cavity. The solid crystals in HPDC and Thixomolding are ESCs and primary globules, respectively. In-cavity
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Fig. 5. Misorientation-angle distributions for all types of specimens. The misorientation angles less than 5° are not plotted in order to represent only the boundaries between regions with significantly different orientations. The fractions are displayed in boundary length. Misorientation-axis distributions are shown for some angular ranges. The Mackenzie curve for each case is shown by a line.
solidified crystals start to form in the early stages of die filling under conditions of strong forced convection and relatively high pressure. The driving force for agglomeration among in-cavity solidified crystals is simply turbulent flow, similar to the shearing of primary globules, but the shearing time for in-cavity solidified crystals here is very short
(less than a tenth of a second) and not all in-cavity solidified crystals have nucleated. For this reason, it is not likely that shear during die filling alone leads to agglomeration among in-cavity solidified grain in this work. In-cavity solidified crystals in both processes are subjected to further shear and compressive stresses during
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Fig. 6. Misorientation profiles relative to the first point along the lines drawn across regions within the same grains: (a) within a primary globule in an AZ91 Thixomolded sample and (b) within an ESC in an AlSi4MgMn HPDC sample.
the intensification stage. A pressure on the order of 10– 100 MPa is applied to the biscuit to feed the solidification shrinkage and compress entrapped air in the cavity. The intensification pressure is transmitted through the runner system to the casting, and subsequently transports mushy material deeper into the die cavity. The deformation of the partially solid material with fs < fsCoh is in the form of suspension flow which promotes collisions between solidifying crystals. In the deforming mush with fs P fsCoh , the intensification pressure pushes crystals into
one another (collisions) and the mechanism of mush deformation is grain rearrangement [22] which causes more collisions (the misorientations between neighboring crystals change during deformation due to crystal rotation and translation within the crystal assembly). This could explain the stronger degree of crystal agglomeration due to compressive loading created by intensification pressure, which has also been observed previously in a compression experiment on partially remelted Al alloy by Verrier et al. [44].
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The very high fraction of LAGBs in the AlSi4MgMn HPDC sample (Figs. 4 and 5c) could be explained by the relatively symmetrical geometry of the fcc crystal structure and therefore two equiaxed crystals have a relatively high probability of colliding with a small misorientation angle. However, other reasons, such as the values of GB energy, would also be expected to have a significant effect. 4.1.3. ESCs in HPDC and primary globules in Thixomolding ESCs in HPDC microstructures can be compared with primary globules in SSM microstructures, such as Thixomolding, as they are both crystals that form before die filling. Since ESCs and primary globules are present in the alloy before die filling, they are exposed to significant mechanical stresses during both die filling and pressure intensification. They experience large shear stresses when they are squeezed through the relatively thin gate during die filling. The velocity at the gate in HPDC and Thixomolding is typically in the range 10–50 m s1. In Thixomolding and other SSM techniques, partially solidified alloy is isothermally sheared, mechanically or electromagnetically, and fragmentation of dendritic crystals occurs due to strong forced convection. Fragmented dendrites subsequently become primary globules [2]. Further shearing has been found to cause agglomeration of these primary globules by the evidence that most GBs between large globules are low-energy GBs (Fig. 3e) [35,45]. The effect of stirring on the degree of crystal agglomeration has been studied previously: Arnberg et al. [46] found that the fraction of low-energy GBs in castings solidifying with electromagnetic stirring is significantly higher than in castings solidifying without stirring. They suggested that shear stresses applied to the semi-solid alloy, e.g. by electromagnetic stirring, increase the crystal collision frequency and therefore promote agglomeration of crystals. Even though primary globules in Thixomolding and ESCs in HPDC are similar in some aspects, there are a number of important differences. First, primary globules are intentionally made and the process conditions, e.g. temperature and stirring parameters, are carefully monitored, while the formation of ESCs is poorly controlled. Second, the shearing time of primary globules in Thixomolding (typically several minutes [19]) is significantly longer than that for ESCs in HPDC (less than tenth of a second during die filling and just several seconds during intensification). Thus, the number of collisions is expected to be significantly higher between globules than ESCs. These differences make the formation of GBs between ESCs significantly different from the formation of GBs between primary globules: agglomeration of primary globules, evidenced in the EBSD mapping results such as in Fig. 3e, was not observed for ESCs in this study. This suggests that external mechanical stresses acting on ESCs during die filling and intensification do not cause agglomeration of ESCs to occur frequently. The combination of a low fraction of ESCs and short period of shear
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must have led to fewer collisions and therefore a low number of ESC–ESC GBs in this study. 4.2. Crystal bending Another possible reason for the presence of a high fraction of LAGBs in the microstructures in this study is crystal bending during solidification. Doherty [47] has suggested that primary crystals can be bent by either mechanical stresses or changes in the growth directions. When mechanical stresses act on solidifying crystals, arising from forced convection [48], solidification contraction [47], gravitational force [7], etc., the crystals can plastically deform leading to bending of their arms. The preferred growth directions of dendrites in certain alloys are known to vary either abruptly or continuously, e.g. primary dendrite growth directions in Al–Zn alloys exhibit continuous changes of direction from h1 0 0i to h1 1 0i over a range of alloy composition [49]. In this case, a change in alloy composition ahead of the solidification front, e.g. due to convection, may lead a change in growth direction leading to crystal bending. If a misorientation between adjacent regions in a bent dendrite arm is within 0–15° (LAGBs), the bent arm still exists. On the other hand, the arm is broken if misorientations exceeds 15° (HAGBs) because cgb > 2cs/l, causing wetting of the liquid at the GB and separation of the crystals. In columnar crystals, crystal bending can be evidenced by in-grain misorientations when sectioned through the longitudinal section of the columnar dendrite. If sectioned through the plane transverse to the primary trunk, dendrite bending can be evidenced by LAGBs between primary and secondary arms. In equiaxed crystals, the growth direction is almost isotropic. If dendrite bending exists in equiaxed structures, it can be observed as either in-grain misorientations or LAGBs. In this study, the primary crystals in the “statically cooled” samples and in-cavity solidified crystals in HPDC and Thixomolding contain only relatively low (<1°) in-grain misorientations (e.g. Fig. 6a). This suggests that the effect of dendrite bending is insignificant or too small to be detected by EBSD. The reason is most likely the compact morphology of the equiaxed crystals in this study. When shear stresses, originating from solidification shrinkage, gravity, external mechanical stresses, etc., act on compact-shaped equiaxed crystals, the crystals may not be able to deform as significantly as in highly branched crystals. ESCs in HPDC and primary globules in Thixomolding are both subjected to more external mechanical stresses than in-cavity solidified grains, particularly at the gate during die filling. The significant in-grain misorientations for the ESCs shown in Fig. 6b suggest that shearing has caused bending of the ESCs. This could be a result of the higher bending moment acting on the longer dendrite arms of ESCs compared to those acting on the more compact globules in the Thixomolded sample and the small, compact incavity solidifying crystals.
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5. Conclusions Microstructural characterization of AlSi4MgMn, AM50 and AZ91 specimens produced by HPDC, Thixomolding and a “static cooling” technique has been performed using EBSD. The study has focused on the influence of external mechanical stresses on agglomeration and bending of equiaxed crystals during solidification. The following conclusions can be drawn. 1. In all samples, the fraction of low-energy GBs is higher than the theoretical prediction for a randomly oriented microstructure. This is attributed to the attractive force between solidifying crystals which have a misorientation relationship favoring the formation of a low-energy GB. 2. In HPDC and Thixomolded samples, the fraction of low-energy GBs among in-cavity solidified grains is significantly higher than in “statically cooled” specimens. This is attributed to the increased number of crystal collisions during HPDC and Thixomolding, which promotes agglomeration of favorably oriented crystals. 3. The distribution of grain misorientations among in-cavity solidified grains in HPDC and Thixomolded specimens are comparable. 4. In-grain misorientations are not significant in primary crystals in all types of specimens, except ESCs in HPDC. This suggests that bending of solidifying crystals is not significant in equiaxed crystals with compact morphology. In contrast, bending was observed in ESCs, which is attributed to the increased bending moment acting on dendrite arms in crystals with a more branched morphology.
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Acknowledgments
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The authors acknowledge the support from the CAST Cooperative Research Centre and the Australian Research Council’s Centre of Excellence for Design in Light Metals. The technical, scientific and financial assistance from the Australian Microscopy and Microanalysis Research Facility (AMMRF) are also acknowledged. We thank Dr. A. Lohmu¨ller (NMF GmbH, Nu¨rnberg, Germany) for supplying the ThixomoldedÒ castings. SO gratefully acknowledges the financial support provided by the Royal Thai Government and a UQ Graduate School Research Travel Grant (GSRTG). SO thanks Drs. Peter Hines and Gwe´nae¨lle Proust (EMU, The University of Sydney) for hospitality and support, and Drs. Tim Williams and Peter Miller (MCEM, Monash University) for the help on preliminary EBSD results.
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