Microstructural evolution of equal channel angular pressed AZ91D magnesium alloy during partial remelting T. J. Chen*, Y. Ma, Y. D. Li, G. X. Lu and Y. Hao In the present work, the authors investigated the microstructural evolution process of equal channel angular pressed (ECAPed) AZ91D magnesium alloy during partial remelting, and the effects of pressing pass, pressing route and heating temperature on the semisolid microstructures. The results indicated that the microstructure evolution could be divided into four steps: the initial coarsening due to the dissolution of interdendritic eutectics, structure separation resulting from the melting of the residual eutectic and the penetration of the firstly formed liquid into the recrystallised boundaries, spheroidisation due to the partial melting of primary particles and final coarsening attributed to Ostwald ripening. Correspondingly, a series of phase transformations occurred in turn: bRa, azbRL and aRL. The variation of the primary particles with heating time obeyed the Lifshitz, Slyozov and Wagner law, D 3ðtÞ 2D 3ð0Þ 5Kt, after the semisolid system was up to liquid–solid equilibrium state. Increasing the heating temperature was beneficial for obtaining an ideal semisolid microstructure because of the decreased tendency of the primary particles to merge. With the increase in pressing pass, the size of the primary particles decreased and their morphology tended to be more spheroidal. Simultaneously, the amount of liquid phase slightly increased because an increased amount of structure melted due to the increased energy stored in the alloy. At a given pressing pass, the semisolid microstructure of the alloy processed by route BC was quite ideal for thixoforming while that of the alloy processed by route A was not completely suitable. In addition, because of the difference in the stored energy, the liquid amount of the former alloy was obviously larger than that of the latter alloy. Keywords: ECAP, Semisolid microstructure, AZ91D alloy, Partial remelting, Thixoforming
Introduction Magnesium alloys are attractive for lightweight structural applications in the transportation industry because of their low density and high specific strength and stiffness.1 However, AZ91D alloy, one of the most commonly used magnesium alloys, suffers from the challenge in meeting the requirements of strength, ductility, fatigue and creep resistance. In order to improve these properties, thixoforming is a promising way by decreasing grain size and shrinkage porosity.2 For thixoforming technology, the key procedure is the production of semisolid ingots with spheroidal primary particles uniformly suspended in liquid phase.2,3 There have been extensive investigations into the microstructural evolution of aluminium and magnesium alloys produced by magnetohydrodynamic or mechanical
Key Laboratory of Gansu Advanced Nonferrous Materials, Lanzhou University of Technology, Lanzhou, China *Corresponding author, email
[email protected]
ß 2010 Institute of Materials, Minerals and Mining Published by Maney on behalf of the Institute Received 8 September 2008; accepted 25 February 2009 DOI 10.1179/174328409X428873
stirring, spray casting, grain refining using refinement or chilling, near liquidus pouring and strain induced melt activation (SIMA).4–8 Alternatively, SIMA process produces the desired structures by deformation and a following heat treatment in the mushy zone and is quite suitable for magnesium alloys because this technique excludes the melting procedure and thus decreases or avoids oxidation.6,7,9–12 Unfortunately, the shape and dimensions of the starting ingots must change after deformation (compressing or rolling), and the strain in the deformed ingots is not quite uniform for the ingots with slightly large dimensions. In addition, it is difficult for magnesium alloys to obtain severe deformation in strain induced step because of the limited number of slip system due to their close packed hexagonal crystal structure. Therefore, a new strain induced procedure is needed in order to obtain severe deformation. It is well known that equal channel angular pressing (ECAP) is a promising technique for obtaining ultrafine grained bulk materials through severe plastic deformation. During ECAP, a billet is pressed through a die that consists of two channels with equal cross-section,
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intersecting at an angle.13,14 Since the cross-sectional shape of billet remains nearly the same, ECAP can be repeated for many passes and very large strain can be introduced in the alloys without cracking. However, investigations on commercial and engineered applications of ECAP are rarely reported, and most of the available studies have been focused on microstructures, mechanical properties and texture of aluminium or magnesium alloys.15–19 Jiang and co-workers used this technology to fabricate semisolid magnesium alloy ingots for thixoforming. Their investigations indicated that a semisolid microstructure with small and spheroidal primary particles could be produced if an equal channel angular pressed (ECAPed) AZ91D alloy ingot was partially remelted.20,21 Recently, Ashouri and coauthors22 also studied the semisolid microstructure of ECAPed A356 aluminium alloy. However, compared with the other technologies for producing semisolid ingots, the investigations about this method are very scarce, and only the above three papers can be found.20–22 Therefore, it is necessary to pay more attention to verify some detailed information, such as the microstructural evolution before liquid formation and phase transformations during partial remelting. Therefore, in the present work, the microstructural evolution process of ECAPed AZ91D alloy and phase transformations during partial remelting are mainly investigated. The effects of processing parameters, such as ECAP pass, ECAP route and heating temperature on the semisolid microstructure are also discussed.
Experimental The material used in this work is commercial AZ91D alloy, and its composition is 9?2Al–0?85Zn–0?35Mn– 0?07Si–0?02Cu–0?002Ni–0?005Fe (wt-%), with a balance of Mg. A quantity of AZ91D alloy was remelted in an electric resistance furnace, degassed by C2Cl6 and poured at 720uC into a permanent mould to form rods with 18 mm diameter and 130 mm length. These rods were machined to small specimens with 9?8 mm diameter and 100 mm length for ECAP. The ECAP die has two equal cross-section channels with a diameter of 10 mm. The intersecting angle between the two channels is 120u, and the angle of the outer intersection arc is 30u. The specimens were heated at 250uC for 1 h in a muffle furnace and then ECAPed. MoS2 was used as lubricant. The ECAPed specimens were cut into small specimens with 6 mm height (along the direction perpendicular to the specimens’ axes) as the starting specimens for partial remelting. The specimens ECAPed for four passes by route BC were heated at 570 or 580uC for different durations, ranging from 2 to 120 min, to investigate on the microstructural evolution with the heating time. Some specimens also ECAPed for four passes by route BC were heated for 10 min at different temperatures, 550, 560, 570 and 580uC, to study the effect of heating temperature on semisolid microstructures. Some specimens were ECAPed for different passes by route BC and then heated at 580uC for 10 min in order to study the effect of the number of ECAP passes on semisolid microstructure. The other specimens were ECAPed for four passes by different routes, such as A, BA, BC and C, heated for 10 min at 580uC to investigate the effect of the ECAP route on semisolid microstructure.
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All of the heated specimens were water quenched quickly. A cross-section (perpendicular to the specimens’ axes) of each specimen was ground and polished. Metallographic observations were firstly carried out using backscattered imaging system S-530 or JSM 6700F scanning electron microscope (SEM), and compositions of phases were examined by an energy dispersive spectroscope (EDS) equipped in S-530 SEM. Then, the specimens were etched by 2 vol.-% HNO3 aqueous solution and observed on a Mef-3 optical microscope. To determine the variations of primary particle size and shape factor during partial remelting, the microstructures of the specimens heated for different durations at 580uC were quantitatively examined. The areas Ai and perimeters Pi of each particle were obtained. The average particle size D was calculated from the formula hX i D~ 2(Ai =p)1=2 =N (1) where N was the total grain numbers in each image. The shape factor F was calculated from the formula X (2) F ~( P2i =4pAi )=N if the particles are perfectly spheroidal, this shape factor has a value of 1; it is larger for less spheroidal particles.23 On each sample, three images with magnification of 100 times were analysed. In order to identify the phase transformations during partial remelting, analyses were carried out on the flat surfaces of the specimens heated for different durations at 570uC using a D8 advance X-ray diffractometer (XRD). It must be noted that the XRD analyses should be carried out immediately after water quenching to avoid natural aging.
Results and discussion Microstructural evolution of ECAPed AZ91D alloy during partial remelting Figure 1a shows the microstructure of the AZ91D alloy ECAPed for four passes in route BC. It can be seen that its microstructure was only little deformed, and compared with the as cast microstructure shown in Fig. 1b, the difference in their microstructural morphologies is very small. This is because the elongated primary dendrites formed during the first and second passes were compressed back to some extent during third and fourth passes along the reverse deformation direction of the first and second passes respectively.24 However, it can be expected that a large strain already has stored in the alloy although its microstructure little changed after being ECAPed for four passes. Figure 2 presents the microstructures of the alloy (ECAPed for four passes in route BC) heated for different durations at 580uC. It can be found that the microstructure turned into large conglomerated structure, and the original deformed dendrite morphology almost disappeared when heated for 2 min (Fig. 2a). This change was seemly due to the merging or coarsening of secondary dendritic arms. In addition, liquid phase started to form (marked by arrows A), and in some local zones, small spheroidal particles presented (marked by arrows B). After 5 min, the large conglomerated structure separated into individual polygonal particles (Fig. 2b). This implied that recrystallisation
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Microstructural evolution of ECAPed AZ91D magnesium alloy
1 Optical micrographs of a AZ91D alloy ECAPed for four passes by route BC and b as cast AZ91D alloy
2 Optical micrographs of AZ91D alloy ECAPed for four passes by route BC followed by heating at 580uC for a 2 min, b 5 min, c 7 min, d 10 min, e 15 min and f 120 min
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3 Variation of temperature in specimen with heating duration heated at 580uC
occurred during this stage, and the firstly formed liquid phase penetrated along the boundaries with high angle that resulted from the recrystallisation, resulting in the formation of the polygonal particles.25,26 As the time was further prolonged, part of the edges of the polygonal particles melted due to the temperature rise in the specimen and the relatively low melting point of the edges,27 which led to the increase in liquid phase amount and spheroidisation of the polygonal particles (Fig. 2c). However, the particle size did not decreased, and some particles singularly coarsened (comparing Fig. 2b and c). This is attributed to the coarsening or merging of some neighbouring particles because these particles might have had a perfectly matching crystallographic orientation and the distance between them was very small.6,28 That is to say that the coarsening and separation were in a competition condition during this stage. When the time was over 7 min, the decrease in size from the melting of the part of the particles was obviously larger than the increase in size from coalescence. This made their size decrease, and the shape become spheroidal. Simultaneously, the amount of liquid phase significantly increased. Thus, a semisolid microstructure with relatively small and spheroidal particles was obtained when heated for 10 min (Fig. 2d). After this, it can be found that the liquid amount basically did not change, and the particles gradually grew (Fig. 2e, f). The temperature in the specimen also remained steady at 580uC (Fig. 3). These implied that the semisolid system reached at the final solid–liquid equilibrium state, and thus the liquid amount did not change. In addition, the distance between the particles increased due to the increase in the liquid amount. Therefore, the coarsening from coalescence further weakened and that from Ostwald ripening became obvious,29 i.e. the particles gradually grew up due to the Ostwald ripening with the further increase in heating time. Therefore, it can be concluded that the microstructural evolution of the ECAPed AZ91D alloy during partial remelting can be divided into four stages: the initial coarsening, structure separation, spheroidisation and final coarsening. As discussed above, the deformed primary grains had evolved into small and spheroidal particles when heated
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4 Variations of primary particle size and shape factor with heating time after semisolid system has reached equilibrium state
for 10 min at 580uC, and then, they coarsened due to Ostwald ripening. Figure 4 shows the variations of the primary particle size and shape factor after being heated for 10 min and longer times at 580uC. It can be clearly seen that the particle size continuously increased with increasing heating time, while the shape factor sharply decreased during period from 10 to 30 min and then slowly decreased. Ostwald ripening is a process in which a system lowers its energy by reducing its interfacial area through the dissolution of small particles and the growth of large particles.28,30 The variation of particle size should obey the formula, D3(t) {D3(0) ~Kt, in the steady equilibrium state, where D(t) is the average particle size (diameter) at time t, D(0) is the initial particle size and K is the coarsening rate constant for a system. The present result is shown in Fig. 5, where D(0) is the average primary particle size at 10 min (taking 10 min heating as the holding time, t50 min). It can be seen that the result obeys the formula well, which demonstrates that the particle coarsening resulted from Ostwald ripening after being heated for 10 min. In addition, Ostwald ripening also includes the dissolution of edges and corners of the particles and the subsequent reprecipitation at sunken zones,4,6 which results in the particles becoming more spheroidal. Therefore, the shape factor sharply decreased during the period from 10 to 30 min
5 Cube of primary particle size versus heating time, taking 10 min as starting time, t50 min
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6 Optical micrographs of as cast AZ91D alloy heated for 15 min at 580uC
(Fig. 4). As the heating time was further prolonged, both the number and curvature of edges and corners of the particles decreased. Thus, the variation of shape factor resulting from Ostwald ripening decreased. From Figs. 2d and 4, it can be seen that a semisolid microstructure with small (an average size of 51 mm) and spheroidal (a shape factor of 1?53) primary particles can be obtained when the ECAPed AZ91D alloy was heated for 10 min at 580uC. However, for the as cast AZ91D alloy, the solid primary phase was very large and irregular, as shown in Fig. 6. Therefore, it can be concluded that the ECAPed followed by partial remelting is an effective method to produce semisolid ingot for thixoforming.
Phase transformations during partial remelting In fact, the microstructural evolution during partial remelting essentially results from some phase transformations. Therefore, the microstructural evolution process can be further clarified by studying the phase transformations occurring during heating. In the present work, the phase transformations were investigated using SEM observation, XRD and EDS analysis. It should be noted that, from Figs. 1a and 2a shown above, it can be seen that the microstructure change was very large during heating from 0 to 2 min, and some details were not obviously presented due to the rapid evolution speed. Therefore, in this section, the microstructures heated at 570uC were observed by SEM. Figure 7 shows SEM micrographs of the ECA pressed AZ91D alloy heated for different times at 570uC. It can be seen that the interdendritic eutectic b phase of the as ECA pressed alloy is not a continuous distribution (Fig. 7a). After being heated for 2 min, the discontinuously distributed b phase almost disappeared, and only some liquid phase can be distinguished from the uniformly grey matrix (Fig. 7b). It is well known that AZ91D alloy is a single phase alloy, and only one Al rich a phase forms at equilibrium solidification condition. However, in practice, including the present work, the solidification condition is not an equilibrium state, so some eutectics always form in interdendritic regions. During the initial stage of partial remelting, the eutectics, especially the Al rich b phase, dissolve into the primary Mg rich a grains because of the temperature
Microstructural evolution of ECAPed AZ91D magnesium alloy
rise in the specimen. This results in the decrease in interdendritic eutectic b phase and the merging of dendritic arms (Fig. 7b). The dissolution of the b phase leads to an increase in the Al content in the primary a phase (Table 1). In addition, this also results in the decrease in b phase and the increase in a phase (comparing the diffraction intensities of the a and b phases of the as ECAPed and heated alloys shown in Fig. 8). This process can be described as the phase transformation of bRa. Because the temperature rise in the specimen was quite rapid (Fig. 3), there was not enough time for the b phase to completely dissolve into the a phase, and then the residual b phase began to melt through the reverse eutectic reaction of azbRL and liquid phase formed (Fig. 7b). As the heating time was further prolonged, the transformation of azbRL further proceeded, which led to the liquid phase increase and the coarsened primary grain separation (Fig. 7c). Simultaneously, the deformed primary grains recrystallised during heating, and small size polygonal grains with subgrain boundaries formed. The firstly formed liquid phase then penetrated along these subgrain boundaries. The arrows marked in Fig. 7c indicate the penetration orientations. Furthermore, the recrystallisation boundaries afforded prompt paths for the dissolution of b phase, and the b phase must agglomerate at these paths. Subsequently, this b phase melted and formed thin and long melted boundaries (Fig. 7c). It can be expected that some b phase could be entrapped within the primary grains due to the mergence of dendritic arms and also melted, forming the liquid pools within the particles (Fig. 7c). In addition, it is impossible that the b phase dissolved into the a grains uniformly distributed within the a grains in a short time. These small size Al rich particles then melted, also forming the liquid pools. During the transformation azbRL, the transformation bRa should also be very active within the period from 2 to 4 min because the amount of a phase continuously increased at this stage (Fig. 8). It is because of these two transformations that the b phase further decreased and even disappeared when heated for 4 min (Fig. 8). Furthermore, it can be expected that the firstly formed Al rich liquid phase would dissolve towards the a phase due to the increase in Al solubility in the a phase from the temperature rise, together with the transformation of bRa, which resulted in the further increase in Al content in the a phase (Table 1). Although both the transformations of azbRL and bRa all operated during the period from 2 to 4 min, the bRa should dominate because the amount of a phase significantly increased (Fig. 8) and the Al content in the eutectics (in fact, the values shown in Table 1 about eutectics were the compositions of eutectic b phase) or liquid phase almost did not change (Table 1). As the heating was prolonged to 7 min, the structure gradually separated into individual polygonal particles due to the penetration and boundary melting driven by the transformation of azbRL (Fig. 7d). Simultaneously, due to the melting of some Mg rich a phase through this transformation, the Al content in the liquid phase sharply decreased (Table 1) and the a phase amount also decreased (Fig. 8). After being water quenched, the liquid phase again solidified into the a and b phases, but the solidification rate during water
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7 Scanning electron micrographs of AZ91D alloy ECAPed for four passes by route BC followed by heating at 570uC for a 0 min (as ECAPed state), b 2 min, c 4 min, d 7 min and e 15 min
quenching was higher than that under the permanent mould casting. It is well known that the more rapid the solidification rate, the more the eutectic amount for the hypoeutectic alloys. Thus, the amount of the eutectic b phase formed under the water quenched condition was more than that formed under the permanent mould casting condition even though the liquid phase amount was the same. That is to say that the more the liquid amount, the more the resulting eutectic b phase. Therefore, the b phase amount increased with increasing heating duration (Fig. 8). As shown in Fig. 2a, a microstructure with individual polygonal particles was obtained when heated for 5 min at 580uC. However, at 570uC, such microstructure was not formed until heated for 7 min. This implied
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that increasing the heating temperature accelerated the speed of microstructural evolution. During heating from 7 to 15 min, part of the primary a phase melted, i.e. the transformation of aRL occurred because of the temperature rise in the specimen. Therefore, both the a phase amount and the Al content in the liquid phase decreased (Table 1 and Fig. 8). Because the melting point of the edges and corners of the polygonal primary particles was relatively low due to their larger curvatures,27 it can be expected that these sites should be preferentially melted, which resulted in the significant increase of the liquid phase and also made the particles spheroidise (Fig. 7e). Furthermore, it should be noted that, according to the Mg–Al binary
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8 X-ray diffractograms of ECAPed AZ91 alloy heated at 570uC for different times and then quenched
phase diagram,31 the Al solubility in the a phase decreases with the temperature increases within the semisolid temperature interval. This is the reason why the Al content in the a particles decreased with the increase in the heating time (Table 1). After being heated for 15 min, the system was up to the solid–liquid equilibrium state. The fraction of liquid or solid was maintained constant, and Ostwald ripening of the dissolution–reprecipitation mechanism dominated. Although the transformations of aRL and LRa were active, the net reaction should be negligible. Therefore, it can be considered that no phase transformations occurred after being heated for 15 min. The fact that the compositions of the a particles and liquid phase and the amounts of the a and b phase did not change demonstrates this standpoint (Table 1 and Fig. 8). From the above discussion, it can be concluded that the microstructural evolution during partial remelting can be well interpreted by the phase transformation process. This process can be divided into four stages, bRa, azbRL, aRL and no phase transformations, which just correspond to the four stages of the Table 1 Chemical compositions of different structures of ECAPed AZ91D alloy during partial remelting at 570uC obtained from EDS analysis Composition, wt-% Heating time, min Structure 0 2 4 7 10 15 20
Mg
Primary a phase 92.22 Eutectics 72.96 Primary a phase 88.78 Eutectics or liquid phase 85.68 Primary a phase 88.38 Liquid phase 83.52 Primary a phase 87.26 Liquid phase 89.58 Primary a phase 83.52 Liquid phase 89.73 Primary a phase 84.28 Liquid phase 92.99 Primary a phase 85.16 Liquid phase 92.12
Al
Zn
5.80 22.47 8.87 12.83 9.34 12.31 11.30 9.15 12.31 8.77 12.25 6.63 12.23 6.60
1.98 3.77 1.21 2.65 1.97 3.64 1.07 1.16 3.64 1.50 2.83 1.18 1.07 1.28
Microstructural evolution of ECAPed AZ91D magnesium alloy
9 Relative X-ray diffraction intensity changes of a(1011) and b of ECAPed AZ91D alloy during partial remelting at 570uC
microstructural evolution, the initial coarsening, structure separation, spheroidisation and final coarsening respectively. According to the changes of the amounts of the a and b phases, the variation of phase transformation with the heating time at 570uC is well depicted by Fig. 9, in which the four stages were marked I–IV respectively.
Effect of heating temperature on semisolid microstructure Figure 10 shows the microstructures of the AZ91D alloy ECAPed for four passes in route BC followed by heating for 10 min at different temperatures. It can be seen that the primary particle size gradually decreased and the shape changed from polygonal to spheroidal as the temperature rose. At the same time, the liquid amount increased. When heated at a relatively lower temperature, the distance between the neighbouring particles was smaller due to the smaller liquid amount. Therefore, the possibility of the particle coalescence should be larger, which resulted in the larger particle size. Furthermore, the needed amount of liquid phase was lower at the lower temperature, so the melted part of the primary particles, such as the edges and corners, was less, which also resulted in larger particle size and especially the irregular shape. It should be noted that the larger possibility of the neighbouring particle merging was the other reason for the irregular shape. As the temperature elevated, the corners and edges of the particles would melt, the distance between particles increased due to the increase of liquid phase and thus the probability of the merging decreased. Therefore, the particle size and shape gradually became small and spheroidal respectively. That is to say that increasing the heating temperature is beneficial for preparing semisolid ingot. However, too high temperature must generate too high liquid fraction, and the advantages of semisolid forming will be lost because this condition is more similar to that of liquid casting. As described in the section on ‘Phase transformations during partial remelting’, the microstructural evolution speed at a higher heating temperature was higher than that at a lower temperature. Moreover, the
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10 Optical micrographs of AZ91D alloys ECAPed for four passes in route BC after being heated for 10 min at a 550uC, b 560uC, c 570uC and d 580uC
microstructural evolution essentially resulted from the phase transformation process. In view of this standpoint, it can be supposed that higher temperature leads to more rapid phase transformation and thus the more rapid microstructural evolution speed.
Effect of pressing pass on semisolid microstructure Figure 11 presents the microstructures of the AZ91D alloy ECAPed for different passes in route BC followed by heating for 10 min at 580uC. Together with the microstructure of the AZ91D alloy ECAPed for four passes shown in Fig. 10d, it can be found, similar to the variations with increasing heating temperature discussed in the section on‘Effect of heating temperature on semisolid microstructure’, that the primary particle size decreased, the shape became more spheroidal and the liquid phase amount increased as the pass increased. For route BC, the equivalent strain eeq generated in the specimen can be calculated by the formula eeq52n/3, where n is the pressing pass.32 After being ECAPed for one to four passes, the corresponding strain was about 0?667, 1?33, 2 and 2?67 respectively. Namely, the strain generated in the specimens increased as the pass increased. Therefore, the size of polygonal primary grains resulted from the recrystallisation during partial remelting decreased with increasing the pass. Thus, the size of final primary particles originated from the polygonal grains also decreased. In addition, it can be expected that the smaller the size of the polygonal particles, the easier they become spheroidal. In other
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words, the smaller the polygonal particle size, the more spheroidal are the final particles. Considering from thermodynamic aspect, the energy stored in the specimens increased as the pass increased. Therefore, the speed of the microstructural evolution should increase, or the residual eutectics and the primary grains would easier melt during partial remelting with increasing the pressing pass. Therefore, it can be easily understood that the liquid phase amount increased with the increase of pressing pass.
Effect of pressing route on semisolid microstructure Figure 12 shows the microstructures of the AZ91D alloy ECAPed for four passes in different routes followed by heating for 10 min at 580uC. It can be seen that the microstructure of the alloy ECAPed in route A, similar to that of the alloy ECAPed for one pass in route BC shown in Fig. 11a, is composed of big, irregular and interconnected particles (Fig. 12a). For the alloy pressed in route BA, its microstructure consists of small and individual particles, and most of them are spheroidal (Fig. 12b). The microstructure of the alloy pressed in route C is composed of individual particles, but all of them are almost in polygonal shape and have relatively large size (Fig. 12c). As described above, for the alloy pressed in route BC, the primary particles are small, uniform and spheroidal (Fig. 10d). According to the deformation mechanisms during ECAP, shear always carries out in one shear plane for
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Microstructural evolution of ECAPed AZ91D magnesium alloy
11 Optical micrographs of AZ91D alloys ECAPed for a one pass, b two passes and c three passes in route BC and then heated for 10 min at 580uC
12 Optical micrographs of AZ91D alloys ECAPed for four passes in routes a A, b BA and c C after being heated for 10 min at 580uC
the alloy pressed for four passes in route A.33 The operated shear systems are few, and the number of recrystallised grain boundaries with large angle is also small. So, the size of the recrystallised grains is large and thus that of the final primary particles in the semisolid microstructure is also large and irregular. Therefore, it can be concluded that the final primary particle size and shape are determined by the number of the operated shear systems during ECAP. For routes BA, C and BC, shear occurred in two planes, one and four planes respectively.34 Thus, the primary particles in the semisolid microstructure produced by route BC should be the most spheroidal and the smallest, and those
prepared by route BA take the second place. Although the shear for routes A and C all operate in one plane, the shear orientations between two passes reversed each other for route C, and equiaxed grains is easier to obtain.34 Therefore, the final primary particles for route C should be smaller and more regular than those for route A. Similarly, considering from thermodynamic aspect, it can be supposed that the energy stored in the specimen should be higher if the grain refining effect is better. In view of this standpoint, it can be easily understood that the liquid phase amount in the semisolid microstructures decreased in turn of routes BC, BA, C and A (comparing Figs. 11a and 12).
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Conclusions 1. Equal channel angular pressing followed by partial remelting is an effective method to produce semisolid AZ91D alloy ingot for thixoforming. A semisolid microstructure with 51 mm spheroidal particles can be obtained if the AZ91D alloy is pressed for four passes in route BC and then heated for 10 min at 580uC. 2. During partial remelting, the microstructure evolution of the ECAPed AZ91D alloy can be divided into four stages, the initial coarsening due to the dissolution of interdendritic, structure separation resulting from the melting of the residual eutectic and the penetration of the firstly formed liquid into the recrystallised boundaries, spheroidisation due to the partial melting of primary particles and final coarsening attributed to Ostwald ripening, which correspond to the phase transformations of bRa, azbRL, aRL and no phase transformations respectively. 3. Proper increase in heating temperature is beneficial to prepare semisolid microstructure with small and spheroidal primary particles. 4. With the increase in pressing pass, the primary particle size decreased and its shape became more spheroidal. Simultaneously, the amount of liquid phase increased. 5. The semisolid microstructures produced by different routes changed from good to bad in turn of BC, BA, C and A. The liquid phase amount decreased in this sequence.
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