Flow localization during severe plastic deformation of AZ81 magnesium alloy: Micro-shear banding phenomenon

Flow localization during severe plastic deformation of AZ81 magnesium alloy: Micro-shear banding phenomenon

Author's Accepted Manuscript Flow localization during severe plastic deformation of AZ81 magnesium alloy: Microshear banding phenomenon P. Changizian...

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Author's Accepted Manuscript

Flow localization during severe plastic deformation of AZ81 magnesium alloy: Microshear banding phenomenon P. Changizian, A. Zarei-Hanzaki, M. Ghambari, A. Imandoust

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S0921-5093(13)00618-7 http://dx.doi.org/10.1016/j.msea.2013.05.069 MSA29967

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Materials Science & Engineering A

Received date: 29 April 2013 Revised date: 24 May 2013 Accepted date: 27 May 2013 Cite this article as: P. Changizian, A. Zarei-Hanzaki, M. Ghambari, A. Imandoust, Flow localization during severe plastic deformation of AZ81 magnesium alloy: Micro-shear banding phenomenon, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2013.05.069 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Flow localization during severe plastic deformation of AZ81 magnesium alloy: Micro-shear banding phenomenon P. Changizian, A. Zarei-Hanzaki∗, M. Ghambari, A. Imandoust

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Email address: [email protected] (Abbas Zarei-Hanzaki) Tel.: +98-21-44284466, Fax: +98-21-88006076

 

Flow localization during severe plastic deformation of AZ81 magnesium alloy: Micro-shear banding phenomenon P. Changizian, A. Zarei-Hanzaki∗, M. Ghambari, A. Imandoust

Abstract The present study deals with the flow localization and its effects on the deformation mechanism as well as texture characteristics of AZ81 magnesium alloy during severe plastic deformation. Towards this purpose, a set of accumulative back extrusion (ABE) was conducted at temperatures ranging from 300°C to 450 °C. The results indicate that the severe imposed strain tends to be localized in the form of micro-shear band due to the intrinsic anisotropic plasticity of the alloy. This banding evolves from the mechanical twinning which has a great contribution in accommodation of severe plastic strain during initial stages of ABE-processing. In addition, the micro-shear banding provides paths by which the extensive operation of slip systems and high dislocation density promote the initiation of dynamic recrystallization and dynamic precipitation of γ particles. The micro shear banding has also resulted in the deviation from the strong fiber basal texture at higher deformation temperature. This is mainly attributed

                                                                                                                                                                                                

to the reorientation of basal planes parallel to the shear plane and widening of the bands via progression of the recrystallization inside the micro-shear bands. Keywords: Micro-shear banding; Severe plastic deformation; Dynamic recrystallization; Twinning; Texture 1. Introduction The plastic instability or plastic flow localization generally is seen as the jerky flow, Luders bands, deformation bands, and shear bands during straining.

This would lead to

premature fracture of the materials during deformation [1, 2]. In the case of magnesium alloys, shear banding is the main manifestation of strain localization [3]. Accordingly, so many efforts have been directed towards characterizing the associated structures [1, 4, 5] and developing theories to explain and predict the shear banding phenomenon [4]. The proposed models in this regard have been mainly classified into two types, both of which are strongly correlated with the dominant deformation mechanisms. The former, proposed by Ion et al. [6], relies on the fact that the shear bands form as a consequence of rotational dynamic recrystallization at the vicinity of grain boundaries. Once the new recrystallized grains are formed, they tend to cluster and develop large banded areas in a way that deformation concentrates in these regions and produces ductile shear zones. The latter, developed by Couling et al. [7] and Wonsiewicz et al. [8], suggest that the shear bands may evolve from double twinning mechanism in the course of deformation. In line to these efforts, the dominant shear banding mechanism has been systematically investigated on the series of AZ magnesium alloy [4, 9]. It has been found that the initial texture and grain size are the two crucial elements in activating the dominant shear banding mechanism.

                                                                                                                                                                                                

The stronger basal texture and bigger initial grain size would increase the influence of mechanical twinning (so triggering the second proposed mechanism, Couling et al. [7] and Wonsiewicz et al. [8]), while weaker texture and smaller grain size would activate the recrystallization mechanism (i.e., the first proposed mechanism, Ion et. al. [6]). In addition, regarding the mutual effect of shear banding and texture evolution during deformation, the shear banding would provide a path by which the orientation of grains are more favored for basal slip; therefore, a lower stress level is necessitated for further deformation inside the shear zones compared to the matrix [10]. This is so called the texture softening effect. Accordingly, the flow localization, in particular shear banding, could have a remarkable influence on the operation of slip and twinning systems, restoration mechanisms, texture evolution, and vice versa. As a matter of fact, the effect of shear banding is more pronounced during severe plastic deformation, due to the limited active slip systems to accommodate the imposed strain. Up to date, the different severe plastic deformation methods such as equal channel angular press (ECAP) [11, 12], accumulative roll bonding (ARB) [13, 14], and friction stir processing (FSP) [15, 16] have been proposed and successfully applied to refine the microstructure of the magnesium alloys in general, and cast AZ magnesium alloys in special. Consequently it was found that the shear banding play a crucial role in recrystallization behavior and texture evolution of magnesium alloys. The occurrence of dynamic recrystallization along and at the vicinity of shear bands has been reported during equal channel angular extrusion (ECAE) of commercially pure magnesium [17]. It has been also shown that recrystallization within shear bands might lead to stronger texture after severe plastic deformation. However, there are few controversial reports in this regard indicating that the rotational recrystallization, which would result in shear banding, might lead to significant decrease of deformation texture [10, 18].

                                                                                                                                                                                                

Accumulative back extrusion (ABE) method has been recently introduced as a new noble SPD process suitable for mass production [19, 20]. This method has brought many advantages such as lower required load and no need of inter-pass operation compared with the conventional SPD processes [21]. Accumulative back extrusion has been successfully employed to severely deform the low Al content AZ magnesium alloy (AZ31) [22]. Based on the obtained results, significant grain refinement has been achieved and the subsequent mechanical properties have been improved.

Furthermore, the shear banding has been found during low temperature

deformation and outstanding grain refinement has been reported inside the shear bands [1]. In spite of the several conducted researches, a great lack of knowledge is realized on the severe plastic deformation behavior of high-Al containing AZ magnesium alloys during accumulative back extrusion processing and the effect of strain localization on their microstructure and texture evolution. Accordingly, the present work has been conducted to clarify the flow localization characteristics in accumulative back-extruded AZ81 magnesium alloy. This study emphasizes on the effect of the strain localization on the restoration process, precipitation of the second phase, and texture evolution.

2. Experimental procedures The experimental material (AZ81 magnesium alloy) was received in as-cast condition, the chemical composition of which is Mg-8.1Al-0.7Zn-0.18Mn (all in wt.%).

The initial

microstructure of as-cast AZ81 Mg alloy consists of α-Mg matrix, coarse eutectic γ-Mg17Al12 network, and γ-precipitates distributed in grain interior and along the grain boundaries (Figs. 1a and b). The as-cast material was homogenized at 415 ± 2 °C for 24h to dissolve the γ phases formed during casting into the α- phase matrix.

As is seen in Fig. 2 the homogenization

                                                                                                                                                                                                

treatment has resulted in outstanding dissolution of γ-Mg17Al12 phase, thereby producing a supersaturated solid solution.

Only a few γ-eutectic particles are seen along the grain

boundaries. The grain size of solution treated material was estimated to be ~300μm. The accumulative back extrusion was carried out at different temperatures in the range of 300-450 °C using cylindrical workpieces in the sizes of Ф18 mm × H9 mm. The schematic illustration of the process is given in Fig. 3. Any ABE pass (cycle) consists of two steps, including a back extrusion followed by a two-dimensional constrained back pressing. In the first step the workpiece is back extruded into the gap between the inner punch and the die (Figs. 3a and b). In the second step the back-extruded material is forged back by the outer punch (Figs. 3c and d). Consequently, at the end of any cycle (pass) the initial shape of the workpiece is reproduced. Any ABE processing was accomplished by preheating and stabilizing the die at at deformation temperature. The graphite spray was also used to reduce the friction between the workpiece and the tool surfaces. The temperature was controlled using a K-type thermocouple, which was firmly set into the die wall (2 mm deep from the wall surface). The workpieces were first heated up to the deformation temperature and held isothermally for 5 min to achieve a uniform temperature prior to any ABE processing. At the end of the process, the products were cut perpendicular to their longitudinal axis (parallel to the extrusion direction) for further microstructural observations and texture analysis (Fig. 4). The sectioned specimens were mounted using cold curing resin, ground and polished step by step with the final polishing by 0.05μm Al2O3 powders. The polished surface was revealed using acetic-picral etchant (4.2g picric acid, 10 ml acetic acid, 10 ml distilled water, and 70 ml ethanol).

Macro-texture measurements were performed, using a Philips X’Pert

diffractometer furnished with a close Eulerian cradle. The measurements were carried out

                                                                                                                                                                                                

employing CuKα radiation at 50 kV with the sample tilt angle ranging from 0° to 90°. The calculated (0002) pole figures and orientation distribution function (ODF) were obtained using the X’Pert Texture software. Bunge notation of the Euler angles was utilized to calculate the ODF maps.

3. Results and discussion 3.1. Micro-shear banding and deformation mechanism The microstructure of the ABE-processed material at 300 °C (as a representative of the developed microstructure) has been shown in Fig. 5. The remarkable contribution of mechanical twinning to accommodate the imposed severe strain is clearly seen. Outstanding amounts of twin bands, which are typically ended to the original grain boundaries or intersect with other twins and form mutual twinning intersections at the grains interiors, are also recognized. As a matter of fact, the initial coarse grain size of the experimental alloy (due to the large-scale heterogeneity [23]) and also the high percentage of Al content (due to its dragging effect on dislocation movement [24]) provide a suitable condition for mechanical twinning. The grain segmentations due to the twinning mutual intersections play a vital role in accommodation of the severe imposed strain through accumulative back extrusion. These segmentations make the grain consistent with the overall deformation, because the plastic flow can be accommodated more easily at reduced scales rather than a total grain. Despite the extensive occurrence of twining, the severe imposed strain tends to be localized in the form of micro shear band as is shown in Fig. 6. This is well justified considering the intrinsic anisotropic plasticity of the magnesium alloys [25]. In line to this effort, the strain localization has been reported in a

                                                                                                                                                                                                

number of studies during severe plastic deformation of magnesium alloys and several explanations have been put forward to describe the formation of these bands. In the present study, the twinning development is thought to be the main responsible for micro-shear banding. This is described as is follows. The proposed sequence of the micro shear banding with emphasize on the role of mechanical twinning is depicted in Fig. 7.

As is well established, due to the restricted

independent slip systems in Mg alloys at relatively low temperatures, severe plastic deformation would activate the different variants of primary twins (tension and compression twinning systems) in a single grain at the initial stages of the process. This can lead to the formation of several parallel twin bands and mutual intersection of twins in some definite grains (Fig. 7b). Considering the high misorientation induced by tension twins (86⁰ [26]) and compression ones (56⁰ [27]), they act as the strong barriers to dislocation movement within original grains [28]. Therefore, the dislocation piles up at twin boundaries may cause relatively high stress concentration and result in distortion of mechanical twins toward the shear direction (Fig. 7c). This may lead to the strain localization and the formation of micro-shear bands along the twin boundaries (Fig. 7d). It is worth to note that these regions are favorable sites for formation of new dynamically recrystallized grains at higher temperatures as is schematically depicted in Fig. 7e. In fact, the concentration of shear strain within the micro shear bands activates more slip systems compared to the matrix, and subsequently provide a suitable condition for the formation of sub-grains and well defined grains [1, 10]. For further clarification, the final microstructures of the processed specimens at the higher temperatures of 350, 400 and 450 °C has been illustrated in Fig. 8. As is evident, the initiation of the dynamic recrystallization from micro shear bands has resulted in the necklace-

                                                                                                                                                                                                

type or sharply bimodal structures formation in which the grain size distribution is highly inhomogeneous (Figs. 8a and b). According to the Fig. 8b, the initiation of dynamic recrystallization from micro-shear bands has resulted in formation of ultra-fine grains along these bands; hence, significant grain refinement has been achieved inside the micro-shear bands. At higher deformation temperatures of 400 and 450 °C the micro shear bands are widen by increasing the volume fraction of the recrystallized grains within them and consuming the parent grains, so the homogeneity has been increased (Figs. 8c, d and 8e, f).

3.2. Micro-shear banding and dynamic precipitation Having a supersaturated solution treated alloy, the severe imposed strain would lead to the strain induced precipitation of dissolved γ-Mg17A112 in the course of deformation. Figure 9a illustrates the microstructure of the processed specimen at 300 °C and depicts the non-uniform precipitation of fine γ-Mg17Al12 particles (i) along the micro-shear bands and also (ii) at the vicinity of the γ-eutectic phase. The latter results from the high concentration of Al solute atoms near the γ-eutectic phase, and the former is rationalized considering the fact that the localized strain within the micro shear bands facilitates the occurrence of dynamic precipitation. In more details, the high dislocations density in consequence of the shear strain concentration provides a path with high diffusion rate, so the heterogeneous precipitation of γ phase initiating from micro shear bands is expected. In addition, a more homogenous microstructure has been obtained at higher deformation temperatures as is shown in Fig. 9b. In fact, the progress of dynamic recrystallization process in the microstructure eliminates the influence of the strain localization and leads to precipitation of the second phase more uniformly at higher temperatures.

                                                                                                                                                                                                

3.3. Micro-shear banding and texture evolution The initial texture of AZ81 experimental alloy is a typical of as-cast magnesium alloys consisting of a random texture. Fig. 10a and b show the (0002) pole figure and the related orientation distribution function (ODF), respectively, which indicate the texture characteristics of the single pass ABE processed specimen at 300 °C. As is observed the dominant deformation texture is strong (0002) basal fiber texture with the basal planes perpendicular to ND, i.e. the caxes are perpendicular to extrusion direction (ED). The strong basal texture is mainly originated from the operation of basal dislocation slip along with the extension twinning which was shown to be activated from the initial stages of deformation (Fig. 5). The same basal texture has been developed after hot deformation of Mg-Al-Zn alloys, where the twinning and basal slip reported as the dominant deformation mechanisms [29, 30]. The characteristics of dominant texture component for the severely deformed specimens at higher temperatures of 350, 400 and 450 ⁰C have been depicted in Fig. 11. As is clear increasing the deformation temperatures has resulted in development of similar basal deformation texture with c-axis perpendicular to the extrusion direction. However, a significant decrease in the maximum intensity of (0002) fiber texture has been observed by increasing the deformation temperature. It is well clarified that the developed recrystallization texture at these relatively high temperatures is close to the deformation texture, so the fiber texture is expected to be intensified. This discrepancy (the deviation from basal texture along with the occurrence of recrystallization) may be justified relying on the formation of micro-shear bands, which accommodates the large strain during ABE-processing. The microshear banding leads to the zones development with orientation which is different from the deformation orientation. In fact, the grains within the micro-shear bands are orientated parallel to the shear plane in a way that being different from the orientation of the matrix grains induced

                                                                                                                                                                                                

via accumulative back extrusion (that is aligned parallel to the extrusion direction). Accordingly, the dynamic recrystallization inside the shear bands resulted in the formation of quite a few new grains with orientation parallel to the shear plane. Therefore, the widening of micro-shear bands due to the recrystallization progress within them is considered as the main reason for weakening the basal texture in the specimens processed at higher temperatures. This justification is consistent with the reported results of Del Valle et al. [10] and Li [18], who studied the texture evolution of AZ61 magnesium alloy during large strain hot rolling and cast Mg–RE alloy under uniaxial compression, respectively. Consequently, the recrystallization within micro shear bands leads to the formation of remarkable number of grains where the basal planes are aligned toward the shear plane. This would provide the zones with high Schmid factor, so the paths for the further easy basal slip are produced. In contrast, having strong basal texture as a dominant deformation texture, in most of the matrix grains, would make the basal slip unfavorite due to the lower Schmid factor. This preference (easy basal slip within the micro shear bands) enables the system to deform under lower stress. The occurrence of texture softening has been also reported by Barrnet et al. [3] during cold rolling of pure Mg, Mg-0.2Ce and Mg-3Al-1Zn.

4. Conclusion The present study has focused on the flow localization characteristics and its effect on the dynamic recrystallization and precipitation behavior as well as the texture evolution of severely deformed AZ81 magnesium alloy employing the accumulative back extrusion technique. The main points resulted from this work are summarized as is follows.

                                                                                                                                                                                                



The strain localization leads to an extensive slip and cross slip within shear bands, which promotes the initiation of dynamic recrystallization from these bands.



The strain induced precipitation takes place non-uniformly along the micro-shear bands, where the high density of dislocations facilitates the aluminum diffusion in magnesium matrix.



The detected strong fiber basal texture after single pass process is attributed to the great contribution of extensive twinning which accommodates the severe imposed strain.



The deviation from basal texture along with the occurrence of recrystallization at high temperatures has been justified relying on the formation of micro-shear bands, which would lead to develop the zones with orientation which is different from the deformation orientation.

                                                                                                                                                                                                

References [1] S. Fatemi-Varzaneh, A. Zarei-Hanzaki, J. Cabrera, J. Alloys Compd. 509 (2011) 3806-3810. [2] F. Kang, J.T. Wang, Y. Peng, Mater. Sci. Eng. A 487 (2008) 68-73. [3] M. Barnett, M. Nave, C. Bettles, Mater. Sci. Eng. A 386 (2004) 205-211. [4] Z. Zhang, M.-p. Wang, N. Jiang, S. Zhu, J. Alloys Compd. 512 (2012) 73-78. [5] Y. Chen, Q. Wang, H. Roven, M. Karlsen, Y. Yu, M. Liu, J. Hjelen, J. Alloys Compd. 462 (2008) 192-200. [6] S. Ion, F. Humphreys, S. White, Acta Metall. 30 (1982) 1909-1919. [7] S. Couling, J. Pashak, L. Sturkey, Trans. ASM 51 (1959) 94. [8] B. Wonsiewicz, G. Chin, Trans. AIME 239 (1967) 1422. [9] A. Chalay-Amoly, A. Zarei-Hanzaki, P. Changizian, S. Fatemi-Varzaneh, M. Maghsoudi, Mater.

Des. 47 (2013) 820-827. [10] J. Del Valle, M.T. Pérez-Prado, O. Ruano, Mater. Sci. Eng. A 355 (2003) 68-78. [11] J. Del Valle, F. Carreño, O.A. Ruano, Acta Mater. 54 (2006) 4247-4259. [12] M. Eddahbi, J.d. Valle, M.T. Pérez-Prado, O.A. Ruano, Mater. Sci. Eng. A 410 (2005) 308-311. [13] J. Del Valle, M.T. Pérez-Prado, O.A. Ruano, Mater. Sci. Eng. A 410-411 (2005) 353-357. [14] M.T. Pérez-Prado, O. Ruano, Scripta Mater. 51 (2004) 1093-1097. [15] A. Feng, Z. Ma, Acta Mater. 57 (2009) 4248-4260. [16] A.H. Feng, B.L. Xiao, Z.Y. Ma, R.S. Chen, Metall. Mater. Trans. A 40 (2009) 2447-2456. [17] S. Suwas, G. Gottstein, R. Kumar, Mater. Sci. Eng. A 471 (2007) 1-14. [18] L. Li, Mater. Sci. Eng. A 528 (2011) 7178-7185. [19] S. Fatemi-Varzaneh, A. Zarei-Hanzaki, Mater. Sci. Eng. A 504 (2009) 104-106. [20] B. Bazaz, A. Zarei-Hanzaki, S. Fatemi-Varzaneh, Mater. Sci. Eng. A 559 (2013) 595–600.

                                                                                                                                                                                                 [21] N. Haghdadi, A. Zarei-Hanzaki, S. Heshmati-Manesh, H. Abedi, S. Hassas-Irani, Mater. Des. 49 (2013) 878–887. [22] S. Fatemi-Varzaneh, A. Zarei-Hanzaki, Mater. Sci. Eng. A 528 (2011) 1334-1339. [23] E. Karimi, A. Zarei-Hanzaki, M. Pishbin, H. Abedi, P. Changizian, Mater. Des.49 (2013)

173-180. [24] P. Changizian, A. Zarei-Hanzaki, A.A. Roostaei, Mater. Des. 39 (2012) 384–389. [25] J. Koike, Metall. Mater. Trans. A 36 (2005) 1689-1696. [26] M. Barnett, Mater. Sci. Eng. A 464 (2007) 1-7. [27] M. Barnett, Mater. Sci. Eng. A 464 (2007) 8-16. [28] P. Changizian, A. Zarei-Hanzaki, H. Abedi, Mater. Sci. Eng. A 558 (2012) 44–51. [29] K. Hantzsche, J. Bohlen, J. Wendt, K. Kainer, S. Yi, D. Letzig, Scripta Mater. 63 (2010) 725-730. [30] A. Styczynski, C. Hartig, J. Bohlen, D. Letzig, Scripta Mater. 50 (2004) 943-947.

Figures Captions Figure 1. a) The optical, and b) the SEM micrographs of as-cast AZ81 magnesium alloy. Figure 2. The optical micrograph of AZ81 magnesium alloy after homogenization. Figure 3. Schematic presentation of accumulative back extrusion process steps. Figure 4. The schematic illustration of the specific plane in processed workpiece used for microstructural and texture analysis. Figure 5. The final microstructure of the single pass processed material at 300 ⁰C (extensive occurrence of twinning). Figure 6. The final microstructure of the single pass processed specimen at 300 ⁰C (micro-shear banding phenomenon). Figure 7. The schematic depiction of micro-shear band formation and the subsequent dynamic recrystallization.

School of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, Tehran, Iran

Abstract The present study deals with the flow localization and its effects on the deformation mechanism as well as texture characteristics of AZ81 magnesium alloy during severe plastic deformation. Towards this purpose, a set of accumulative back extrusion (ABE) was conducted at temperatures ranging from 300°C to 450 °C. The results indicate that the severe imposed strain tends to be localized in the form of micro-shear band due to the intrinsic anisotropic plasticity of the alloy. This banding evolves from the mechanical twinning which has a great contribution in accommodation of severe plastic strain during initial stages of ABE-processing. In addition, the micro-shear banding provides paths by which the extensive operation of slip systems and high dislocation density promote the initiation of dynamic recrystallization and dynamic precipitation of γ particles. The micro shear banding has also resulted in the deviation from the strong fiber basal texture at higher deformation temperature. This is mainly attributed to the reorientation of basal planes parallel to the shear plane and widening of the bands via progression of the recrystallization inside the micro-shear bands.                                                                                                                                                                                                 

Figure 8. The optical microstructures of the single pass ABE-processed material at different temperatures of (a, b) 350, (c, d) 400, (e, f) 450 ⁰C Figure 9. The SEM micrograph of dynamically precipitated Mg17Al12 particles in the single pass ABE-processed material at different temperatures of a) 300 and b) 450 ⁰C. Figure 10. The texture charactristics of the processed specimen at 300 ⁰C, a) (0002) pole figure, and b) oreintation distribution function (ODF). Figure 11. (0002) pole figures of the specimens processed at different tempertures of a) 350, b) 400, and c) 450⁰C.

Keywords: Micro-shear banding; Severe plastic deformation; Dynamic recrystallization; Twinning; Texture

Figure(s)

(a)

(b) γ-eutectic

γ- precipitates

Fig. 1. a) The optical, and b) the SEM micrographs of as-cast AZ81 magnesium alloy.

Figure(s)

Fig. 2. The optical micrograph of AZ81 magnesium alloy after homogenization.

Figure(s)

(a)

(b)

(c)

(d)

Fig. 3. Schematic presentation of accumulative back extrusion process steps.

Figure(s)

Fig. 4. The schematic illustration of the specific plane in processed workpiece used for microstructural and texture analysis.

Figure(s)

(a)

(b)

Fig. 5. The final microstructure of the single pass processed material at 300 ⁰C (extensive occurrence of twinning).

Figure(s)

(a)

(b)

Micro-shear bands Fig. 6. The final microstructure of the single pass processed specimen at 300 ⁰C (micro-shear banding phenomenon).

Figure(s)

(a)

(b)

(d)

(c)

(e)

Fig. 7. The schematic depiction of micro-shear band formation and the subsequent dynamic recrystallization.

Figure(s)

(b)

(a)

DRX Initiation from micro-shear bands

Micro-shear bands

(c)

(d)

Widened micro-shear band

(e)

(f)

Fig. 8. The optical microstructures of the single pass ABE-processed material at different temperatures of (a, b) 350, (c, d) 400, (e, f) 450 ⁰C.

Figure(s)

(a) γ-eutectic

γ-Mg17Al12 particles

Trace of micro-shear band

(b)

Fig. 9. The SEM micrograph of dynamically precipitated Mg 17Al12 particles in the single pass ABE-processed material at different temperatures of a) 300 and b) 450 ⁰C.

Figure(s)

(a)

ED

Max: 12.5

0

TD

(b) 0

Fig. 10. The texture charactristics of the processed specimen at 300 ⁰C, a) (0002) pole figure, and b) oreintation distribution function (ODF).

Figure(s)

ED

(a)

Max: 11.2

0

TD

(b)

ED

Max: 10.8

(c)

ED

Max: 9.1

0

0

TD

TD

Fig. 11. (0002) pole figures of the specimens processed at different tempertures of a) 350, b) 400, and c) 450 ⁰C.