Ex-situ EBSD analysis of hot deformation behavior and microstructural evolution of Mg–1Al–6Y alloy via uniaxial compression

Materials Science & Engineering A 775 (2020) 138978

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Ex-situ EBSD analysis of hot deformation behavior and microstructural evolution of Mg–1Al–6Y alloy via uniaxial compression Mengna Zhang, Jinhui Wang *, Yunpeng Zhu, Lei Zhang, Peipeng Jin Qinghai Provincial Key Laboratory of New Light Alloys, Qinghai Provincial Engineering Research Center of High-Performance Light Metal Alloys and Forming, Qinghai University, Xining, 810016, PR China

A R T I C L E I N F O

A B S T R A C T

Keywords: Magnesium alloy Texture Work hardening Twinning Dislocation slip

The microstructure and texture evolution of as-extruded Mg-0.96 wt%Al-5.82 wt%Y (denoted as Mg–1Al–6Y) alloy with strain at a temperature of 350 � C and a strain rate of 0.1 s 1 were studied using ex-situ electron backscatter diffraction (EBSD). Moreover, the deformation mechanism was systematically analyzed. The results shown that weak <2-1-11> rare-earth texture and no twinning were observed in the as-extruded Mg–1Al–6Y alloy. The hot compression curve exhibited obvious work-hardening characteristics with a yield strength of 98 MPa. In the initial stage of deformation, {10–12} extension twinning was formed in some grains. As the strain increased, the number of {10–12} extension twinning and texture intensity gradually decreased, whereas the amount of low-angle grain boundaries increased. The Schmid factor (SF) of pyramidal slip remained at ~0.4. In addition, the TEM characterization revealed that as the strain increased, dislocations multiplied, and the pyra­ midal dislocation was always observed. Therefore, the reduction of twinning and activation of pyramidal slip lead to the dominance of slip during the deformation. Hence, work hardening appeared in the compression curve of Mg–1Al–6Y alloy.

1. Introduction Magnesium alloys with low density, high specific strength, and specific stiffness have potential applications in electronics, automotive, aeronautics, and astronautics fields [1,2]. However, because magnesium has the hexagonal close-packed (hcp) crystal structure [3], only two independent slip systems are present on the basal plane. This is insuf­ ficient to accommodate c-axis deformation at room temperature [4,5], while prismatic and pyramidal slip can be activated at a higher tem­ perature [6,7]. In addition, because of low critical resolved shear stress (CRSS), during the compression, the {10–12} extension twin can be easily activated [8]. Therefore, magnesium alloy exhibits low strength and ductility at a high temperature, which significantly limits its broad applications under the high-temperature conditions. A large number of experimental and simulation studies [9–11] have shown that due to the interactions between {10–12} extension twinning and dislocation slips [12], the thermal compression curve usually ex­ hibits two types of characteristics: one is the yielding platform caused by softening of twinning [13,14], another is the “S” curve dominated by work-hardening [15]. Zhou et al. [16] have studied the thermal defor­ mation behavior of AZ31 magnesium alloy at different strain rates and

temperatures, and they found that the deformation mechanism was dominated by twin deformation in the early stage of low-temperature deformation. However, Mg–Al and Mg–Al–Zn alloys exhibited positive strain rate sensitivity and negative temperature effect during high-temperature compression, which shows obvious softening charac­ teristics. The deformation mechanism was mainly dominated by twin­ ning and dynamic recrystallization (DRX) [17–20]. However, although the high-temperature deformation mechanism of Mg–Al alloy was ob­ tained, the poor high-temperature mechanical properties of Mg–Al alloy were still not improved. Recent studies have found that alloying (especially the addition of rare-earth elements) can effectively improve the high-temperature me­ chanical properties of magnesium alloys. Zhang et al. [21–24] found that rare-earth element Y has a high solubility in magnesium alloy and significantly promotes the bonding force between atoms and effectively inhibits dislocation slip. Thus, the strain-hardening rate and high-temperature mechanical properties of magnesium alloys will be improved during the hot deformation. Therefore, the addition of rare-earth element Y probably has a favorable influence on the defor­ mation mechanism of magnesium alloy. Huang et al. [25] have studied the hot deformation behavior of magnesium alloys containing rare-earth

* Corresponding author. E-mail address: [email protected] (J. Wang). https://doi.org/10.1016/j.msea.2020.138978 Received 9 October 2019; Received in revised form 17 January 2020; Accepted 18 January 2020 Available online 20 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.

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element Y. It was found that the addition of Y leads to the formation of <2-1-11> rare-earth texture in the alloy and then effectively improves the compressive strength of alloy. Although a large number of studies have focused on the effects of deformation temperatures and strain rates on the high-temperature mechanical properties of magnesium alloys. However, the deformation behaviors, especially microstructure and texture evolution, of magnesium with rare-earth elements at high-temperature, has not been studied in depth. Therefore, in present study, the texture evolution, activation behavior of twinning and slip system, and the characteristics of continuous work hardening of as-extruded Mg–1Al–6Y alloy during high-temperature deformation were studied.

obtained at a higher magnification and the corresponding selected-area electron diffraction (SAED) pattern are shown in Fig. 1(b-1), which further demonstrated that the precipitate was the Al2Y phase, and the size and the volume fraction of Al2Y phase were very small. However, only 2 low-intensity Al2Y phase diffraction peaks were observed in the XRD pattern, suggesting a small content of Al2Y phase. Fig. 1 (c) shows the inverse pole figure (IPF) map of as-extruded Mg–1Al–6Y alloy. The alloy mainly consists of fine recrystallized grains and coarse unrec­ rystallized grains, surrounded by a number of smaller grains. The average grain size was ~13 μm, and a weakly preferred orientation was observed. Fig. 1(d) shows that the as-extruded Mg–1Al–6Y alloy had a weak <2-1-11> rare-earth texture with an intensity of 6.26. In addition, the c-axis of most grains tilted about 45� from ED to transverse direction (TD), which was not a typical basal fiber texture.

2. Experimental procedure In present study, a Mg–1Al–6Y (wt.%) alloy ingot was prepared from industrial pure Mg (99.95 wt%), pure Al (99.99 wt%), and Mg–Y master alloy (30 wt%). The preheated alloy was melted in an electric resistance furnace and kept at 750 � C for 20 min, and then the liquid alloy was poured into a steel mold with a size of φ 500 mm � 1000 mm when cooled down to 730 � C. To prevent melt oxidation or combustion during smelting, a mixture of 99 %CO2 and 1 %SF6 was used as a protective gas. The ingot was preheated at 500 � C for 1.5 h and hot-extruded on a fourcolumn hydraulic press at a extrusion ratio of 13:1. Compression sam­ ples of φ 8 mm � 12 mm were prepared along the extrusion direction (ED) and annealed at 500 � C for 2 h, followed by quenching with cold water. The chemical compositions of the alloy were determined by induc­ tive coupled plasma emission spectrometer (ICP-OES) which were shown in Table 1. The specimens were compressed along the ED at a temperature of 350 � C with a strain rate of 0.1 s 1 and true strains (ϵ) of 0.05, 0.10, 0.15, and 0.20 using a Gleeble-3500 thermomechanical simulator. After the thermal compression was completed, the sample was immediately quenched with water within 5 s to retain the com­ pressed microstructure. The phase of as-extruded sample was examined using a Brooke XD8 ADVANCE A25 X-ray diffractometer (XRD) using copper target Kα line (λ ¼ 0.1541 nm) at a scan rate of 0.02� /s and 2θ range from 20� to 90� . The microstructure of as-extruded sample, second phase, and dislocation multiplication during the deformation were observed and analyzed using a Zeiss ZEISS-6035 field-emission scanning electron microscope (SEM) and JEM-2100F transmission electron mi­ croscope (TEM), respectively. Electron backscatter diffraction (EBSD) analysis was performed on selected areas using Oxford HKL Channel 5 software.

3.2. Flow curve of as-extruded Mg–1Al–6Y alloy Fig. 2 shows the true stress-strain curve and strain-hardening rate curve of the as-extruded Mg–1Al–6Y alloy at a temperature of 350 � C and a strain rate of 0.1 s 1. For comparison, the figure also exhibits the corresponding curve of the same sample compressed at room tempera­ ture. It can be seen that the flow stress increased significantly after yielding at both two temperatures, which is a typical characteristic of work hardening. Although the high temperature strength of Mg–1Al–6Y alloy is lower than room temperature strength, it is interesting that the phenomenon of work hardening characteristics at high temperature is different from that of AZ31 alloy. Moreover, the yield strength of asextruded Mg–1Al–6Y alloy was 98 MPa, and no obvious yielding plat­ form and peak value were observed when compressed at 350 � C and a stain rate of 0.1 s 1. To further analyze the hot deformation behavior of Mg–1Al–6Y alloy, the corresponding strain-hardening rate curve was shown in Fig. 2(b). During the hot compression, the initial high strainhardening rate sharply decreased at the initial stage of deformation (stage I) and then slowly decreased with increasing strain (stage II). It’s worth noting that the strain-hardening rate did not increase during the entire deformation (Fig. 2(b)). It indicates that the twinning was not activated in a large amount, and thence the slip dominated the entire deformation stage [26]. 3.3. Microstructure and microtexture evolution of as-extruded Mg–1Al–6Y alloy by varying deformation To further study the work-hardening phenomenon and deformation mechanism of as-extruded Mg–1Al–6Y alloy during hot compression deformation, the microstructure and texture evolutions of the deformed alloy at 350 � C and 0.1 s 1 strain rate under various strains were studied using ex-situ EBSD. Fig. 3 illustrates the EBSD microstructures and misorientation angle distribution maps of as-extruded Mg–1Al–6Y alloy under different deformations. As shown in Fig. 3(a), when the strain was 0.05, fine recrystallized grains were present around the original coarse grains, and the coarse grains underwent a large amount of deformation. Twinning was observed inside a part of the coarse grains, but few twinning was observed in fine grains. In Fig. 3(b), the misorientation angle was mainly concentrated in the low-angle GBs ((LAGBs <10� ), while extremely low at the high-angle GBs (HAGBs>10� ). This is prob­ ably related to the dislocation pile-up caused by unfavorable LAGBs migration with strain increase [27]. The presence of a misorientation angle peak at 86� confirmed the occurrence of twinning in Fig. 3(a). When the compressive strain reached 0.10, the grain was broken to some extent, the twinning in original coarse grains was consumed, and furthermore lead to the amount of twinning decreased. The HAGBs gradually reduced, meanwhile the LAGBs increased, and the peak value at 86� decreased (Fig. 3(d)). Fig. 3(e) shows that as the strain increased to 0.15, the fragment of original coarse grains aggravated, and the number of twinning further decreased. The peak frequency of LAGBs increased to 0.5, and the peak at 86� gradually disappeared. Fig. 3(g)

3. Results 3.1. Microstructure and texture of as-extruded Mg–1Al–6Y alloy Fig. 1 shows the microstructure, XRD pattern, and EBSD maps of the as-extruded Mg–1Al–6Y alloy. As shown in Fig. 1(a), a small amount of white precipitate was distributed at the grain boundaries (GBs) in ma­ trix. The size of precipitates was in the submicron range, and the area fraction was 0.49%, indicating extremely low distribution density. The inset in Fig. 1(a) reveals the corresponding energy disperse spectroscopy (EDS) spectrum. As can be seen from Fig. 1(b), in addition to the diffraction peak of α-Mg phase, the diffraction peak of Al2Y phase existed in the alloy. Moreover, the TEM image of the precipitates Table 1 Alloy compositions analysis. Nominal composition (wt.%) Mg–1Al–6Y

Analyzed composition (wt.%) Mg (%)

Al (%)

Y (%)

93.22

0.96

5.82

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Fig. 1. Microstructures and texture of as-extruded Mg–1Al–6Y alloy. (a) SEM image, (b) XRD pattern, (b-1) TEM image of a precipitate and selected electron diffraction pattern in the corresponding region, (c) inverse pole figure map, and (d) (0001) basal-plane pole figure.

Fig. 2. (a) Ture stress-strain curves at room temperature and high temperature, and (b) corresponding strain-hardening rate curve of as-extruded Mg–1Al–6Y alloy when compressed at 350 � C and stain rate of 0.1 s 1.

indicates that when the strain reached 0.20, the fine grains distributed around the large grains, and sawtooth GBs appeared, resulting from lattice distortion and subgrain formation caused by dislocation [28]. No twinning was observed at the same time, and the peak at 86� dis­ appeared (Fig. 3(h)). Twinning types and their fractions during compression have signif­ icant influences on the strain hardening and texture evolution. There­ fore, it is necessary to further investigate the twinning behavior. The twin boundary maps and {10–12} extension twinning volume fraction distribution of as-extruded Mg–1Al–6Y alloy with compressive strains along the ED of 0.05, 0.1, 0.15, and 0.2 were illustrated in Fig. 4. At a compression strain of 0.05, the area fraction of {10–12} extension twinning was 11.2% in the interior of some coarse grains. The twinning was substantially parallel to each other, and twinning intersection existed in a few grains. Meanwhile, the {10–11} contraction twinning

and {10–11}-{10–12} double twinning were not observed. When the applied strain increased from 0.05 to 0.2, the volume fraction of twin­ ning gradually decreased from the initial 11.2%–0.95% (Fig. 4(e)). The SF can affect the activity of {10–12} extension twinning during hot compression will be discussed later. Fig. 5 shows the pole figures (PF) of Mg–1Al–6Y alloy at strains of 0.05, 0.1, 0.15, and 0.2 along the ED. Compared with Fig. 1(d), it can be seen that the grain orientation deflected towards the TD to some extent in Fig. 5(a), while the <2-1-11> rare-earth texture still occurs, which results in the texture intensity of 5.6 was lower than the initial state. Fig. 5(b) reveals that when the strain increased to 0.1, the texture in­ tensity decreased to 5.39. As the strain increased to 0.2, the original grains in the alloy fragmented; the preferred orientation of grains further weakened; the texture intensity decreased to 5.06 (Fig. 5(d)).

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Fig. 3. EBSD microstructure and misorientation angle distribution maps of as-extruded Mg–1Al–6Y alloy compressed to (a, b) ε ¼ 0.05, (c, d) ε ¼ 0.10, (e, f) ε ¼ 0.15, and (g, h) ε ¼ 0.20.

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Fig. 4. Twin boundary maps of as-extruded Mg–1Al–6Y alloy compressed to (a) ε ¼ 0.05, (b) ε ¼ 0.10, (c) ε ¼ 0.15, (d) ε ¼ 0.20, and (e) area fraction distribution of {10–12} extension twinning.

Fig. 5. Pole figures of as-extruded Mg–1Al–6Y alloy deformed to compressive strains of (a) 0.05 (b) 0.1 (c) 0.15 and (d) 0.2.

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4. Discussion

metal undergoes plastic deformation, additional stress is generated at the GBs to ensure plastic strain compatibility between adjacent grains [37]. This leads to an increase in the geometrically necessary disloca­ tions (GNDs) density near the GBs. Gao et al. [38] revealed that the density of GNDs ρGND ¼ 2θ ub has a positive correlation with the local misorientation angle (KAM value). Therefore, as the amount of strain increases, the average KAM increases sharply, and the density of GNDs formed locally by microscopic plastic instability increases continuously. Therefore, when the strain was increased to 0.2, a large orientation deviation occurred in both the interior and GBs, it is therefore inferred that the large density of dislocations were formed. It is well known that the activity of the slip systems and {10–12} extension twinning is strongly depended on the grain orientation. In general, the activation of slip systems and twinning follows the Schmid law [39], i.e., the grains with higher SF are more favored for the acti­ vation of slip systems and twinning. Fig. 7 shows the average SF dis­ tribution of basal slip, prismatic slip, pyramidal slip and {10–12} extension twinning of as-extruded Mg–1Al–6Y alloy with strains of 0, 0.05, 0.10, 0.15, and 0.2. As shown in Fig. 7, the average SFs of basal slip, prismatic slip, pyramidal slip and {10–12} extension twinning in the as-extruded sample were 0.35, 0.33, 0.38, and 0.30, respectively, indicating that the orientation of most grains in the initial as-extruded alloy facilitated the activation of slip systems and twinning. When strain was 0.05, the SF value of the slip is slightly changed, and the basal slip is dominant, the flow stress increases linearly at same time. When dislocation slip is hindered during hot deformation, twins will be generated in the region of stress concentration and control the slippage in crystals. Thence, the fraction of {10–12} extension twinning reached 11.2% when strain was 0.05 (Fig. 4(a)). Fig. 7(b and c) further gives the TEM microstructure of twinning and dislocations when strain was 0.05. It is worth mentioning that many dislocations can be observed in the structure of twins, which are mainly on slip planes and contain sharp twin boundaries. Moreover, some dislocations are entangled at the boundary of the twinning zone. It can be inferred that slip and twinning have synergistic effects. Roberts et al. [40] proposed that the specific strain provided by twinning can be calculated using the following equation:

4.1. Effect of twins and slips on deformation Second phases, texture, solutes content and the resultant deforma­ tion mode have been generally suggested to induce the variation in curves characters during the compression of Mg alloys. Kwak [29] have founded that the volume fraction of second phase was related to the hot workability. The volume fraction of Al2Y phase was very small, only 0.49% in this work (Fig. 1). This is mainly because of the large solid solubility of Y in magnesium, and most of the Y element dissolved in the magnesium matrix. Therefore, the strengthening effect of undissolved Al2Y particles by dispersion strengthening was limited. It was worth noting that the basal texture of the as-extruded Mg–1Al–6Y alloy was weak. According to the previous report [30], RE addition to magnesium alloys have great effects on the strength and the type of texture. Moreover, due to the recovery and recrystallization of grains during annealing heat treatment after extrusion, which weakened the basal texture [31]. Furthermore, the formation of rare-earth texture also weakened the basal texture to some extent [32]. Therefore, It resulted in tensile stress along the c-axis of some coarse grains and thus facilitated the activation of {10–12} extension twinning [33] during compression. Moreover, the larger the grain size, the smaller the con­ straining force produced by adjacent GBs, and thence the twinning deformation mechanism is more favorable. Dobron et al. [34,35] found that the nucleation and growth of {10–12} extension twinning prefer­ entially occurred in coarse grains during the compression of as-extruded alloys. Fig. 6 displays the kernel average misorientation (KAM) images of Mg–1Al–6Y alloy with compression strains of 0.05, 0.10, 0.15, and 0.2 and the variations in average KAM with increasing strain, which quali­ tatively reflects the homogeneous degree of plastic deformation. The high values indicate larger plastic deformation or high defect density [36]. The average KAM of annealed alloy after extrusion was very low, i. e., 0.264, due to the dislocations formed during hot extrusion annihi­ lated through dynamic recovery (DRV) and DRX after the annealing treatment. As shown in Fig. 6(a), when the strain was 0.05, the dislo­ cations multiplied, and more and more dislocations were piled up, resulting in a significant increase of the average KAM from 0.264 to 0.573. As the strain increased to 0.15, the strain energy inside the ma­ terial accumulated thence the average KAM linearly increased to 1.25. Furthermore, the KAM figure exhibits that at relatively early stage of deformation, local dislocations were highly concentrated near the GBs, i. e., at strains of 0.05 and 0.1 (Figs. 6(a) and (b)). When a polycrystalline

εtwin ¼ f twin ⋅m⋅γtwin where εtwin , f twin , m, and γtwin are the twinning strain, area fraction of twinning, average SF (SF) of {10–12} extension twinning, and charac­ teristic twinning shear, respectively. Therefore, when the strain was 0.05, f twin was 11.2%, m was 0.315, and γtwin was 0.129. εtwin was calculated to be 0.455%, indicating that the softening effect of {10–12}

Fig. 6. Kernel average misorientation maps of as-extruded Mg–1Al–6Y alloy with compressed to (a) ε ¼ 0.05, (b) ε ¼ 0.10, (c) ε ¼ 0.15, (d) ε ¼ 0.20, and (e) average KAM of as-extruded Mg–1Al–6Y alloy. 6

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Fig. 7. (a) Variations in average SFs for basal slip, prismatic slip, pyramidal slip and {10–12} extension twinning for as-extruded Mg–1Al–6Y alloy with strains of 0, 0.05, 0.1, 0.15, and 0.2; (b),(c) TEM micrographs with a strain of 0.05.

extension twinning during deformation was small. Nevertheless, the activation of {10–12} extension twinning lead to crystal orientation rotation and alleviation of local stress concentration generated during the deformation, resulting in promoting the activation of various slip systems and dislocation slippage in the crystal [41]. Therefore, the SF of pyramidal slip was up to 0.39 with the strain increasing to 0.05, indicating that the pyramidal slip might be already activated at the initial stage of deformation during a high-temperature compression. In order to investigate the dislocation micrographs, Fig. 7(c) gives the TEM result. Based on the SAED, it is determined to be pyramidal dislocations. Related study have shown that dislocations with type of Burgers vector can accommodate c-axis strain [42]. Therefore the activation of nonbasal slip systems ensures strain compatibility during plastic deformation. This is beneficial for dislocation slip and

then for the improvement of alloy properties. With the increase of strains, the activity of basal slip and pyramidal slip show an increasing trend while the activity of prismatic slip de­ creases. As the strain increased from 0.05 to 0.2, the average SF of basal slip and pyramidal slip increased from 0.35 to 0.37, 0.387 to 0.40, while that of prismatic slip and {10–12} extension twinning gradually decreased from 0.316 to 0.298, 0.318 to 0.306, respectively. It can be found that the average SF factors remained higher than 0.3 during the compression. Jiang et al. [43] revealed that a lower area fraction of twins have little effect on the texture randomization during the hot deformation of AZ31 Mg alloy. Therefore, the minor change can be mainly attributed to the lattice distortion caused by activation of non­ basal slips, which decreased the amount of grains with the initial preferred orientation, consistent with the result of {0001} PF in Fig. 5.

Fig. 8. TEM bright-field images of as-extruded Mg–1Al–6Y alloy with compressed to (a) ε ¼ 0.05, (b) ε ¼ 0.10, (c) ε ¼ 0.15, and (d) ε ¼ 0.20. 7

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To further explore the hot deformation behavior of alloy, TEM was carried out on the investigated samples. Fig. 8 shows the TEM brightfield images of samples with strains of 0.05, 0.10, 0.15 and 0.2. The �! dislocation morphology can be characterized by ! g ⋅ b ​ ¼ 0 invisible criterion [44], and the only pyramidal dislocation can be observed in the diffraction mode of ! g ¼ 0002. Therefore, in the deformation samples, the pyramidal dislocation was clearly observed. As shown in Fig. 8(a), when the strain was 0.05, several long and straight-shaped dislocations (as indicated by the arrows) were formed inside the grains, indicating that the pyramidal dislo­ cation was activated at a small strain, and its average SF was 0.39, as shown in Fig. 7(e). However, the amount of slip systems that could be activated was limited. As the strain increased, more pyramidal dislocations occurred, and the dislocation density increased (Fig. 8 (a–d)). The dislocation entanglement formed by multiple dislocation slips suppressed the growth and movement of twinning, thus reduction the number of twinning (Fig. 4). This indicates that in the early stage of deformation, the nonbasal slip was activated. The slip was the dominant deformation mechanism during the entire deformation, dislocation multiplication and entanglement contribute to the work-hardening ef­ fect, leading to a material exhibiting work-hardening behavior all the time. The work-hardening effect yields at a higher strength while ensuring plasticity and preventing premature failure of material.

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5. Conclusions The microstructure evolution and deformation behaviors of asextruded Mg–1Al–6Y alloy at a temperature of 350 � C and a strain rate of 0.1 s 1 under uniaxial compression were investigated using exsitu EBSD analyses. (1) During the compression, Mg–1Al–6Y alloy exhibited workhardening behavior with no obvious yielding platform, and the yield strength was 98 MPa. (2) In the early stage of deformation, the presence of initial <2-111> rare-earth texture resulted in a small amount of {10–12} extension twinning only formed in a part of coarse grains. As the strain increased, the dislocation density increases and the twins were consumed, which resulted in a decrease in the volume fraction of twins. (3) The pyramidal slip was activated at the beginning of deformation, and the slip was the dominant deformation mech­ anism. In addition, the lattice rotation caused by the activation of non-basal slip leads to a gradual decrease in the texture strength of {0001} PF. (4) During the high-temperature compression of as-extruded Mg–1Al–6Y alloy, a gradual decrease in the number of twinning and activation of pyramidal slip result in work-hardening behavior of alloy. Author contributions section Mengna Zhang: conceived and designed the study. Jinhui Wang: conceived and designed the study. Yunpeng Zhu: performed the exper­ iments. Lei Zhang: performed the experiments. Peipeng Jin: reviewed and edited the manuscript. All authors read and approved the manuscript Prime novelty statement Firstly, as-extruded Mg–1Al–6Y alloys were successfully fabricated via casting process. Secondly, the microstructural and textural evolution of as-extruded Mg–1Al–6Y alloy as a function of strain at a temperature of 350 � C and strain rate of 0.1 s 1 were studied using ex-situ electron backscatter diffraction (EBSD). Thirdly, the combined effect of 8

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