Plasticity Induced by Twin Lamellar Structure in Magnesium Alloy

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ScienceDirect J. Mater. Sci. Technol., 2013, 29(12), 1123e1128

Plasticity Induced by Twin Lamellar Structure in Magnesium Alloy Xiyan Zhang*, Chao Lou, Jian Tu, Qing Liu School of Materials Science and Engineering, Chongqing University, Chongqing 400030, China [Manuscript received December 12, 2012, in revised form February 2, 2013, Available online 22 October 2013]

Effect of {10e12} twins on the mechanical properties of magnesium alloy has received considerable research interest. A hot-rolled AZ31 Mg alloy sheet was subjected to dynamic plastic deformation with the aim of introducing {10e12} twin lamellar structure. It has been found that higher strength and better ductility are obtained when tensile loading is perpendicular to the c axis of twin region of the twin lamellar structured sample, indicating that the plasticity improvement caused by twins depends on the special strain path. The fracture morphology of the twin lamellar structured sample shows a dimple fracture mode under tensile loading perpendicular to the c axis, while the cleavage fracture with river pattern has been observed in other fractured samples. Above experimental results indicate that the interaction of dislocations and twin lamellae may play an important role in improving mechanical properties of Mg alloy. KEY WORDS: Magnesium alloy; Twin lamellar structure; Mechanical behavior; Plasticity

1. Introduction Due to the characteristic of the hexagonal close packed (HCP) structure, magnesium alloys have the limited number of slip systems, which results in the poor formability at room temperature. Consequently, deformation twinning plays an important role in plastic deformation of polycrystalline magnesium alloys[1e3]. Although several types of twins have been found in magnesium alloys, {10e12} twinning is thought to be activated more frequently[4e6]. In general, the critical resolved shear stress (CRSS) for {10e12} twinning is greater than that for basal slip, but less than that for non-basal slip[7]. This results in easier activation of {10e12} twinning during uniaxial deformation at room temperature for Mg alloys with conventional grain sizes. At the later stages of deformation, other twin such as secondary twins can also be activated, but it is usually accompanied by microcrack attributable to the interactions of twinningeslip, twinningetwinning and twinegrain boundary[1,3,8]. Barnett[4] found that {10e12} twinning can increase the uniform elongation. Therefore, compared with other twinning mechanisms, {10e12} twinning plays a more important role in improving the workability of Mg alloys[5,9,10]. The common view is that twinning cannot improve strength and plasticity simultaneously. But an interesting phenomenon

that both strength and plasticity increased significantly with decreasing lamellar thickness has been found in nanocrystalline Cu with twin lamellar structure. For example, when twin lamellar thickness is 15 nm, the tensile strength of 1.0 GPa and an elongation of up to 13% can be obtained[11]. The reason for improving plasticity is that these twin boundaries are not only as obstacles to dislocation gliding, but also as gliding planes to store dislocation. The new full/partial dislocations caused by the interaction of dislocation in the lamella and TB (twin boundary) glide accommodate deformation[11e13]. Another result has been reported[14,15] that twins can facilitate the plasticity and refine grains when a face-centered cubic (fcc) metals with a metastable stacking fault energy is subjected to severe plastic deformation. During plastic deformation of Mg alloys, parent grains also exhibit segmentation by {10e12} twins, and this grain refinement seems to peak at certain plastic strain. In this paper we report that twin lamellar structure can improve both plasticity and strength of AZ31 alloy. These results thus reveal a route to improve workability of hcp magnesium alloy. High strain rate can enhance the activation of {10e12} twinning dramatically[16,17]. So to facilitate twinning, dynamic plastic deformation (DPD) was employed to achieve pre-deformation of samples. 2. Experimental

Corresponding author. Prof., Ph.D.; Tel./Fax: þ86 23 65112154; E-mail address: [email protected] (X. Zhang). 1005-0302/$ e see front matter Copyright Ó 2013, The editorial office of Journal of Materials Science & Technology. Published by Elsevier Limited. All rights reserved. http://dx.doi.org/10.1016/j.jmst.2013.10.017 *

2.1. Initial microstructure of material The hot-rolled AZ31 Mg alloy sheet (Mge3%Ale1%Zn) has twin-free equiaxed grain structure (Fig. 1(a)) and an average grain size of 34 mm evaluated by linear intercept method. The

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Fig. 1 (a) Microstructural characteristics of the hot-rolled material; (b) (0001) pole figure of the as-rolled material. RD and TD represent the rolling direction and the transverse direction respectively.

alloy sheet was compressed at 593 K to get a strong basal texture (the c axis of lattice parallel to the normal direction (ND), as shown in Fig. 1(b)), then annealed at 573 K for 2.5 h to eliminate deformation twins. 2.2. DPD treatment Parallelepiped (30 mm  30 mm  22 mm in rollinge transverseenormal directions) was cut from the hot-rolled AZ31 sheet. Samples were subjected to DPD just once with Instron Dynatup 8120 testing machine at room temperature (w293 K) and the strain rate is about 103 s1. The loading direction (LD) is

perpendicular to the c axis of crystal lattice in order to introduce {10e12} twinning lamellar structure. The sample was placed on a lower anvil and was compressed by an upper impact anvil. Details of the process were described in literature[18]. The deformation strain is defined as ε ¼ ln(L0/Ld), where L0 and Ld are the initial and final height of the samples, respectively. The samples were deformed to ε ¼ 5%. The initial and pre-deformed textures were measured by X-ray diffraction in a Rigaku D/max2500 PC. Scanning electron microscopyeelectron back-scattered diffraction (SEMeEBSD) was utilized to analyze the microstructure and texture of samples before and after deformation. EBSD scans were performed by FEI Nova400 FEG-SEM on an

Fig. 2 (a) (0001) pole figure and (b) the {10e12} twin lamellar structure of DPD sample with 5% strain; (c) the statistical distribution of twin lamellae thickness; (d) misorientation distribution of boundaries. LD represents the loading direction.

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area of 300  300 with a step size of 1 mm and EBSD information was acquired with Channel 5 software. A basal texture (the c axis almost parallel to LD) and twin lamellar structure were obtained due to the activation of {10e12} twins, as shown in Fig. 2(a) and (b), respectively. Fig. 2(c) shows the statistical distribution of the lamellar thickness. The average thickness of lamellae is about 3.2 mm. In magnesium alloys, the coalescence of different {10e12} twin variants can form the 60 <10e10> and 60.4 <8e1e70> boundaries[19,20], causing the TB fraction decreased. But this phenomenon usually happens in large strain deformation. As shown in Fig. 2(d), the 86 <1e210> boundary is dominant in the twin structure of the sample with ε ¼ 5%.

referred to as HR-90 (tension along RD) and DPD-90 (tension along the initial TD), respectively. The similar naming is used for tensile loading parallel to the c axes of two samples, referred to as HR-0 (tension along ND) and DPD-0 (tension along LD), respectively. Fig. 3(deg) shows the changes of mechanical property after DPD. The tensile curves of samples with loading direction perpendicular to the c axis are shown in Fig. 3(d), while the tensile curves of samples with loading direction parallel to the c axis are shown in Fig. 3(f). Fig. 3(e) and (g) shows good repeatabilities of tensile results.

2.3. Tensile tests

A mechanical anisotropy arises in the samples containing twin lamellar structure. As shown in Fig. 3(d), the pre-deformed sample exhibits higher strength and better plasticity (DPD-90 curve), compared to the hot-rolled sample (HR-90 curve) when the tensile direction is perpendicular to the c axis. However, both strength and plasticity of the pre-deformed sample decrease compared to that of the hot-rolled sample when the tensile direction is parallel to the c axis (Fig. 3(f)). The tensile fracture surface has been examined to analyze the fracture characteristics, as shown in Fig. 4. When the tensile direction is perpendicular to the c axis, the fracture morphology of the twin lamellar

At room temperature, quasi-static tensile tests were performed in both the hot-rolled and pre-deformed AZ31 samples up to failure, using SHIMADZU AG-X10KN machine at an initial strain rate of 102 s1. Tensile tests were performed in dogboneshaped samples with a gage length of 5 mm and a width of 1.26 mm. The sampling method of tensile samples is shown in Fig. 3(aec). The tensile directions are designed to be perpendicular to the c axis of crystal lattice of the initial hot-rolled sample and the c axis of twin of the pre-deformed sample,

3. Results and Discussion

Fig. 3 Sampling method of tensile samples: (a) the hot-rolled sample; (b) the twin lamellar structured sample; (c) schematic diagram of tensile sample; typical tensile flow curves of samples with ε ¼ 5%: the loading directions are perpendicular (d) and parallel (f) to c axis, respectively. A repeatability of tensile result of the twin lamellar structured samples with ε ¼ 5% is shown in (e) and (g), respectively. LD represents the loading direction.

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structured sample shows a mainly dimple fracture mode (Fig. 4(a)) and the hot-rolled sample presents a cleavage fracture with river pattern (Fig. 4(b)), indicating the operation of different deformation mechanisms. When the tensile direction is parallel to the c axis, the cleavage fracture with river pattern is also observed in both the twin lamellar structured and hot-rolled samples, as shown in Fig. 4(c) and (d), respectively. To explore the reasons for improving plasticity in the samples with twin lamellar structure in Figs. 3(d) and 4(a), the workhardening rate and texture change during deformation have been studied. Fig. 5(a) and (b) shows the work-hardening rate curves of the samples in which the tensile direction parallel and perpendicular to the c axis, respectively. Fig. 5(a) shows a typical twinning dependence of deformation behavior[20e22]. However, for the twin lamellar structured sample twin strengthening reaches the maximum at a smaller strain of about 5.6%, indicating that the earlier texture change caused by pre-deformation promotes the activation of slip system. An interesting result is shown in Fig. 5(b). It can be seen that for the twin lamellar structured sample the work-hardening rate (red line, DPD-90 ) decreases more quickly than that of the hot-rolled sample (black line, HR-90 ), though the shapes of these two curves are similar. It is known when the tensile direction is parallel to the basal

plane, the basal slip system is difficult to be activated because Schmid factor is nearly zero. And prismatic slip plays an important role in the plastic deformation of Mg alloys[23], while for pure zirconium gliding on pyramidal plane accommodates deformation[24]. However, there always exist some grains which have a deviation from ideal orientation of basal texture in wrought Mg alloys, and basal slip can accommodate deformation in these grains because the CRSS of basal slip of Mg alloys is low (about 2 MPa/mm2[25]). With increasing plastic strain, more dislocations would be activated[3,23]. Therefore, dislocation slip dominates the plastic deformation when tensile loading is perpendicular to the c axis, as shown as black line in Fig. 5(b). However, for the twin lamellar structured samples there are some factors to complicate dislocation motion. Firstly, sufficient residual dislocations caused by pre-deformation make slip easier. Secondly, the TBs (or twin dislocation) act as barriers to dislocation gliding[26e28]. Although work hardening is caused by TBs as barriers to dislocations, twins have the favorable effects on the uniform elongation of magnesium[4] and titanium[29]. In fact, due to the reactions between gliding dislocations and TB dislocation, TB can promote dislocation motion in some case[30], which can be clarified in crystallography[31]. Even thus it is unclear whether

Fig. 4 Tensile fractographs of samples: (a) DPD-90 ; (b) HR-90 ; (c) DPD-0 ; (d) HR-0 .

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Fig. 5 Work-hardening rate curves of the samples in which the tensile direction is parallel (a) and perpendicular (b) to the c axis, respectively.

this case is similar to that in nano-twin lamellar structured Cu[11,14,15]. Furthermore, the parallel twin bands can divide a grain into a lamellar structure. This means that twin lamellar structure can lead to grain refinement[30,32,33], which would improve both strength and ductility of materials. Another research emphasis focuses on the texture changes of the twin lamellar structured and hot-rolled samples during tension. Fig. 6 presents the pole figures of samples before and after deformation. As shown in Fig. 6(a), similar textures were obtained before and after tension of the hot-rolled sample when tensile loading is perpendicular to the c axis, indicating that grain orientation has little change during deformation. But the same deformation condition leads to the slight weakening in the twinning texture of the twin lamellar structured sample (Fig. 6(b) and (c)), which may be due to that the dislocation slip results in

some grains rotated away from the LD. When the tensile direction is parallel to the c axis of crystal lattice of the hot-rolled sample, twinning leads to an w86 rotation of c axis of grains (Figs. 1(b) and 6(d)). An obvious texture change happens in the twin lamellar structured sample when tensile loading is parallel to the c axis. As shown in Fig. 6(e) and (f), the twinning texture caused by pre-deformation disappears and the initial texture is recovered. This is because tensile reloading causes an untwinning process[10,34,35]. The development of texture plays an important role in determining the workability of Mg alloys during plastic deformation[36,37]. For hcp single crystals the tensile loading parallel to the c axis can lead to the activation of {10e12} twinning, while the tensile loading perpendicular to the c axis leads to the activation of {10e11} twinning. However, for magnesium alloy the tensile deformation perpendicular to the c

Fig. 6 Texture distributions in (0001) pole figures: (a) the hot-rolled sample after HR-90 tension; the twin lamellar structured samples before (b) and after (c) DPD-90 tension; (d) the hot-rolled sample after HR-0 tension; the twin lamellar structured samples before (e) and after (f) DPD0 tension. LD represents the loading direction. T represents tensile loading.

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axis is mainly controlled by dislocation slip, though some compression twins such as {10e11} twins have also been found[20e23]. Above experimental results indicate that the interaction of dislocations with twin lamellae may improve both strength and ductility of Mg alloy when tensile loading is perpendicular to the c axis. 4. Conclusion A hot-rolled AZ31 Mg alloy sheet was subjected to DPD parallel to the rolling direction with the aim of introducing {10e 12} twins. The tensile experimental results show that both strength and ductility are improved when the tensile direction is perpendicular to the c axis of twin region of the twin lamellar structured sample, indicating that the plasticity improvement caused by twins depends on the special crystal direction. The fracture morphology of the twin lamellar structured sample shows a mainly dimple fracture mode under the tensile loading perpendicular to the c axis, while the cleavage fracture with river pattern has been observed in other fractured samples, indicating the operation of different deformation mechanisms. Above experimental results indicate that the improvements of strength and ductility of Mg alloy may be attributed to the interaction of dislocations with twin lamellae. Acknowledgments This work was supported by the National Natural Science Foundation of China (Grant Nos. 51071183, 50890170 and 51271208), Basic Research of China (No. 2010CB631004) and Fundamental Research Funds for the Central Universities (No. CDJXS11132225). REFERENCES [1] [2] [3] [4] [5]

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