Microstructural changes during tension of friction-stir welded AZ31 magnesium alloy

Accepted Manuscript Microstructural changes during tension of friction-stir welded AZ31 magnesium alloy

S. Mironov, T. Onuma, Y.S. Sato, S. Yoneyama, H. Kokawa PII: DOI: Reference:

S1044-5803(17)30212-7 doi: 10.1016/j.matchar.2017.05.016 MTL 8677

To appear in:

Materials Characterization

Received date: Revised date: Accepted date:

23 January 2017 8 May 2017 10 May 2017

Please cite this article as: S. Mironov, T. Onuma, Y.S. Sato, S. Yoneyama, H. Kokawa , Microstructural changes during tension of friction-stir welded AZ31 magnesium alloy, Materials Characterization (2017), doi: 10.1016/j.matchar.2017.05.016

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ACCEPTED MANUSCRIPT Microstructural changes during tension of friction-stir welded AZ31 magnesium alloy S. Mironova1, T. Onumaa, Y.S. Satoa, S. Yoneyamab, H. Kokawaa a

Department of Materials Processing, Graduate School of Engineering, Tohoku University,

6-6-02 Aramaki-aza-Aoba, Sendai 980-8579, Japan b

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Department of Mechanical Engineering, Aoyama Gakuin University, 5-10-1 Fuchinobe,

Sagamihara 252-5258, Japan

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In this work, microstructural changes occurring during transverse tension tests of friction-stir welded AZ31 magnesium alloy were systematically studied in order to clarify the

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microstructure-properties relationship. To this end, digital image correlation and electron backscatter diffraction (EBSD) were employed. As expected, plastic deformation in the stir

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zone was found to be preferentially concentrated at its advancing- and retreating sides. In both these locations, the microstructure evolution was shown to be a complex process consisting of two stages. The first stage was characterized by a profuse {10 1 2} twinning

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giving rise to a {0001}  1120  texture component. In the second stage, enhanced slip activity led to extensive formation of low-angle boundaries and to the development of a fiber

{hkil}  1120  texture.

{0001}  1120 

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The twin-induced

orientation was shown to promote double

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{10 1 1}  {10 1 2} twinning which is known to lead to final failure. It was therefore suggested that the cause of the well-known premature fracture of magnesium welds at the stir zone extremity may be due to the profuse {10 1 2} twinning.

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Keywords: Magnesium; Friction-stir welding; Microstructure-properties relationship; Electron-backscatter diffraction (EBSD); Digital image correlation

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Corresponding

author;

[email protected]

Phone:

+81-22-795-7353,

FAX:

+81-22-795-7352,

E-mail:

ACCEPTED MANUSCRIPT Microstructural changes during tension of friction-stir welded AZ31 magnesium alloy

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1. Introduction Magnesium alloys exhibit a promising combination of properties, including low density, high specific strength, and good dumping capacity. However, poor weldability of these materials by conventional fusion techniques restricts their widespread industrial application. In this regard, friction-stir welding (FSW), an advanced solid-state process, is particularly attractive for the joining of magnesium alloys. Initial studies in this field demonstrated the good feasibility of FSW for production

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of defect-free welds in these materials [1], which gave rise to interest in the underlying microstructural and textural changes. Extensive research has demonstrated that FSW produces a very specific microstructure in the stir zone which is characterized by relatively fine grain size [2-11], an exceptionally strong basal texture (up to ~50 mrd) [12-15], and an extremely inhomogeneous texture distribution [2, 10, 15-27]. The sharp basal texture has been shown to dictate misorientation distribution, characterized by a relatively high (~35%) fraction of low-angle boundaries as well as crystallographic preference of 30o<0001> misorientations [12-15]. The joint efficiency of magnesium friction-stir welds is well accepted to be

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governed by inhomogeneous texture distribution, which favors slip near the stir zone edge but inhibits it in the stir zone center [2, 10, 15-25]. This is commonly believed to lead to strain localization and premature fracture during transverse tensile tests and thus deteriorates mechanical properties of the welds [2, 10, 17-20, 22-23, 24-28]. Considering the well-known high sensitivity of the mechanical behavior of magnesium to texture, this conception seems to be entirely obvious. As a result, to the best of the authors’ knowledge, the actual microstructure evolution of magnesium joints during tension tests has never been studied systematically. It is worth noting, however, that magnesium is prone to mechanical twinning, which may significantly change the crystallographic texture and thus fundamentally affect the slip activity. Two twin modes are believed to be of particular importance in magnesium [29-30]: {10 1 2} tensile twinning and {10 1 1} contraction twinning, which provides extension and compression along the <0001> axis, respectively. Indeed, extensive tensile twinning has been reported to operate during tension of the magnesium friction-stir joints [10, 23-25]. Obviously, activation of the twinning mechanism should complicate the above-mentioned texture-properties relationship. In an attempt to provide further insight into this issue, microstructure and texture evolution during transverse tension tests of friction-stir welded magnesium was studied in this work. To this end, the digital image correlation technique was used for

ACCEPTED MANUSCRIPT accurate strain measurements and electron backscatter diffraction (EBSD) was employed for microstructure and texture observations.

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2. Material and Experimental Procedures The material used in the present investigation was a commercial AZ31 magnesium alloy with a nominal chemical composition of Mg-3.0Al-1.0Zn (wt. %). This is a typical and comparatively simple magnesium alloy whose deformation behavior is well documented. The material was produced by extrusion at 350 oC2 followed by annealing at 300oC for 2 hours and supplied as 4-mm thick sheets. It was

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characterized by a relatively coarse-grained (~200 m) microstructure and a moderately strong {hkil}  1210  fiber texture. In more details, the initial

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microstructure is described elsewhere [12]. The supplied material was friction-stir welded in a bead-on-plate configuration at a tool rotational speed of 2000 rpm and a tool travel speed of 200 mm/min. The welding tool was fabricated from a tool steel and consisted of a concave shoulder 15 mm in diameter and an M5 threaded cylindrical pin 3.7 mm in length. The principal directions of welding geometry are denoted throughout as welding direction (WD), transverse direction (TD), and normal direction (ND). Other details of the FSW process are described elsewhere [12].

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To examine the mechanical behavior of the welds, tensile specimens were machined transverse to the welding direction. The specimens were centered at the weld line, had a gauge section 35 mm in length and 7 mm in width, and included all of the characteristic microstructural zones developed during the FSW process. To achieve a uniform thickness, the upper surfaces of the tensile specimens were mechanically polished. Accordingly, their thickness was reduced from 4 mm to 3.7 mm. Tension tests were conducted at ambient temperature and a constant crosshead velocity corresponding to a nominal strain rate of 10-3 s-1 using an Instron 5969 testing apparatus. To examine the microstructural evolution, a series of tests to a total tension elongation of 4%, 7%, 10%, and 13%, as well as to final failure (~16%), was performed. To evaluate the strain distributions, the digital image correlation technique was employed [31, 32]. To this end, a random ink pattern was applied to the sample surfaces and a high-speed XiQ MQ042MG-CM digital camera equipped with a Nikon AI AF Micro-Nikkor 200 mm f/4D IF-ED lens was used for image recording. The images were taken at a rate of 1 image per 5.25 seconds (or ~0.5% of nominal strain). A subset size of 15 × 15 pixels (or 450 × 450 m2) was used for the strain calculations. 2

The extrusion ratio had not been provided by material supplier (Osaka Fuji Kogyo, Japan)

ACCEPTED MANUSCRIPT Microstructural observations were performed on the front through-thickness side of the weld (TD×ND plane) and were primarily conducted by EBSD. A suitable surface finish for EBSD was obtained by electro-polishing in a commercial AC2 Struers polishing solution at ~10oC (ice bath) with an applied potential of 40 V. The EBSD analysis was conducted using a Hitachi S-4300SE field-emission gun scanning electron microscope equipped with a TSL OIMTM EBSD system and operated at an accelerated voltage of 25 kV. Depending on the particular purpose, different scan step sizes were used during the mapping. Specifically, low-resolution (overview) EBSD

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maps were acquired using a scan step size of 5 m, whereas high-resolution maps were obtained using a step size of 1 or 0.5 m. On each diffraction pattern, nine Kikuchi-bands were used for indexing, thus minimizing misindexing error. To ensure reliability of the EBSD data, all small grains comprising three or fewer pixels were automatically removed from the EBSD maps using the grain-dilation option of the TSL software. Furthermore, to eliminate spurious boundaries caused by orientation noise, a lower-limit boundary misorientation cut-off of 2o was used. A 15o criterion was applied to differentiate low-angle boundaries (LABs) and high-angle boundaries (HABs). 3. Results

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3.1. As-welded material A composite optical image showing a typical transverse cross section of the obtained welds is given in Fig. 1a. In the figure, AS and RS are advancing side and retreating side, respectively. The base material, stir zone and thermomechanical affected zone can be seen in the figure. On the other hand, no distinct heat-affected zone was found, this effect being thought to be associated with the well-annealed condition of the base material. To examine texture distribution, a sample-scale EBSD map was obtained, as shown in Fig. 1b. In this map, individual grains are colored according to their crystallographic orientations relative to the WD3; the color code triangle is shown in the bottom right corner. As expected, the (0001) basal planes in the stir zone were aligned with the profile of the welding tool and thus the texture distribution was inhomogeneous. This result agrees well with scientific literature [2, 15- 25]. Based on the EBSD data, the distribution of Schmid factors for the (0001)  1120  basal slip (i.e., the dominant slip mode in magnesium alloys) assuming transverse tension was calculated, as shown in Fig. 1c. As expected, the basal slip was favored at the stir zone edge but inhibited along its centerline. This observation is consistent with recent simulations reported by Xin et al. [10, 19]. 3

Here and hereafter, a reader is referred to on-line version of this paper to see figures in color.

ACCEPTED MANUSCRIPT Moreover, the Schmid factor was also found to be relatively high at the AS of the upper section of the stir zone (Fig. 1c).

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3.2. Strain distributions The distributions of extensional strains measured on the front through-thickness side of the weld as a function of global tension elongation are shown in Fig. 2. It is clear that the strain distribution in the stir zone was inhomogeneous in almost the entire range of tensile deformations. Specifically, the strain was preferentially concentrated at the AS and RS of the stir zone; an increased strain was also observed at the AS of the upper portion of the stir zone (Figs. 2b-d). These

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observations are in good agreement with variation of the Schmid factor for the basal slip (Fig. 1c). This correlation perhaps indicates that plastic strain during the tension tests primarily resulted from the basal slip, as is normally anticipated for magnesium alloys.

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3.3. Microstructure evolution As shown in the previous section, the strain in the stir zone was primarily clustered at the AS and RS. Therefore, to track microstructural evolution in the stir zone, microstructural changes in these two locations were studied. In both cases,

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however, they were found to be broadly similar to each other. For the sake of simplicity, therefore, only the microstructure evolution at the AS is considered below. The microstructural changes observed at the RS are provided in supplementary Figs. S1-S4, and Table S1.

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3.3.1. Microstructure morphology Selected portions of EBSD grain-boundary maps illustrating microstructure evolution are summarized in Fig. 3. In the maps, LABs, HABs, and {10 1 2} twin

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boundaries are depicted as red, black, and blue lines, respectively. In the as-welded condition, the microstructure was dominated by nearly equiaxed grains containing a significant proportion of LABs (Fig. 3a). As shown in previous works [12-14], the latter effect was attributable to a relatively strong crystallographic texture produced during FSW. Remarkably, the microstructure contained almost no twins. At relatively low tensile strains (Fig. 3b), profuse twinning was found. Misorientations across the twin boundaries were typically close to 86o  2 1 1 0  , thus indicating the activation of the {10 1 2}  10 1 1  tensile twinning mode. This observation agrees well with previously reported results [10, 23-25]. Importantly, only one or two twin variants (out of six) were present in the grain interior, thus evidencing

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the occurrence of variant selection during the twinning. The twins were usually lenticular- or wedge-shaped and extended across entire grains. Of particular interest was the observation that the twins in neighboring grains were often connected end to end with each other, thus forming extended twin chains. This effect is frequently observed in magnesium alloys being typically attributable to the relatively easy twin-twin transfer across grain boundaries [e.g. 24, 33-38]. In friction-stir welded material, this phenomenon has been reported by Liu et al [24]. In the present study, two sets of the twin chains oriented by ~±55o to the tension axis can be identified (Fig. 3b). Another interesting point was the formation of irregular and poorly-developed LABs4 in some grains (examples are circled in Fig. 3b).

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At a moderate level of tensile strain (Fig. 3c), twinning became very pronounced. Each grain typically contained two twin variants which belonged to the above two sets of the twin chains. After relatively large tensile straining (Fig. 3d), misorientation of the twin boundaries was found to frequently deviate from the exact twin-matrix relationship with some of the boundary segments transforming into random (non-twin) HABs (an example is arrowed). The progressive development of this process leads to gradual transformation of the twins into regular grains; examples of completely transformed twins are circled in Fig. 3d. As described in the scientific literature [39-41], this

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phenomenon is attributable to strain-induced crystallographic rotation of twins and matrix from their initial orientations. In addition to the twin transformations, deformation-induced LABs became very pronounced in grain interiors at the large strains (Fig. 3d). These LABs were relatively short, curved, and tended to concentrate near original grain (or twin-) boundaries.

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3.3.2. Misorientation distribution To provide additional insight into grain structure development, misorientation data were extracted from the EBSD maps and arranged as misorientation-angle and misorientation-axis distributions (Figs. 4 & 5 and Table 1). The misorientation angle data in Fig. 4 were characterized in terms of the total grain boundary length for a given misorientation angle (or small range of the misorientation angles) divided by the area of the EBSD map. This metrics provide a direct comparison of grain-boundary characteristics for different tensile strains, thus enabling more reliable determination of the key physical mechanisms governing the microstructure evolution. In agreement with previous works [12-15], the misorientation distribution of the as-welded material was characterized by the crystallographic preference of 10-30o<0001> boundaries (Figs. 4 & 5a). This effect may be explained in the terms of 4

i.e. short and curved LABs having relatively low misorientation

ACCEPTED MANUSCRIPT the formation of a very strong basal texture during FSW [12-15]. The extensive {10 1 2} twinning observed at low-to-moderate tensile strains (≈ 0.06-0.13) gave rise to a strong peak near 86o in the misorientation-angle distribution (Fig. 4) and clustering of misorientation axes near  2 1 1 0  (Fig. 5b & c). The total fraction of the twin boundaries achieved ~50% of the total grain boundary area (Table 1). The interaction between the (10 1 2)  (01 1 2) twin variants (an example is indicated by an arrow in Fig. 3b) promoted additional enlargement of the area fraction

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of 60o  10 1 0  misorientations (Table 1, Figs. 4 & 5b-c). A subtle increase of the boundaries with the lowest detectable misorientations of 2-3o is also noteworthy (Fig. 4). The latter effect mirrors the formation of poorly developed LABs in the grain

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interior (examples are circled in Fig. 3b). At relatively large tensile strains (≈0.16-0.20), the {10 1 2} twin-induced peaks tended to be obscure (Figs. 4 & 5d) and the fraction of the {10 1 2} twin boundaries

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rapidly decreased (Table 1). These effects reflect the transformation of twins into regular grains, as discussed in the previous section, and perhaps evidence a gradual suppression of the twinning strain mode. On the other hand, activation of subtle {10 1 1} and {10 1 3} twinning was found (Table 1). Of particular interest was the activation of {10 1 1}  {10 1 2} double twinning5 (Table 1). In magnesium alloys, this mechanism is well accepted to be a precursor of

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cracking and subsequent failure during uniaxial tension [30, 43, 54]. In contrast to the gradual disappearance of the twin boundaries, an abrupt increase of the LAB area was observed (Fig. 4). This perhaps indicates that slip became the dominant strain mode at the large tensile strains. The enhanced slip activity at this stage also follows from Kernel-average-misorientation maps (Supplementary Fig. S5). In this context, the clustering of misorientation axes of the

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LABs near  2 1 1 0  (arrow in Fig. 5d) is of interest. The rotation axes of (low-angle) dislocation boundaries impose a constraint upon the slip modes which are potentially active during deformation and thus may provide useful information about the slip activity. It is believed that the slip activity of the p system rotates the crystalline lattice around the axis wp  bp  n p , where b p is the Burger’s vector and n p is the slip plane normal. Among all slip modes possible in magnesium, only the basal slip results in rotation around the  2 1 1 0  axis. Thus, the concentration of the misorientation axes of LABs near this pole indirectly indicates the operation of the basal slip.

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In accordance with numerous reports in the scientific literature [2, 16-25], material in the as-welded condition had a strong basal texture (Fig. 6a). It is seen that the basal texture is rotated and that this rotation is associated with local orientation of the shear direction at the AS [16]. At low-to-moderate tensile strains ( ≈ 0.06-0.13), formation of a deformation-induced {0001}  1120  texture component was found (Figs. 6b & c).

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This effect was obviously attributable to the profuse {10 1 2} twinning, as noted in the

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previous two sections. Importantly, the activation of all possible twin variants should provide effective texture scattering. Therefore, the development of strong {0001}  1120  orientation indicates the occurrence of a variant selection during

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twinning. This conclusion agrees with the limited number of twin variants revealed within the grains (Figs. 3b & c), as mentioned in Section 3.3.1. It is noteworthy that selection of twin variants is often observed during deformation of magnesium alloys [24, 33-38, 45-54]; the origin of this phenomenon is discussed in Section 4. In friction-stir welded material, this effect has been reported by Liu et al [24].

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After relatively large straining (≈0.20), the remnants of the original basal texture as well as the twin-induced {0001}  1120  component tended to transform into {hkil}  1120  fiber texture (Fig. 6d). Remarkably, texture transformation seems to be most pronounced in the twinned areas (Fig. 6d) and its mechanism is of particular interest. At least partially, the observed reorientation is attributable to the double {10 1 1}  {10 1 2} twinning, as mentioned in the previous section. However, the area

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fraction of the appropriate twin boundaries is low (Table 1), and thus it is unlikely that the contribution of this mechanism was large. On the other hand, the twin-induced {0001}  1120  orientation is characterized by a relatively low Schmid factor for the basal slip, and therefore the {hkil}  1120  fiber texture may originate from operation of non-basal slip. Indeed, the activation of the {1 1 00}  1120  prism slip and the {10 1 1}  1120  pyramidal slip is predicted during uniaxial tension of magnesium at ambient temperature [55]. To verify the above idea, crystal plasticity modelling is required. However, such detailed textural analysis is somewhat out of the scope of the present work, and it is believed to be worthy of a separate study. 4. Discussion Based on the experimental observations summarized in Section 3, two stages of microstructure evolution during transverse tension of friction-stir welded magnesium

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the stir zone edge [19, 22]. It should be emphasized that the twinning was governed by strict variant selection which promoted the formation of a relatively strong {0001}  1120  texture component (Fig. 6c). In magnesium, the variant selection during twinning is well accepted to be

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consistent with Schmid law, i.e. the twins are activated from the variant pair with the highest Schmid factor [24, 33-37, 46-47, 49-54]. In friction-stir welded material, this effect has been demonstrated by Liu et al [24]. Therefore, considering the strong basal texture in the as-welded material (Fig. 6a), the observed variant selection appears to be a well-expected result. On the other hand, the twins in magnesium are often arranged into twin chains, and thus the selection of the twin variants is well established to be additionally influenced by local (i.e., grain-to-grain) strain accommodation [24, 33-37, 45-46, 49-50, 52, 54]. In the present study, however, the twin chains were not restricted by two neighboring grains; instead, they formed a

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macro-scale pattern with two sets oriented by ± 55o to the tension axes (Fig. 3c). It may be suggested, therefore, that the formation of the twin chains (and thus the variant selection) was rather associated with strain accommodation on a macro-scale,

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thus being attributable to the extremely inhomogeneous strain distribution in the stir zone (Fig. 2). This idea, however, requires experimental verification. Also, considering preferential strain concentration in the locations with high Schmid factor for the basal slip (Figs. 1c vs 2b-d), the first stage of the microstructure

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evolution was presumably characterized by the prevalence of this slip mode. In the second stage (true strain > ~0.15), the {10 1 2} twinning was suppressed. Moreover, the twin boundaries were gradually transformed into random (non-twin) HABs, thus disappearing as a structural element. These effects are well known and are usually attributable to an increased activity of slip at high strains. Indeed, textural observations suggest activation of non-basal slip at this stage, as discussed in Section 3.3.3. The increased slip activity gave rise to an abrupt enlargement of the LAB area (Fig. 4) and promoted the formation of the fiber {hkil}  1120  orientation (Fig. 6d). The relatively fast texture transformations observed in the twinned areas (Fig. 6d) indirectly indicate relatively high slip activity within the twins. It is important to emphasize that failure of magnesium alloys during uniaxial tension is often attributed to double {10 1 2}  {10 1 1} twinning [30, 43-44]. The {10 1 1}

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during the first stage of the microstructure evolution (as well as the strict variant selection governing this process) creates preconditions for the subsequent failure. Due to inhomogeneous texture distribution, the {10 1 2} twinning in friction-stir welded magnesium is most feasible at the stir zone extremity [19, 22]. This perhaps explains the abundantly reported fracture of the welded material in this

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microstructural region [2, 10, 17-20, 22-28].

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5. Summary In this work, microstructural changes occurring during transverse tension tests of friction-stir welded magnesium alloy AZ31 were systematically studied to clarify the microstructure-properties relationship. To this end, digital image correlation was used for accurate strain measurements and EBSD was employed for thorough examination of microstructure evolution and texture development. The main conclusions derived from this study are as follows.

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As expected, plastic strain in the stir zone was preferentially concentrated in locations with favorable crystallographic orientation for basal slip, i.e., at the AS and the RS. In both locations, the microstructure evolution was found to be a relatively complex process consisting of two stages. In the first stage (true strain <~0.15), the microstructure evolution was characterized by profuse {10 1 2} twinning. Importantly,

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this process was governed by strict variant selection promoting the formation of a strong {0001}  1120  texture component. During the second stage (true strain >~0.15), {10 1 2} twinning was suppressed and slip became the dominant strain mode. This gave rise to abrupt enlargement of the LAB area and resulted in development of the fiber {hkil}  1120  texture. Twin-induced {0001}  1120  orientation was shown to be prone to subsequent {10 1 2}  {10 1 1} double twinning which, in turn, is known to typically lead to final failure. Therefore, the profuse {10 1 2} twinning occurring during the first stage of the microstructure evolution (as well as variant selection governing this process) creates preconditions for fracture of magnesium friction-stir welds. Acknowledgments The authors are grateful to Mr. A. Honda for technical assistance. One of the authors

ACCEPTED MANUSCRIPT (S. Mironov) would like to express his hearty thanks to Tohoku University for providing a scientific fellowship. References [1] R.S. Mishra, Z.Y. Ma, Friction stir welding and processing, Mater. Sci. Eng. R 50 (2005) 1–78, doi:10.1016/j.mser.2005.07.001 [2] S.H.C. Park, Y.S. Sato, H. Kokawa, Effect of micro-texture on fracture location in friction stir weld of alloy

AZ61

during

tensile

test,

Scr.

Mater.

49

(2003)

PT

Mg

http://dx.doi.org/10.1016/S1359-6462(03)00210-0.

161–166,

[3] W. Woo, H. Choo, M.B. Prime, Z. Feng, B. Clausen, Microstructure, texture and residual stress in a AZ31B

magnesium

alloy,

Acta

http://dx.doi.org/10.1016/j.actamat.2007.12.020

Mater.

56

(2008)

RI

friction-stir-processed

1701–1711,

subsequent

compression,

SC

[4] C.J. Lee, J.C. Huang, X.H. Du, Improvement of yield stress of friction-stirred Mg– Al–Zn alloys by Scr.

Mater.

(2007)

875–878,

NU

http://dx.doi.org/10.1016/j.scriptamat.2007.01.041.

56

[5] W. Yuan, R.S. Mishra, B. Carlson, R.K. Mishra, R. Verma, R. Kubic, Effect of texture on the mechanical behavior of ultrafine grained magnesium alloy, Scr. Mater. 64 (2011) 580–583,

MA

http://dx.doi.org/10.1016/j.scriptamat.2010.11. 052.

[6] W. Xunhong, W. Kuaishe, Microstructure and properties of friction stir butt welded AZ31 magnesium alloy, Mater. Sci. Eng. A 431 (2006) 114–117, http:// dx.doi.org/10.1016/j.msea.2006.05.128.

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[7] W. Yuan, R.S. Mishra, Grain size and texture effects on deformation behavior of AZ31 magnesium

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alloy, Mater. Sci. Eng. A 558 (2012) 716–724, http://dx.doi. org/10.1016/j.msea.2012.08.080. [8] W. Yuan, R.S. Mishra, B. Carlson, R. Verma, R.K. Mishra, Material flow and microstructural evolution during friction stir spot welding of AZ31 magnesium alloy, Mater. Sci. Eng. A 543 (2012) 200–209, http://dx.doi.org/10.1016/j. msea.2012.02.075.

grained

CE

[9] A. Mohan, W. Yuan, R.S. Mishra, High strain rate superplasticity in friction stir processed ultrafine Mg–Al–Zn

alloys,

Mater.

Sci.

Eng.

A

562

(2013)

69–

76,

AC

http://dx.doi.org/10.1016/j.msea.2012.11.026. [10] R. Xin, L. Sun, D. Liu, Z. Zhou, Q. Liu, Effect of subsequent tension and annealing on microstructure evolution and strength enhancement of friction stir welded Mg alloys, Mater. Sci. Eng. A 602 (2014) 1-10, doi:10.1016/j.msea.2014.02.062 [11] R. Xin, D. Liu, B. Li, L. Sun, Z. Zhou, Q. Liu, Mechanisms of fracture and inhomogeneous deformation on traverse tensile test of friction-stir-processed AZ31 Mg alloy, Mater. Sci. Eng. A 565 (2013) 333–341, http://dx.doi.org/ 10.1016/j.msea.2012.12.053. [12] S. Mironov, T. Onuma, Y.S. Sato, H. Kokawa, Microstructure evolution during friction-stir welding of AZ31 magnesium alloy, Acta Mater. 100 (2015) 301-312; doi:10.1016/j.actamat.2015.08.066 [13] U.F.H.R. Suhuddin, S. Mironov, Y.S. Sato, H. Kokawa, C.-W. Lee, Grain structure evolution during friction-stir

welding

of

AZ31

magnesium

alloy,

Acta

Mater.

57

(2009)

5406-5418;

ACCEPTED MANUSCRIPT doi:10.1016/j.actamat.2009.07.041. [14] S. Mironov, Q. Yang, H. Takahashi, I. Takahashi, K. Okamoto, Y.S. Sato, H. Kokawa, Specific character of material flow in near-surface layer during friction stir processing of AZ31 magnesium alloy, Metall. Mater. Trans. A 41 (2010) 1016-1024; doi:10.1007/s11661-009-0158-x. [15] S. Mironov, Y. Motohashi, R. Kaibyshev, H. Somekawa, T. Mukai, K. Tsuzaki, Development of grain structure during friction stir welding of a ZK60 magnesium alloy, Mater. Trans. 50 (2009) 610-617; doi:10.2320/matertrans.MRA2008192.

PT

[16] S.H.C. Park, Y.S. Sato, and H. Kokawa, Basal plane texture and flow pattern in friction stir weld of a magnesium alloy, Metall. Mater. Trans. A 34 (2003) 987-994; doi:10.1007/s11661-003-0228-4 [17] W. Woo, H. Choo, D.W. Brown, P.K. Liaw, Z. Feng, Texture variation and its influence on the

RI

tensile behavior of a friction-stir processed magnesium alloy, Scripta Mater. 54 (2006) 1859–1864; doi:10.1016/j.scriptamat.2006.02.019

SC

[18] L. Commin, M. Dumont, R. Rotinat, F. Pierron, J.-E. Masse, Influence of the microstructural changes and induced residual stresses on tensile properties of wrought magnesium alloy friction stir

NU

welds, Mater. Sci. Eng. A 551 (2012) 288-292; doi:10.1016/j.msea.2012.05.021 [19] R. Xin, D. Liu, Z. Xu, B. Li, Q. Liu, Changes in texture and microstructure of friction stir welded Mg

178-187; doi:10.1016/j.msea.2013.06.005

MA

alloy during post-rolling and their effects on mechanical properties, Mater. Sci. Eng. A 582 (2013)

[20] J. Chen, H. Fujii, Y. Sun, Y. Morisada, R. Ueji, Fine-grained Mg-3Al-1Zn alloy with randomized

doi:10.1016/j.msea.2013.05.044

D

texture in the double-sided friction stir welded joints, Mater. Sci. Eng. A 580 (2013) 83-91;

PT E

[21] D. Liu, R. Xin, L. Sun, Z. Zhou, Q. Liu, Influence of sampling design on tensile properties and fracture behavior of friction stir welded magnesium alloys,. Mater. Sci. Eng. A576 (2013) 207–216; doi:10.1016/j.msea.2013.03.037

[22] R. Xin, D. Liu, B. Li, L. Sun, Z. Zhou, Q. Liu, Mechanisms of fracture and inhomogeneous

CE

deformation on traverse tensile tests of friction-stir-processed AZ31 Mg alloy, Mater. Sci. Eng. A 565 (2013) 333-341; doi:10.1016/j.msea.2012.12.053

tension

AC

[23] D. Liu, R. Xin, Y. Xiao, Z. Zhou, Q. Liu, Strain localization in friction stir welded magnesium during and

compression

deformation,

Mater.

Sci.

Eng.

A

609

(2014)

88-94;

doi:10.1016/j.msea.2014.04.089 [24] D. Liu, R. Xin, H. Yu, Z. Liu, X. Zheng, Q. Liu, Comparative examinations on the activity and variant selection during tension and compression of magnesium alloys, Mater. Sci. Eng. A 658 (2016) 229-237; doi:10.1016/j.msea.2016.01.098 [25] Z. Liu, R. Xin, D. Liu, X. Shu, Q. Liu, Textural variation in triple junction region of friction stir welded Mg alloys and its influence on twinning and fracture, Mater. Sci. Eng. A 658 (2016) 185-191; doi:10.1016/j.msea.2016.01.097 [26] N. Afrin, D.L. Chen, X. Cao, M. Jahazi, Microstructure and tensile properties of friction stir welded AZ31B magnesium alloy, Mater. Sci. Eng. A 472 (2008) 179-186; doi:10.1016/j.msea.2007.03.018

ACCEPTED MANUSCRIPT [27] S.M. Chowdhury, D.L. Chen, S.D. Bhole, X. Cao, Tensile properties of friction stir welded magnesium alloy: effect of pin tool thread and weld pitch, Mater. Sci. Eng. A 527 (2010) 6064-6075; doi:10.1016/j.msea.2010.06.012 [28] J. Yang, B.L. Xiao, D. Wang, Z.Y. Ma, Effects of heat input on tensile properties and fracture behavior of friction stir welded Mg-3Al-1Zn alloy, Mater. Sci. Eng. A 527 (2010) 708-714; doi:10.1016/j.msea.2009.09.044 [29] M.R. Barnett, Twinning and the ductility of magnesium alloys Part I: “Tension” twins, Mater. Sci.

PT

Eng. A 464 (2007) 1–7, doi:10.1016/j.msea.2006.12.037 [30] M.R. Barnett, Twinning and the ductility of magnesium alloys. Part II. “Contraction” twins, Mater. Sci. Eng. A 464 (2007) 8-16; doi:10.1016/j.msea.2007.02.109

RI

[31] H.A. Bruck, S.R. McNeill, M.A. Sutton, W. H. Peters III, Digital image correlation using Newton-Raphson method of partial differential correction, Experimental Mechanics, 29-3 (1989),

SC

261-267; doi:10.1007/BF02321405

[32] S. Yoneyama, Basic principle of digital image correlation for in-plane displacement and strain Advanced

Composite

Materials,

NU

measurement,

doi:10.1080/09243046.2015.1129681

25-2

(2016),

105-123;

[33] R. Xin, C. Guo, Z. Xu, G. Liu, X. Huang and Q. Liu, Characteristics of long {10-12} twin bands in rolling

of

a

magnesium

alloy,

MA

sheet

Scripta

Mater.

74

(2014)

96-99,

http://dx.doi.org/016/j.scriptamat.2013.11.008

[34] B. Song, R. Xin, Y. Liang, G. Chen, Q. Liu, Twinning characteristic and variant selection in

D

compression of a pre-side-rolled Mg alloy sheet, Mater. Sci. Eng. A 614 (2014) 106-115,

PT E

http://dx.doi.org/10.1016/j.msea.2014.07.026

[35] C. Guo, R. Xin, C. Ding, B. Song, Q. Liu, Understanding of variant selection and twin patterns in compressed Mg alloy sheets via combined analysis of Schmid factor and strain compatibility factor, Mater. Sci. Eng. A 609 (2014) 92-101, http://dx.doi.org/10.1016/j.msea.2014.04.103

CE

[36] R. Xin, C. Ding, C. Guo, Q. Liu, Crystallographic analysis on the activation of multiple twins in rolled AZ31 Mg alloy sheets during uniaxial and plane strain compression, Mater. Sci. Eng. A 652

AC

(2016) 42-50, http://dx.doi.org/10.1016/j.msea.2015.11.033 [37] H. Chen, T. Liu, D. Hou, D. Shi, Study on the paired twinning behavior in a hot rolled AZ31 magnesium alloy via interrupted in situ compression, Mater. Sci. Eng. A 667 (2016) 402-408, http://dx.doi.org/10.1016/j.msea.2016.05.006 [38] A. Fernandez, A. Jerusalem, I. Gutierrez-Urrutia, M.T. Perez-Prado, Three-dimensional investigation of grain boundary–twin interactions in a Mg AZ31 alloy by electron backscatter diffraction and

continuum

modeling,

Acta

Mater.

61

(2013)

7679-7692,

http://dx.doi.org/10.1016/j.actamat.2013.09.005 [39] E.V. Nesterova, V.V. Rybin, Mechanical twinning and fragmentation of commercial purity titanium during large plastic deformations, Phys. Met. Metall. 59 (1985) 169–179.

ACCEPTED MANUSCRIPT [40] S. Mironov, Y.S. Sato, H. Kokawa, Structural response of superaustenitic stainless steel to friction stir welding, Acta Mater., 59 (2011) 5472-5481; http://dx.doi.org/10.1016/j.actamat.2011.05.021 [41] S. Mironov, K. Inagaki, Y.S. Sato & H. Kokawa, Development of grain structure during friction-stir welding

of

Cu–30Zn

brass,

Phil.

Mag

94

(2014)

No.

27

pp.

3137-3147,

doi:10.1080/14786435.2014.951712 [42] M.D. Nave, M.R. Barnett, Microstructures and textures of pure magnesium deformed in plane-strain

compression,

Scr.

Mater.

51

(2004)

881–885,

http://dx.doi.org/10.1016/j.scriptamat.2004.07.002.

PT

[43] W.H. Hartt and R.E. Reed-Hill, Internal deformation and fracture of second-order {10 1 1}  {10 1 2} twins in magnesium, Trans. Met. Soc. AIME 242 (1968) 1127-1132.

formation

in

AZ31

magnesium

alloy,

RI

[44] D. Ando, J. Koike, Y. Sutou, Relationship between deformation twinning and surface step Acta

58

(2010)

4316-4324;

SC

doi:10.1016/j.actamat.2010.03.044

Mater.

[45] J.J. Jonas, S. Mu, T. Al-Samman, G. Gottstein, L. Jiang, E. Martin, The role of strain

NU

accommodation during the variant selection of primary twins in magnesium, Acta Mater 59 (2011) 2046-2056, http://dx.doi.org/10.1016/j.actamat.2010.12.005

[46] S. Mu, J.J. Jonas, G. Gottstein, Variant selection of primary, secondary and tertiary twins in a

MA

deformed Mg alloy, Acta Mater. 60 (2012) 2043-2053, http://dx.doi.org/10.1016/j.actamat.2012.01.014 [47] S. Godet, L. Jiang, A.A. Luo and J.J. Jonas, Use of Schmid factors to select extension twin variants

in

extruded

magnesium

alloy

tubes,

Scripta

Mater.

55

(2006)

1055–1058,

D

doi:10.1016/j.scriptamat.2006.07.059

PT E

[48] Y. Pei, A. Godfrey, J. Jiang, Y.B. Zhang, W. Liu, Q. Liu, Extension twin variant selection during uniaxial compression of a magnesium alloy, Mater. Sci. Eng. A 550 (2012) 138– 145, https://dx.doi.org/10.1016/j.msea.2012.04.046 [49] B.L. Wu, Y.D. Zhang, G. Wan, M. Humbert, F. Wagner, C. Esling, Primary twinning selection with

CE

respect to orientation of deformed grains in ultra-rapidly compressed AZ31 alloy, Mater. Sci. Eng. A 541 (2012) 120-127, https://dx.doi.org/10.1016/j.msea.2012.02.012

re-twinning

AC

[50] A. Chapuis, Y. Xin, X. Zhou, Q. Liu, {10-12} Twin variants selection mechanisms during twinning, and

detwinning,

Mater.

Sci.

Eng.

A

612

(2014)

431-439,

http://dx.doi.org/10.1016/j.msea.2014.06.088 [51] H.T. Huang, A. Godfrey, J.P. Zheng, W. Liu, Influence of local strain on twinning behavior during compression

of

AZ31

magnesium

alloy,

Mater.

Sci.

Eng.

A

640

(2015)

330-337,

http://dx.doi.org/10.1016/j.msea.2015.06.009 [52] X. Liu, J.J. Jonas, B.W. Zhu, T. Wang, L.X. Li, Variant selection of primary extension twins in AZ31 magnesium

deformed

at

o

400 C,

Mater.

http://dx.doi.org/10.1016/j.msea.2015.10.020

Sci.

Eng.

A

649

(2016)

461-467,

ACCEPTED MANUSCRIPT [53] J. Jiang, A. Godfrey, W. Liu and Q. Liu, Identification and analysis of twinning variants during compression,

of

a

Mg–Al–Zn

alloy,

Scripta

Mater.

58

(2008)

122–125,

doi:10.1016/j.scriptamat.2007.09.047 [54] E. Martin, L. Capolungo, L. Jiang, J.J. Jonas, Variant selection during secondary twinning in Mg–3%Al, Acta Mater. 58 (2010) 3970–3983, http://dx.doi.org/10.1016/j.actamat.2010.03.027 [55] S.-B. Yi, C.H.J. Davis, H.-G. Brokmeier, R.E. Bolmaro, K.U. Kainer, J. Homeyer, Deformation and texture evolution in AZ31 magnesium alloy during uniaxial loading, Acta Mater. 54 (2006) 549-562;

AC

CE

PT E

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MA

NU

SC

RI

PT

https://dx.doi.org/10.1016/j.actamat.2005.09.024

ACCEPTED MANUSCRIPT Figure captions

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Figure 1. Low magnification overview of a transversal cross section of the stir zone: (a) optical image, (b) EBSD map showing orientations relative to the welding direction (color code triangle is shown in the bottom right corner), (c) EBSD map showing Schmidt factors for basal slip for transverse tensile tests (color code is shown in the bottom right corner). WD, ND and TD are welding, normal and transverse directions, respectively; AS and RS are advancing and retreating side, respectively. Figure 2. The distribution of extensional strain in the stir zone after tensile elongation of (a) 0.5%, (b) 4.9%, (c) 9.9% and (d) 16.1% (i.e., immediately before

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fracture). In all cases, the advancing side is on the left and the retreating side is on the right. Tension direction is horizontal.

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Figure 3. EBSD grain-boundary maps showing microstructure at the AS of the stir zone in (a) the as-welded state and after local true strain of (b) 0.06, (c) 0.13, and (d) 0.20 (failure). LABs, HABs and {10 1 2} twin boundaries are depicted as red, black

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and blue lines, respectively. The reference frame for all maps is shown in the bottom right corner of (a). Tension direction is horizontal. Note the difference in scales. See text for details Figure 4. Effect of tensile strain on misorientation-angle distribution at the AS of the stir zone

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Figure 5. Effect of tensile strain on misorientation-axis distribution at the AS of the stir zone in (a) as-welded state, and after true strain of (b) 0.06, (c) 0.13, and (d) 0.20 (failure) Figure 6. The 0001 and {1120} pole figures showing texture at the AS of the stir

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zone in (a) as-welded state and after true strain of (b) 0.06, (c) 0.13, and (d) 0.20 (failure). The reference frame for the pole figures is shown in the bottom right corner.

ACCEPTED MANUSCRIPT Table 1. Evolution of twin boundary fraction at AS of the stir zone (within 5-degree tolerance) Twin mode

True local strain

Misorientation [42]

0

0.06

0.09

0.13

0.16

0.20 (failure)

Primary twinning {10 1 2} 86o  2 1 1 0  {10 1 1} 56o  2 1 1 0 

0.3 0.0

37.1 0.0

50.6 0.0

52.2 0.0

37.6 0.0

64o  2 1 1 0 

0.0

0.0

0.1

0.1

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0.0 0.0

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1.0 0.4

1.0 0.4

38  2 1 1 0 

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0.0

0.1

30  2 1 1 0 

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0.1 0.0 0.0

0.2 0.2 0.1

0.4 0.2 0.1

{10 1 3}

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{10 1 2} twin variants (10 1 2)  (01 1 2)

60o  10 1 0 

(10 1 2)  (0 1 12)

60  8170  o

Double twinning o

o

{10 1 1}  {10 1 2}

67 o  3472  70o  14;7;7;3 

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Note: The relatively high twin boundary fractions are highlighted with gray

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9.9 0.5 1.4

1.9 1.4

0.6 0.7

0.0

0.0

0.4

0.2 0.2 0.1

0.1 0.1 0.1

0.4 0.3 0.2

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texture.  Twin-induced

{0001}  1120 

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 Microstructural changes in stir zone concentrate on advancing- and retreating sides.  Microstructure evolution consists of 2 stages: (1) twin-induced and (2) slip-induced.  Twinning is governed by strict variant selection giving rise to {0001}  1120  orientation promotes double

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twinning which leads to final failure.

{10 1 2}  {10 1 1}