Microstructure evolution and mechanical properties of an Mg–Gd alloy subjected to surface mechanical attrition treatment

Materials Science & Engineering A 630 (2015) 146–154

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Microstructure evolution and mechanical properties of an Mg–Gd alloy subjected to surface mechanical attrition treatment X.Y. Shi a, Y. Liu b,c, D.J. Li a, B. Chen d, X.Q. Zeng a,n, J. Lu c,e,nn, W.J. Ding a a

The State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai 200240, PR China Key Laboratory of Near Net Forming of Jiangxi Province, Nanchang University, Nanchang 330031, PR China c Department of Mechanical and Biomedical Engineering, City University of Hong Kong, Kow-loon, Hong Kong, PR China d School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, PR China e Centre for Advanced Structural Materials, City University of Hong Kong, Shenzhen Research Institute, Shenzhen 518057, PR China b

art ic l e i nf o

a b s t r a c t

Article history: Received 9 October 2014 Received in revised form 29 December 2014 Accepted 8 February 2015 Available online 17 February 2015

Nanometer-sized grains (around 50–100 nm) were generated in the surface layer (0–30 μm) of an Mg– 3Gd alloy by means of surface mechanical attrition treatment (SMAT) at room temperature. The deformation process and the formation mechanism of nano grains were investigated by using transmission electron microscopy. The results show that twinning dominates the initial stage of the plastic deformation when the dislocation slips are obstructed. And till all the coarse grains are divided into substructures by twin–twin interactions and twin-dislocation arrays intersections, dislocation slips and stacking faults begin to play an important role in impelling subgrains to nanograins by lattice rotating through dislocation arrays slipping. Dynamic rotation recrystallization is the primarily formation mechanism of the nanocrystalline of Mg–3Gd alloy by SMAT, which is distinct from the migration recrystallization normally observed in severe plastic deformation process. & 2015 Elsevier B.V. All rights reserved.

Keywords: Surface mechanical attrition treatment Mg–3Gd alloy Microstructure evolution Mechanical properties Dynamic rotation recrystallization

1. Introduction Grain refinement is an effective method to improve the strength of metal materials so that substantial severe plastic deformation (SPD) methods are conducted to produce ultrafine grains and nanocrystalline grains in bulk materials over the past decades [1]. Fang et al. [2] found that free-standing nano-grained (NG) metals usually exhibit a high strength but a very low elongation, while the gradient NG copper film confined by a coarse-grained (CG) substrate with a grain-size transition from NG to CG exhibits a twice higher yield strength and a tensile plasticity comparable to that of the CG substrate. Wu et al. [3] believed that the grain-size gradient under uniaxial tension induces a macroscopic strain gradient and converts the applied uniaxial stress to multiaxial stresses due to the evolution and incompatible deformation along the gradient depth, which renders a unique extra strain hardening. Surface mechanical attrition treatment (SMAT) has been proven

n

Corresponding author. Tel.: þ 86 21 54742301; fax: þ86 21 34203730. Corresponding author at: Department of Mechanical and Biomedical Engineering, City University of Hong Kong, Kow-loon, Hong Kong, PR China. Tel.: þ 852 3442 9811; fax: þ852 3442 0295. E-mail addresses: [email protected] (X.Y. Shi), [email protected] (Y. Liu), [email protected] (D.J. Li), [email protected] (B. Chen), [email protected] (X.Q. Zeng), [email protected] (J. Lu), [email protected] (W.J. Ding). nn

http://dx.doi.org/10.1016/j.msea.2015.02.009 0921-5093/& 2015 Elsevier B.V. All rights reserved.

to be an effective way to obtain gradient microstructure with surface layer containing nano grains [4–6]. In addition, optimization of the surface structure would greatly enhance the overall properties of materials since most failures occur on the surface, such as behavior of fatigue [7,8], corrosion and wear [9,10]. It is the key point to understand the grain refinement mechanism during SMAT in order to get NG layer by controlling the parameters of SMAT. There are two significant factors governing the grain refinement process during SMAT deformation: the stacking fault energy (SFE) and the number of slip systems in the metals [11]. For face-centered cubic (fcc) and body-centered cubic (bcc) metals, dislocation slips are in charge of the refinement process if SFE of the metal is relatively high while twinning plays an important role in the initial deformation stage if SFE of the metal is medium or low [12–15]. For hexagonal close packed (hcp) pure Co [16] that has a low SFE and α-Ti [17] with a very high SFE, twinning appears to be the dominant refinement mechanism at the early deformation stage in order to satisfy the von Mises criterion. Therefore, as a typical hcp structured metal, Mg and its alloys have SFEs within the range of 60–78 mJ/m2 and are expected to be refined by combination of dislocation slips and twins, which is testified by the AZ91D Mg alloy investigated by Sun et al. [11]. However, Wei et al. [18] found that the grain refinement of SMAT AZ91D Mg alloy is attributed to the activity of dislocation but without twins. The two contradictory experiment results indicate that the refinement mechanism of Mg alloy during SMAT process is not very

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clear or well-illustrated by the two factors mentioned above. This excites our interest to carefully and systematically investigate the twinning criteria of Mg alloy and other factors effecting grain refinement mechanism during SMAT process. Since the activation of deformation twinning is affected significantly by deformation temperature [19–23], strain rate [19–21,24–27], stacking fault energy [19,20], generalized planar fault energies [28–31] and grain size [32–40]. These will be investigated and discussed in the following future work. SMAT is regarded as a mechanical deformation method but quite different from SPD in Mg alloy system. For Mg alloys with poor deformability at room temperature [41], SPD methods are always applied at high temperatures and low strain rates (lower than 10 s  1) while SMAT method are implemented at ambient temperature and extremely high strain rates (104–105 s  1) [42–44]. The difference is probably responsible for the truth that SMAT could be used to produce a nanocrystalline structure which SPD methods can hardly obtain in Mg alloys. Moreover, dynamic recrystallization (DRX) plays both important but slightly different roles in these two means according to the different results. However, the details of DRX during SMAT process have not been observed or discussed before. In this paper, a model alloy of Mg–3Gd binary alloy with an average grain size of 43 μm was designed to form the gradient structure by SMAT. The deformation mechanism and DRX process of nanocrystalline grain was explored in detailed.

147

test was conducted in the cross section from the surface to center of sample with a load of 50 g and a holding time of 15 s. The microstructure evolution was observed using an optical microscope (OM, Zeiss, Axio Vert.A1), electron backscattering diffraction (EBSD) and transmission electron microscopes (TEM, JEOL, JEM-2100 and JEM-2100F) operating at 200 kV. The OM samples were mechanically polished and then etched in an acetic acid glycol solution. The TEM samples were electro-polished using a twin-jet machine (Struers, TenuPol-5) in an ethanol solution containing 4 vol% perchloric acid at a voltage of 50 V and a temperature of 233 K (  40 1C).

3. Results As shown in Fig. 1, the original as-extruded Mg–3Gd alloy with homogeneous equiaxed grains contains a single α-Mg phase with hcp structure. The average grain size is around 43 72 μm and all grains are randomly oriented without obvious texture (texture intensity 5) and without twins. In order to avoid the temperature severe increasing on the surface of the alloy due to the mechanical heat, the SMAT time is controlled below 4 min. And the thin specimen is designed aiming at thin-wall castings in the future industrial applications.

3.1. Mechanical properties 2. Experimental methods The extruded binary Mg–3Gd alloy (with chemical composition of Gd is 3.25 wt%) was cut into plate specimens with dimensions of 80 (along extrusion direction)  15  1 mm3. The SMAT process was performed at room temperature for 2 and 4 min with a vibrating frequency of 20 KHz. The detailed procedures of the SMAT process have been published in the previous literature [42,43]. It is worthy to note that there was no obvious temperature increasing during these SMAT processes with treatment time not more than 4 min although the localized heat generation might be inevitable. The bone-like tensile specimens with dimensions of 3 mm width, 1, 0.8 and 0.7 mm thickness (for original alloy, SMAT-2 min alloy and SMAT-4 min alloy, respectively) and 18 mm gauge length were tested at ambient temperature using a Zwick-100 kN material testing machine with a crosshead speed of 0.3 mm/min, corresponding to strain rate of 2.8  10  4 s  1. The SMA-Treated surfaces as the width  length area of tensile specimens were kept and the thickness  length area of the specimens cut by electric discharge machining (EDM) were polished before tensile testing. An extensometer was used during mechanical testing, and the stress strain curves obtained were engineering stress-engineering strain. Vickers micro-hardness

Fig. 2 shows the mechanical properties of as-extruded and SMATed (2 min and 4 min) Mg–3Gd alloy. The tensile testing results shown in Fig. 2(a) indicate that tensile yield strengths (TYS) of SMAT Mg–3Gd alloys for 2 min and 4 min (128.2 MPa and 152.4 MPa respectively) are much higher than that of initial extruded Mg–3Gd alloy (70.0 MPa), the ultimate tensile strengths (UTS) of SMAT alloys are about 50 MPa higher than that of original alloy. However, the significant increase of the strengths comes along with the sharp decrease of the elongation for Mg–3Gd alloy after the SMAT process. This is mainly attributed to the strain hardening produced by the accumulation and interaction of dislocations. In addition, the TYS and UTS increase more and elongation decreases more along with the increase of the treatment time. As can be seen in Fig. 2(b), the variation of microhardness tested on the cross section of the SMAT alloys shows a “U” type trend, which means the micro-hardness of the topmost layer is much higher than that in the central region. And the highest hardness value of the SMAT-4 min sample (89 HV) is around 7 HV (or 8%) higher than that of the SMAT-2 min alloy (82 HV), while the lowest points of these two samples are almost the same. Furthermore, even the lowest micro-hardness value of

Fig. 1. (a) Optical microstructure and (b) EBSD image of as-extruded Mg–3Gd alloy.

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Fig. 2. Mechanical properties of extruded Mg–3Gd alloy, SMA-Treated alloys for 2 min and 4 min: (a) tensile testing curves at ambient temperature; (b) variation of microhardness with distance from the topmost surfaces to the center layer.

the SMAT alloys in the central region is 12 HV higher than that of the as-extruded alloy with value of 49 HV. 3.2. Optical microscopy observations The cross sectional optical microstructures of the SMAT specimens for 2 min and 4 min are given in Fig. 3. It can be seen that a large number of twins appear in the SMAT alloys. Grain boundaries in the severely deformed (SD) zone are so indistinct that the grain size could not be ascertained by OM. However, the average grain size in the center regions of the two samples is as the same as the original alloy due to the slight deformation. Compared to SMAT2 min sample, the thickness of severely deformed zone in SMAT4 min specimen increases a little, corresponding to parabola-like (x ¼y2, y40) tendency instead of linear growth. Since the central microstructure of the two samples are almost the same, the difference of the severe deformation zone between the two specimens is responsible for the tensile testing results in which the TYS increment of SMAT-4 min sample compared to SMAT-2 min sample is much less than that of SMAT-2 min specimen compared to extruded alloy. Moreover, the increments in micro-hardness of central regions in SMAT specimens compared to original alloy shown in Fig. 2(b) are mainly attributed to the plenty of twins which provide many boundaries to impede dislocation slips. And the remarkable promotion of the hardness is caused by the grain refinement as well as strain hardening due to the residual stress. 3.3. TEM investigations of the grain refinement process It is necessary to examine the microstructure evolution under different plastic strain during SMAT process in order to analyze the grain refinement mechanism. In the same SMAT specimen, the sections at different depths represent the different strain situations since the deformation always begins from the top surface layer and extends to adjacent substrate regions [11]. Therefore, SMAT-4 min alloy is chosen to characterize the microstructure evolution and the layers at different depths from the surface are taken out to provide information about the strain-induced grain refinement process. 3.3.1. The deformed region Since the thickness of the SMAT-4 min sample is around 700 μm, the layer that is 300–400 μm below the treated surface is regarded as the center layer of the sample whose bright TEM images are shown in Fig. 4. As the strain and strain rate sustained by this layer are the lowest among all layers in the sample, there

SD layer

SMAT Direction

SMAT Direction

SD layer

Fig. 3. Cross-sectional optical micrographs of the SMA-Treated Mg–3Gd alloy for (a) 2 min and (b) 4 min (SD layer means severely deformed layer).

are a few twins and a lot of dislocations appear to coordinate the plastic strain. Some twins appear as shown in Fig. 4, and they are supposed to interact with dislocations and hinder dislocation slips. It is known that deformation twins are the supplement of dislocation slips for strain accommodation in Mg alloys and the twin boundaries become into another kind of block for dislocation slips in turn. Fig. 5 shows the bright field TEM images of the region that is 150– 200 μm below the treated surface. It can be seen from Fig. 5(a) that the density of twins goes up obviously owing to the increase of strain and strain rate in this layer, and the thicknesses of twins range from around 100 nm to 700 nm. At the same time, some twin boundaries (TBs) start to curve, which might be in the charge of the accumulation of dislocations at TBs [45] and twin growth by TBs moving [46]. In addition, the bent incoherent TBs also act as sources of emitting dislocations [47,48]. As shown in Fig. 5(b), some deformation twins derive from grain boundaries and even double twin forms in this region (red line), which makes contribution to the subdivision of twins. Along with strain increasing, a large number of dislocations are accumulated inside the twin lamellae leading to the nonuniform orientations of lamellae and divide them into several irregular substructures, as given in Fig. 5(c). There are many areas with minor misorientation, such as A and C, D and F, E and G, as shown in Fig. 5 (c), since the density of dislocations in the SMAT Mg alloy is lower

X.Y. Shi et al. / Materials Science & Engineering A 630 (2015) 146–154

than that in bcc and fcc treated metals containing dislocation walls and dislocation tangles [49,50]. Fig. 6 shows the bright field TEM images of the layer that is 100– 150 μm below the treated surface. Twin lamellae lying at different directions encountered and interacted with each other and then they are hindered by each other. However, some parts of twin lamellae changed their orientations because of the partial intersection between twins under the higher strain and strain rate. The mechanism for a twin–twin intersection is based on the disclination character of twins and the disclinations are linear defects with a rotational displacement field [51]. Twin–twin intersection needs a higher identical twinning

Twins

149

shear [52] which could be probably provided by a high-strain-rate deformation due to the high kinetic energy. In addition, twin–twin interactions were found to correlate with mechanical hardening [53–55]. The junctions could retard detwinning due to the unfavorable dissociation of twin–twin boundary (TTB) dislocations [56] and would lead to the grain refinement by dividing large grains into sub structures with slightly different orientations. Interacted twins are frequently observed in this layer, which indicates that interaction between twin lamellae plays an important role in grain refinement. 3.3.2. The substructure region Above the deformed region comes the substructure region, as a result of the elevated stress and strain compared with the deformed region. Fig. 7 shows the bright field TEM image of the layer which is 30–80 μm below the treated surface and the corresponding selected area electron diffraction (SAED) patterns. Subgrains with size of 100–200 nm are formed in this layer and still a lot of dislocations inside the subgrains. The SAED pattern shown in Fig. 7(b) was obtained using a large selected area aperture, which selects an area with a diameter of 1.02 μm, and it suggests that the grain-like microstructure shown in Fig. 7(a) are substructures with almost the same orientation and low angle boundaries since the SAED pattern is not a typical circle. However, the spots are not very sharp and elongated or overlapped to some extent, which indicates that the original large grains and twin lamellae start being divided into smaller structures with different orientations.

500 nm

Fig. 4. Bright field TEM images of the layer that is 300–400 μm below the treated surface of SMAT-4 min Mg–3Gd alloy.

3.3.3. The nanocrystalline region The nano-grained region corresponds to the layer less than 30 μm below the topmost surface in this alloy. Fig. 8 gives a typical example of this region. The SAED pattern shown in Fig. 8(b) was also obtained

H

D

E I

GB

F A

G B C

Double Twin

500 nm

500 nm

Fig. 5. Bright field TEM images of (a) parallel twins; (b) double twin; (c) dislocations in twins in the region that is 150–200 μm below the treated surface of SMAT-4 min Mg–3Gd alloy.

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500 nm

500 nm

Fig. 6. Bright field TEM images of interlaced lamellae in the region that is 100–150 μm below the treated surface of SMAT-4 min Mg–3Gd alloy.

Fig. 7. (a) Bright field TEM image of subgrains in the layer that is 30–80 μm below the treated surface of SMAT-4 min Mg–3Gd alloy; (b) corresponding SAED pattern of (a).

Fig. 8. (a) Bright field image of nano grains in the topmost layer with the thickness of 30 μm; (b) corresponding SAED pattern of (a).

using a large selected area aperture, which selects an area with a diameter of 1.02 μm. The continuous diffraction rings in the selected area electron diffraction pattern in Fig. 8(b) demonstrate that randomly oriented grains are obtained. The grain size determined from the TEM image given in Fig. 8(a) ranges from 50 to 100 nm.

Figs. 9 and 10 show the lattice defects in the nanograins. The bright field TEM image of subgrains microstructure taken from the substructure layer is shown in Fig. 9(a) from which it can be seen that the circled structure can be regarded as a nanograin since its orientation is different from the surrounding grains and its grain

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151

Basal plane (0001)

Stacking faults

Fig. 9. (a) Bright field TEM image of nano subgrains; (b) magnified image of a grain circled in (a); (c) SAED patterns of the grain in (b) with electron beam // [2110]; (d) HRTEM image of the rectangular area marked in (b); (e) magnified image of the square region marked in (d) (the dash line represents the missing half plane which results in stacking fault; the inset offers schematic illustration of stacking faults).

size is slightly bigger than 100 nm. The magnified image and the SAED pattern of this grain are given in Fig. 9(b) and (c), respectively. According to the SAED pattern with electron beam paralleling with direction of [2110], the basal plane is marked in Fig. 9(b) where some line-shaped structures along the basal plane can be observed. High resolution transmission electron microscopy (HRTEM) was applied in order to figure out these line-like structures. The HRTEM of rectangular area marked in Fig. 9(b) is magnified as Fig. 9(d) and the magnified image of the square region marked in Fig. 9(d) is given in Fig. 9(e). The inset in Fig. 9(e) offers schematic illustration of stacking faults. The perfect stacking sequence is ABABAB. , and one kind of intrinsic basal-plane faults is formed by removal of a basal layer, which produces a very high energy fault, usually followed by slip of 1/3〈1010〉 of the crystal above this fault to reduce the energy. The dash line in Fig. 9(e) represents the missing half plane which results in stacking fault. It is known that magnesium, with (0001) basal plane as the close-packed plane in hcp structure, has very limited slip systems, and basal slip is the easiest deformation way to be activated. In addition, the stacking faults are the product of dissociation of perfect dislocations in order to lower energy during crystal sliding. Therefore, stacking faults paralleling to basal plane could be observed frequently in magnesium alloys. These line-like structures are testified as stacking faults in the nanograin interior, which were also observed before by Wei et al. [18]. These stacking faults do not pass across the whole grain, but stop in the grain interior with Shockley partial dislocations and these stacking faults

with partial dislocations involved are believed to have a significant effect on deformation behavior especially at the nanometer scale [57]. Fig. 10(a) gives the bright field TEM image of a nano grain with a lot of lattice defects and its HRTEM image is shown in Fig. 10(b). The nano grain with the grain size around 100 nm shown in Fig. 10(a) could not be regarded as a “clean” grain mentioned in previous work [11] of SMAT AZ91D Mg alloy, since diffraction contrast inside the grain is not uniform, which indicates the small misorientations between interior parts exist in this grain. Meanwhile, HRTEM image shows lattice distortions inside the grain, which also demonstrates that the orientation is not uniform in the grain. The magnified HRTEM image of the boundary marked in Fig. 10(a) is shown in Fig. 10(c) where the square area is further magnified as Fig. 10(d). In Fig. 10(d) there is an edge dislocation array acting as a tilt boundary, in which a 5.81 of lattice rotation is obtained as a result. According to the measurement of interplanar distance, the crystal plane observed is confirmed parallel with (1100), from which the direction of 〈1100〉 can be obtained since it is perpendicular to the plane. Therefore, the burgers vector of the dislocation array is parallel with 〈1100〉 or contains a component vector of 〈1100〉, and the former demonstrates that the dislocation array might be a non-basal slip system, such as {1122}/〈1100〉 [58]. The extremely high strain rate at room temperature like what is carried out in the present work may affect the operation of non-basal slip system [59,60]. It is worthy to mention that no matter what kind of dislocation slip systems, a high density of dislocation arrays could be found in this specimen.

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Lattice distortion

0.278 nm

Matrix

(1100)

< 1100 > 5.8

o

Grain

Fig. 10. (a) Bright field TEM image of a nano grain; (b) HRTEM image of the nano grain; (c) HRTEM image of the boundary marked in (a); (d) magnified image of the square area marked in (c).

4. Discussions 4.1. The dynamic rotation recrystallization process during the SMAT process Recrystallization concurrent with deformation is believed to occur by two different mechanisms: rotation recrystallization and migration recrystallization. The former is identified by a gradual increase in the misorientation of a stable dynamic recovery dislocation sub-cell microstructure and eventually the misorientation is such that the sub-boundaries are indistinguishable from high angle grain boundaries. The latter is characterized by a migration of pre-existing high angle grain boundaries through the deforming microstructure leaving a strain free region in their wake [61]. However, as shown in Figs. 9 and 10, there are a lot of defects instead of strain-free region inside the nano grain or subgrain. And the lattice rotation near the boundary is related to the dislocation arrays, as shown in Fig. 10. The existence of dislocation arrays including basal dislocations, prismatic dislocations and pyramidal dislocations will lead to the formation of the low angle boundaries and the lattice rotation when dislocation arrays inside the grain slip to the boundaries [62,63]. It could be expected that dislocation arrays at different planes will result in different rotating directions of the lattice in subcells. In addition, dislocations accumulate or annihilate in the boundaries by gliding through the grain interior in order to accommodate the plastic strain, which is a process of internal energy decrease. Therefore, the misorientation between two subcells gradually increases through

dislocation slips and eventually the low angle boundary becomes into a high angle boundary. The SMAT processes in this paper were conducted at ambient temperature and no obvious temperature increment was tested after treatments though the heat might generate in some localized areas. Dynamic recrystallization (DRX) is usually recognized to happen within temperature ranges from 0.5–0.6 to 0.9–0.98 Tm [64,65] and the recrystallization temperatures is reduced under severe plastic deformation conditions [66]. Furthermore, Kaibyshev et al. [67] claimed that DRX occurred at room temperature in pure Mg when the strain is quite high. During SMAT process, the repeated multidirectional impact of flying steel balls generates severe plastic deformation and the strain rate at the topmost surface could reach up to 104 s  1 which is super high. As a result, a high density of defects inside the substructures appears and nanograins are formed after rotation recrystallization process at room temperature. Therefore, dynamic rotation recrystallization plays a significant role in the formation of nanograins. 4.2. Mechanism of surface nanocrystallization of Mg–3Gd alloy during the SMAT process Based on the above observations, the process and mechanism of surface nanocrystallization of Mg–3Gd with the average grain size of 43 μm could be summarized as a schematic diagram in Fig. 11. At the beginning, there are a great number of dislocations and a few twin lamellae appearing inside the coarse grains due to

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153

Fig. 11. Schematic illustration of the grain refinement process of Mg–3Gd alloy during SMAT.

the very low strain deformation. Along with the increase of the strain and strain rate, the high density of twin lamellae is responsible for the grain refinement since twin boundaries divide the grains into finer twin platelets. The high density of twins is attributed to the higher strain rate and the lower treatment temperature comparing to SPD methods because the twins are easily formed under high strain rate but are difficult to grow at low temperature. As the SMAT process continues, dislocation slips are blocked by twin lamellae and accumulated inside the twin platelets. Meanwhile, some twin lamellae interact with other twin lamellae because of the higher density of twins comparing to the last stage. Therefore, subcells are formed by dislocation arrays intersecting lamellar blocks and twin–twin interactions. It is difficult to form the high-order twins in twin fragments when the twin fragment sizes are dropped to a certain value due to the high critical shear stress, thus the dislocation slips becomes the main plastic deformation mode [68]. Under a higher strain and strain rate, nano-scaled subgrains with a lot of defects and low angle boundaries start to rotate owing to dislocation slips through the grain interior and lattice rotation. Different dislocation slip systems will cause different rotation direction. To accommodate the higher strain, there would inspire the gradual rotation of substructures and increase of misorientation between subgrains, which eventually results in the high angle boundaries. Finally, the nano-sized grains about 50 nm were formed in that the high density of twins and the dynamic rotation recrystallization at ambient temperature. The mechanism of grain refinement of SMATed Mg–3Gd alloy is similar to the previous work of Cu–Zn alloy [45]. However, a high density of nano twins, stacking faults and even secondary twins and stacking faults could be easily achieved in Cu–30 wt% Zn alloy due to the very low stacking fault energy (  7 mJ/m2 [69,70]). And these stacking faults and twin boundaries play a key role in forming subgrains in grain refinement progress. In SMATed Mg–3Gd alloy with the hcp structure and a slightly higher stacking fault energy (  60–78 mJ/m2 [11]), twins are also prevailing because of the limited slip systems. Instead of the high density of secondary twins and stacking faults, twin–twin interactions are more easily observed in SMATed Mg– 3Gd alloy during the grain refinement process.

5. Conclusions (1) Nanometer-sized grains (around 50–100 nm) can be generated in the topmost surface of Mg–3Gd alloy by SMAT process. After treatment for only 4 min, the thickness of nanostructured layer is about 30 μm. (2) The gradient structure of Mg–3Gd alloy can be obtained via SMAT, in which the variation of micro-hardness along the cross section shows a “U” type characterization. The enhancement of the strengths of SMAT Mg–3Gd alloys could be attributed to the grain refinement in the surface layer, the high density of twins in the center region and the strain hardening. (3) The nanocrystallization process of Mg–3Gd by SMAT can be clarified as following three stages: the increase of dislocations followed by the high density of twins; dislocations intersecting lamellar blocks and twin–twin interactions to form subcells; substructures with low angle boundaries rotating through dislocation arrays slipping to generate high angle boundaries, which is classified as dynamic rotation recrystallization.

Acknowledgements The authors gratefully acknowledge the financial supports of the National Natural Science Foundation of China (Nos. 51171113, 51301107, 51474149 and 51464034), the National Key Basic Research Program of the Chinese Ministry of Science and Technology (2012CB932203), the Hong Kong Collaborative Research Fund (CRF) Scheme (C4028–14G) and the Croucher Foundation (No. 9500006). Dr. Long Zeng from Shanghai Jiao Tong University is especially thanked for the support on the casting and extrusion experiments. References [1] R.Z. Valiev, R.K. Islamgaliev, I.V. Alexandrov, Prog. Mater. Sci. 45 (2000) 103–189. [2] T.H. Fang, W.L. Li, N.R. Tao, K. Lu, Science 331 (2011) 1587–1590.

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