Ultrasonic spot welded AZ31 magnesium alloy: Microstructure, texture, and lap shear strength

Materials Science & Engineering A 569 (2013) 78–85

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Ultrasonic spot welded AZ31 magnesium alloy: Microstructure, texture, and lap shear strength V.K. Patel, S.D. Bhole n, D.L. Chen Department of Mechanical and Industrial Engineering, Ryerson University, 350 Victoria Street, Toronto, Ontario Canada M5B 2K3

a r t i c l e i n f o

a b s t r a c t

Article history: Received 6 December 2012 Accepted 17 January 2013 Available online 24 January 2013

The structural application of ultra-lightweight magnesium alloys inevitably involves welding and joining. A solid-state welding – ultrasonic spot welding (USW) – was conducted on AZ31B-H24 Mg alloy sheet by varying a key parameter of welding energy in this study, aiming to identify the changes in the microstructure, crystallographic texture, and lap shear tensile strength. The grain size was observed to increase with increasing welding energy. Crystallographic texture determined via X-ray diffraction showed a significant change that corresponded well to the change in the deformation and recrystallization mechanisms of the Mg alloy. The lap shear strength first increased with increasing energy input, reached the maximum value at a welding energy of 2000 J, then decreased. The welds generally fractured along the joint interface when the welding energy was lower than 2000 J, while fracture occurred at the periphery of the joint (button pull-out) for the samples made with an energy input of higher than 2000 J. & 2013 Elsevier B.V. All rights reserved.

Keywords: Magnesium alloys X-ray diffraction Welding Texture Microstructure Recrystallization

1. Introduction Recently, there has been an increasing demand for lightweight and high specific strength materials in the automotive industry because of resource and environmental concerns [1–3]. Magnesium (Mg) alloys have been receiving a considerable interest in this regard, since Mg is approximately 75% and 34% lighter than steel and aluminum, respectively [4,5]. When Mg alloys are used for the construction of a structure, welding and joining procedures are required. Conventional fusion welding for Mg alloys often produces coarse grains and porosities in the weld metal, which deteriorate the mechanical properties of the joint. A number of papers have recently reported on the joining of Mg alloys by friction stir welding (FSW), friction stir spot welding (FSSW) and resistance spot welding (RSW)[6–8]. However, very little work has been reported for joining Mg alloys using ultrasonic spot welding (USW), which is a solidstate joining process that produces coalescence through the simultaneous application of localized high-frequency vibratory energy and moderate clamping forces. Many physical, chemical and mechanical properties of crystals depend on their crystalline orientations, and it follows that directionality or anisotropy of these properties will result wherever a texture exists in polycrystalline materials [9]. Since preferred orientations of the grains are very common phenomena, crystalline texture sometimes plays an important role in numerous mechanical

n

Corresponding author. Tel.: þ1416 979 5000x7215; fax: þ1416 979 5265. E-mail address: [email protected] (S.D. Bhole).

0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.01.042

and physical behaviors. Crystallographic textures in the hexagonal metals have attracted significant interest over the years because of the use of some hexagonal metals and alloys [10]. It is noted that, for hcp crystalline structure, such as Mg alloys, deformation is dominated by the active deformation modes of slip or twinning, and, hence, the development of texture [11]. The large difference in the critical resolved shear stress between the basal and non-basal slip results in significant anisotropic mechanical properties of Mg alloys [12], and hence the textures that develop during USW would have a strong influence on the mechanical properties. Except for the work done by Zhu et al. [13,14]on the texture evolution of the USWed aluminum foil, most research in the area of USW have worked on the microstructure and the tensile properties of the USWed joints. In addition, a number of studies have reported on the influence of friction stir processing (FSP), FSW, RSW etc. on crystallographic texture and tensile behavior at different positions within the work piece, namely the stir zone (SZ), the thermo-mechanically affected zone and the heataffected zone (HAZ) [15–19]. However, studies so far have not led to a fundamental understanding of the deformation mechanism and dynamic recrystallization during USW and their influence on crystallographic texture formation and mechanical properties of USW Mg alloys. The purpose of this study was, therefore to evaluate the effect of welding parameters on the microstructure, crystallographic texture and lap shear tensile strength of USWed AZ31B-H24 Mg alloy sheets. The macro texture of the USWed AZ31B-H24 Mg alloy sheets was studied by pole figure measurements, and the crystallographic orientations of the basal plane, prismatic plane, and pyramidal plane are discussed.

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2. Materials and experimental procedure In the present study, commercial 2 mm thick sheet of AZ31B-H24 (3%Al, 1%Zn, 0.6%Mn, 0.005%Ni, 0.005%Fe and balance Mg) Mg alloy provided by General Motors Company was selected for the USW. The specimens of 80 mm long and 15 mm wide were sheared, with the faying surfaces ground using 120 emery papers in parallel to the rolling direction (RD), and then washed using acetone followed by ethanol and dried before welding. The welding was done with a dual wedge-reed 20 kHz SonobondMH2016 HP-USW system. The samples were welded at energy inputs ranging from 500 to 3000 J at a constant power setting of 2000 W, an impedance setting of 8 on the machine, and pressure of 0.414 MPa. Four samples were welded in each welding condition. One of them was used for microstructural examination, the other was used for observing crystallographic texture, and the other two were used for the lap shear tensile tests. Cross-sectional samples for scanning electron microscopy (SEM) were polished RD, XRD X- axis

NZ HAZ

BM

TD, Vibration direction

Mg

Mg

79

using diamond paste and MasterPrep. Peak intensities from (1 0 1 0), (0 0 0 2), (1 0 1 1), (1 0 1 2), (1 0 2 0), (1 0 1 3), (0 0 0 4) planes of NZ, at 5 mm distance from the center of the NZ (potential HAZ), BM, and fracture surface (FS) (Fig. 1) of USWed samples were obtained by X-ray diffraction (XRD), using CuKa radiation at 45 kV and 40 mA. The diffraction angle (2y) varied from 301 to 901 with a step size of 0.051 and 2 s in each step. Crystallographic texture distribution of USWed samples was measured by a Panalytical XRD via micro-diffraction with an incident beam with a 300 mm diameter using CuKa radiation (wavelength l ¼0.15406 nm) at 45 kV and 40 mA, and a sample tilt C angle ranging from 0 to 701with a step size of 51 and 3 s in each step. The pole figures were created and analyzed by ‘X-pert texture’ software. Coordinate axes for the pole figures were the rolling direction (RD), transverse direction (TD), and normal direction (ND) of the plate. To evaluate the mechanical strength of the joints and establish the optimum welding conditions, tensile shear tests of the welds were conducted to measure the lap-shear failure load using a fully computerized United testing machine with a constant crosshead speed of 1 mm.min  1 in air at room temperature. In the tensile lap shear testing, restraining shims or spacers were used to minimize the rotation of the weld joints and maintain the shear loading for as long as possible.

15 mm

3. Results and discussion 80 mm

3.1. Microstructure

25 mm

5 mm

2 mm 20 mm

FS Fig. 1. Schematic diagram of the lap joint.

Microstructure characterization was focused on the NZ at the center of the weld samples. Fig. 2 depicts the microstructure of the AZ31B-H24 base metal (BM) and 1000 J, 2000 J, and 3000 J welded samples. The recrystallized grain morphology of the BM, although not typical of the H24 condition, has also been reported elsewhere [20]. The microstructure in the welded samples was fully recrystallized. The mean grain size across all welded samples

40 μm

40 μm

40 μm

40 μm

Fig. 2. Optical micrograph of AZ31 BM and welded samples at different energy levels. (a) AZ31 BM, (b) 1000 J, (c) 2000 J, (d) 3000 J.

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(1011)

20

Intensity, a. u (103)

12

FS

8

NZ HAZ

4

(2022)

(1014)

(1122) (2021) (0004)

(2020)

(1013)

(1120)

(1012)

(1010) (0002)

16

BM 0 30

The tensile properties of the weld of Mg alloy also depend on several microstructural factors, such as grain size and dislocation density [16]. However, besides grain size and dislocation density, crystallographic orientation also strongly influences plastic deformation during the tensile test because the plastic deformation arises from slip on the close-packed planes with the maximum critical resolved shear stresses [16]. Initially, effort has been taken to analyze the texture behavior by XRD on the top surface (at different locations such as BM, HAZ, and NZ of the USWed joint made at 1000 and 2000 J energy inputs, along with the fracture surface (FS) of the USWed sheet (Fig. 1)). HAZ and BM analysis points were located at 5 and 30 mm from the center of the NZ (Fig. 1), while the FS point was located at the center of the NZ in-between the coupons. Fig. 3 shows the computer simulated X-ray diffraction for completely random Mg powder. The (1 0 1 0) and (0 0 0 2) peak heights are about one half of the (1 0 1 1) peak height, all being present within the 2y angles of 30–401. The XRD diffraction patterns of the as-received BM, HAZ, NZ and FS at 1000 and 2000 J energy inputs are shown in Fig. 4(a) and (b). In the BM, intensity of (0 0 0 2) basal plane is higher than any of other planes, which is in complete contrast to the Mg powder data. This suggests that the BM itself exhibits a typical basal rolling texture, owing to the mechanical deformation during the cold rolling of Mg sheet. It can be seen from Fig. 4(a) that in NZ and FS (1000 J), the intensity of (0 0 0 2) plane was approximately two times higher than that in the BM, suggesting, basal texture became stronger after the welding at 1000 J energy input. Similarly, at 2000 J energy input (Fig. 4(b)) on the FS, the intensity of (0 0 0 2) plane was approximately two times higher than that in the BM, which could be the cause of lap shear tensile deformation However, in NZ, the intensity of (1 0 1 1) plane was higher than the intensity of (0 0 0 2) plane (similar to Mg powder data), which indicates a nearly random grain orientation, reflecting its fully recrystallized coarse grain structure as shown in Fig. 2. Similar

60

70

80

90

20

16

Intensity, a. u (103)

3.2. Crystallographic texture

50

2θ, degree

Fig. 3. Simulated X-ray diffraction pattern for the completely random Mg using Cu-Ka radiation. The vertical axis is for the X-ray intensity. Within the 2y angles of 30–401, the (1 0 1 0) and (0 0 0 2) peak heights are about one half of the (1 0 1 1) peak heights.

was coarser than the mean grain size of the as-received samples (5.36 mm) and it can be seen that the grain size increases with welding energy from 7.8 to 10.1 mm owing to an increase in temperature. In addition, as energy inputs increase from 1000 to 3000 J, the aspect ratio of the grains also increases slightly from 1.55 to 1.63.

40

12

FS

8

NZ HAZ

4

BM 0 30

40

50

60

70

80

90

2θ, degree Fig. 4. X-ray diffraction patterns obtained from BM, HAZ, NZ and FS at (a) 1000 (b) 2000 J USWed AZ31B-H24 Mg alloy.

result was also reported by Chang et al. in the FSP of the Mg alloy with increasing rotational speed from 800 to 1800 rpm [21]. It can be noticed that HAZ does not exhibit the texture changing, hence, it has similar texture profile as BM which is in agreement with some literatures reported that USW does not produce HAZ [22,23]. For more detailed analysis of the above mentioned results, pole figures have been plotted for all region of the welded joint (NZ, HAZ, BM and FS (Fig. 1)) and shown in Fig. 5(a) and (b). The (0 0 0 2) basal, (1 0 1 0) prismatic, and (1 0 1 1) pyramidal plane pole figures obtained from the top surface of the USWed joint made at 1000 and 2000 J energy inputs. It can be seen from Fig. 5(a) and (b) that the BM region in the USWed weld coupon of 1000 and 2000 J samples has a typical strong basal texture, where basal plane (0 0 0 2) normal was largely parallel to the ND with some grains slightly tilted in the RD (Fig. 5(a) and (b)), which is in agreement with the XRD results. It has been reported that two major types of texture components, (0 0 0 2)/1 1 2 0S and (0 0 0 2)/1 0 1 0S, are available, depending on the activation of slip systems in the basal plane (either /1 1 2 0S single slip or /1 0 1 0S double slip oriented in the RD) in the Mg alloys [24,25]. The combination of the two components generates a pole figure of rolled Mg (BM of USW) that is almost a (0 0 0 1) fiber texture. This (0 0 0 1) fiber texture can be easily seen by prismatic

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RD

TD

Base Metal (BM)

Basal plane (0002)

Prismat ic plane (10 1 0)

Pyrami dal plane

Heat Affected Zone (HAZ)

Nugget Zone (NZ)

Fracture Surface (FS)

1 26 51 76 100

Intensity Color 0.147 1.898 3.650 5.401 7.082

1 26 51 76 100

Intensity Color 0.095 2.473 4.850 7.228 9.510

1 26 51 76 100

Intensity Color 0.235 6.117 11.999 17.880 23.527

1 26 51 76 100

Intensity Color 0.285 4.755 9.225 13.694 17.985

1 26 51 76 100

Intensity Color 0.345 0.787 1.229 1.671 2.095

1 26 51 76 100

Intensity Color 0.275 0.778 1.280 1.783 2.265

1 26 51 76 100

Intensity Color 0.346 1.304 2.261 3.219 4.139

1 26 51 76 100

Intensity Color 0.341 0.925 1.510 2.095 2.656

1 26 51 76 100

Intensity Color 0.230 0.681 1.133 1.584 2.017

1 26 51 76 100

Intensity Color 0.100 0.616 1.132 1.649 2.144

1 26 51 76 100

Intensity Color 0.255 0.697 1.139 1.581 2.005

1 26 51 76 100

Intensity Color 0.405 0.736 1.067 1.398 1.715

(10 1 1)

Fig. 5. (0 0 0 2) basal plane (1 0 1 0) prismatic plane and (1 0 1 1) pyramidal plane pole figures of USWed AZ31B-H24 Mg alloy obtained from BM, HAZ, NZ and FS at (a) 1000 J (b) 2000 J energy inputs.

(1 0 1 0) plane (circular ring type structure at 901 from the center of the pole figure). Moreover, the crystallographic fibering in the BM could be even better seen from the pyramidal planes (1 0 1 1), where the pyramidal planes (1 0 1 1) formed a nearly complete ring at approximately 451 from the center of the pole figure. Similar (0001) fiber texture component in the rolled AZ31 Mg alloy has also been reported in [26,27]. From the (0 0 0 2) basal (1 0 1 0) prismatic and pyramidal (1 0 1 1) planes of the HAZ region, similar BM (0 0 0 1) fibering characteristics have been observed. It should be noted that both BM and HAZ crystallographic texture results obtained from the pole figures are in

agreement with the peak intensities observed from the XRD analysis (Fig. 4). In NZ of 1000 J energy input, a strong basal texture with (0 0 0 2) normal parallel to the ND was found (Fig. 5(a), NZ-(0 0 0 2) plane). The intensity of (0 0 0 2) plane was much higher than that of the BM, suggesting that after the welding more grains followed the basal texture. Such a texture change occurring in the NZ was attributed to the intense localized shear plastic flow during USW. It can be seen from the (1 0 1 0) prismatic plane pole figure of the NZ for 1000 J USWed sample (Fig. 5(a), NZ-(1 0 1 0) plane) that it breaks the fiber texture (which was in BM) and the prismatic

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Base Metal (BM)

Basal plane (0002)

Prismat ic plane (10 1 0)

Pyrami dal plane

Heat Affected Zone (HAZ)

Nugget Zone (NZ)

Fracture Surface (FS)

1 26 51 76 100

Intensity Color 0.230 2.202 4.173 6.144 8.037

1 26 51 76 100

Intensity Color 0.203 2.169 4.134 6.099 7.986

1 26 51 76 100

Intensity Color 0.073 1.325 2.577 3.829 5.030

1 26 51 76 100

Intensity Color 0.235 5.836 11.438 17.040 22.418

1 26 51 76 100

Intensity Color 0.420 0.833 1.247 1.660 2.057

1 26 51 76 100

Intensity Color 0.453 0.833 1.213 1.592 1.957

1 26 51 76 100

Intensity Color 0.219 0.689 1.159 1.629 2.081

1 26 51 76 100

Intensity Color 0.207 1.058 1.910 2.761 3.578

1 26 51 76 100

Intensity Color 0.417 0.757 1.096 1.436 1.762

1 26 51 76 100

Intensity Color 0.339 0.741 1.143 1.545 1.931

1 26 51 76 100

Intensity Color 0.202 0.733 1.264 1.795 2.304

1 26 51 76 100

Intensity Color 0.331 0.789 1.248 1.706 2.146

(10 1 1)

Fig. 5 b. Continued.

plane with /1 0 1 0S becomes parallel to the RD of the sheet surface. Interestingly, after the lap shear tensile test, on the fracture surface, all the grains take their original orientation (as in the BM) and exhibit a fiber basal texture with higher intensity than that in the BM (Fig. 5(a), FS-(1 0 1 0) and (1 0 1 1) planes). In NZ of 2000 J (Fig. 5(b), NZ-(0 0 0 2) plane), intensity became very low and creates 901rotation of the basal plane so that its normal is parallel to the TD at an angle of 601 and slightly tilted at an angle of 5 –101 towards the RD. Since the crystallographic texture changed from BM to 1000 to 2000 J USWed samples, it could be noted that the texture at NZ is affected by welding heat input. From our previous study, temperature increases from 357 to 488 1C with increasing energy input from 1000 to 2000 J, respectively [28]. In 2000 J energy input sample, with high temperature and slow cooling rate, more complete dynamic recrystallization occurred. Thus, some new grains may nucleate and grow as shown in Fig. 2 with the result that the texture becomes weak. Dynamic

recrystallization is produced by the large strain introduced by the ultrasonic vibration and high temperatures developed from friction and plastic deformation during USW. Moreover, dynamic recrystallization promotes stress concentration at grain boundaries [29], which is effective for the activation of non-basal slip system and serves to break down the basal texture. In addition, Jain and Agnew [30] stated that critical resolved shear stresses (CRSS) of non-basal slip systems decline dramatically at elevated temperatures, and non-basal slip systems are more easily activated and tend to split the basal texture. Eventually, in the present study, the basal texture at the NZ becomes weak. Similar results were also found by Chang et al. in FSP [21] and Yu et al. in fiber laser welding [31] of Mg alloys. It is suggested here that in USW the plastic deformation due to the ultrasonic vibrations leads to texture formation while the increase in temperature due to the USW heat generation leads to dynamic recrystallization and randomization of texture; these two phenomena are in competition with one another. Thus, in 1000 J

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lap shear strength was found to be 89 MPa at 2000 J energy input (calculated as the maximum lap shear fracture load divided by the nugget area of 8  6 mm). As shown in Fig. 6, with increasing energy input, lap shear strength increased in the beginning owing to high temperatures and strain rate, which accelerate diffusion between the Mg alloy sheets, but at very high energy inputs, it was decreased. This may be because, at low welding energies (1000 J), Mg has not been sufficiently deformed plastically to fill the valleys of the knurl patterns of the sonotrodes, as shown in Fig. 7(a). In contrast, at 2000 J energy input (Fig. 7(b)), the knurl pattern valleys fill up with the plastically deformed Mg alloy. In addition, Yu et al. [32], correlated Zener–Hollomon parameter (Z) (widely used to show the combined effect of temperature and strain rate [32–34]) with the tensile properties of friction stir processed Mg alloy, indicating that samples with high Z value show low strength/high ductility and samples with low Z value show high strength/low ductility. Similar relationship has also been established in the present study, where the values of Z for 1000 J and 2000 J were 5.79  1013 and 1.38  1012 [28], with the strength being 67 and 89 MPa, respectively. It can be seen from Fig. 7(c) and (d), at 2000 J energy input, more severe bending takes place, while for the weld produced by lower energy inputs (1000 J), the separated sheets remain almost flat. Interestingly, at the highest energy input of 2500 J, weld specimens experienced large compressive stress developed at the edge of the weld spot (Fig. 7(e)), resulting in extensive local plastic deformation and nugget pull-out during the lap shear test experiments.

60

3.4. Fractography

50

In the lap-shear tensile tests, the welds generally fracture along the weld interface for input energies of less than and equal to 2000 J, while fracture occurs at the periphery of the weld (button pullout) for input energies of greater than 2000 J. The typical failure locations for the USWed of AZ31B-H24 Mg alloy are shown in Fig. 8. It can be seen from Fig. 8(a) that in 1000 J energy input sample the whole fracture surface is flat. In contrast, a big crack was found (Fig. 8(b)) at the periphery of the sonotrode tip indent in 2000 J energy input sample due to the large compressive stress as mention earlier. This crack can penetrate into the

USWed sample the plastic deformation dominates giving a sharp texture while in 2000 J USWed sample the dynamic recrystallization dominates giving a weak texture. The pole figure of prismatic plane of NZ in 2000 J sample (Fig. 5(b), NZ-(1 0 1 0) plane) is almost similar to that of the 1000 J sample. However, the prismatic plane is more concentrated along the RD in the 1000 J sample than the 2000 J sample. As a result, its intensity was one half of the prismatic plane of NZ in 1000 J USWed sample. By observing the texture on the fracture surface of 2000 J USWed sample (Fig. 5(b), FS-(0 0 0 2), (1 0 1 0) planes), it seems that the basal plane regained their original orientation (as in the BM) with higher intensity than BM and NZ. However, from (1 0 1 1) pyramidal plane of fracture surface, (0 0 0 2)/1 1 2 0S texture component was mainly observed (Fig. 5(b), FS-(1 0 1 1)). 3.3. Lap shear strength of USWed joints The lap shear strength of the USWed coupons is plotted against welding energy inputs as shown in Fig. 6. The maximum

100

Lap shear strength, MPa

83

90 80

40 30 0

500

1000

1500

2000

2500

3000

Welding energy, J Fig. 6. Lap shear strength of USWed similar AZ31-H24 Mg alloy at different energy inputs.

Fig. 7. Effect of the knurl pattern on the USWed sample at (a) 1000 J (b) 2000 J and bending effect during the failure at (c) 1000 J and (d) 2000 J and nugget pull out effect (e) energy inputs.

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Fig. 8. Typical SEM images of tensile fracture surface of USWed similar AZ31-H24 Mg alloy, (a) and (b) overall view, (c) and (b) at lower magnification, (e) and (f) at higher magnification of 1000 J and 2000 J energy inputs, respectively.

material more deeply with higher energy. In particular, due to this crack, the base materials were partly pulled out from the welded interface. Thus, further increase in energy input (2500 J) resulted in the whole nugget pull out as shown in Fig. 7(e). Similar nugget pull out phenomenon in the USW process was also reported [35]. A magnified region for 1000 J energy input sample (Fig. 8(c) and (e)) shows the failure mode in the center region of the fracture surface. Apart from some small isolated dimple-rupture failure zone, the dominant feature of the region is the parallel scrubbing lines. Owing to high frequency ultrasonic vibration, these scrubbing lines resulted from abrasion or galling of the sheet surfaces. In contrast, a characteristic dimple-rupture failure mode was found at 2000 J energy input as shown in Fig. 8(d) and (f).

4. Conclusions

1. USW of similar AZ31B-H24 Mg alloy sheets was successfully achieved. 2. After USW, the grain size was observed to increase with increasing welding energy. 3. The scatter of the (0 0 0 2)/1 0 1 0S and (0 0 0 2)/1 1 20S generates a pole figure of rolled Mg that is almost a (0 0 0 1)

fiber texture, which was mainly observed in the BM region of the USWed Mg alloys. After USW of Mg alloy at 2000 J, dynamic recrystallization creates 901 rotation of the basal plane in the NZ so that its normal is parallel to the TD in the NZ at the angle of 601 and slightly tilted at the angle of 5–101 towards the RD. This may because the induced higher temperature rise due to the friction results in more complete dynamic recrystallization. 4. With increasing energy input, lap shear strength increased and failure occurred by interface fracture. However, at very high energy inputs, the lap shear strength decreased and failure occurred by weld button pullout. 5. This study shows a good correlation among the processing parameter, texture development, and deformation mechanism for the USWed Mg alloy AZ31B-H24. The results could be useful for the design of desired final properties of USWed Mg alloys.

Acknowledgments The authors would like to thank the Natural Sciences and Engineering Research Council of Canada (NSERC) and AUTO21 Network of Centers of Excellence for providing financial support. This investigation involves part of Canada–China–USA Collaborative

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Research Project on the Magnesium Front End Research and Development (MFERD). The authors thank Dr. A.A. Luo, General Motors Research and Development Center for the supply of test materials. One of the authors (D.L. Chen) is grateful for the financial support by the Premier’s Research Excellence Award (PREA), NSERC-Discovery Accelerator Supplement (DAS) Award, Canada Foundation for Innovation (CFI), and Ryerson Research Chair (RRC) program. The assistance of Q. Li, A. Machin, J. Amankrah, and R. Churaman in performing the experiments is gratefully acknowledged. References [1] E. Aghion, B. Bronfin, D. Eliezer, J. Mater. Process. Technol. 117 (2001) 381–385. [2] Q. Schiermeier, Nature 316 (2011) 470–471. [3] L.M. Liu, H.Y. Wang, Mater. Sci. Eng. A 507 (2009) 22–28. [4] T.M. Pollock, Science 328 (2010) 986. [5] X. Cao, M. Jahazi, J.P. Immarigeon, J. Wallace, J. Mater. Process. Technol. 204 (2006) 171–188. [6] X. Cao, M. Jahazi, Mater. Des. 30 (2009) 2033–2042. [7] W. Yuan, R.S. Mishra, B. Carlson, R. Vermac, R.K. Mishra, Mater. Sci. Eng. A 543 (2012) 200–209. [8] J.C. Feng, Y.R. Wang, Z.D. Zhang, Sci. Technol. Weld. Joining 11-2 (2006) 154–162. [9] Y.N. Wang, J.C. Huang, Mater. Chem. Phys. 81 (2003) 11–26. [10] A. Kouadri, L. Barrallier, Mater. Sci. Eng. A 429 (2006) 11–17. [11] A. Galiyev, R. Kaibyshev, G. Gottstein, Acta Mater. 49 (2001) 1199–1207.

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