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Severe plastic deformation of a TWIP steel
Materials Science & Engineering A 593 (2014) 163–169
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Severe plastic deformation of a TWIP steel I.B. Timokhina a,n, A. Medvedev b, R. Lapovok b a b
Institute for Frontier Materials, Deakin University, Geelong, Victoria 3217, Australia CAHM, Materials Engineering Department, Monash University, Clayton, Victoria 3800, Australia
art ic l e i nf o
a b s t r a c t
Article history: Received 3 October 2013 Received in revised form 4 November 2013 Accepted 7 November 2013 Available online 17 November 2013
The severe plastic deformation of a Twinning Induced Plasticity (TWIP), 0.61C–22.3Mn–0.19Si–0.14Ni– 0.27Cr (wt%) steel by Equal Channel Angular Pressing (ECAP) at elevated temperatures was used to study the deformation mechanism as a function of accumulated strain and processing parameters. The relationship between the microstructures after different deformation schedules of ECAP at the temperatures of 200, 300 and 400 1C, strain hardening behavior and mechanical properties was studied. The best balance between strength and ductility (1702 MPa and 24%) was found after two passes at 400 1C and 300 1C of ECAP. It was due to the formation of deformation microbands and twins in the microstructure. The twinning was observed after all deformation schedules except after one pass at 400 1C. The important finding was the formation of twins in the ultrafine grains. Moreover, the stacking faults were observed in the subgrains with the size of 50 nm. It is also worth mentioning the formation of nano-twins within the micro-twins at the same time. It was found that the deformation schedule affects the dislocation substructure with formation of deformation bands, cells, subgrains, two variants of twins that, in turn, influence the strain-hardening behavior and mechanical properties. & 2013 Published by Elsevier B.V.
Keywords: Twinning Induced Plasticity steels Equal Channel Angular Pressing Mechanical properties Transmission electron microscopy Micro/nano-twins Dislocation substructure
1. Introduction The great interest in Twinning Induced Plasticity (TWIP) steels from the automotive industry is inspired by their high hardening rate during deformation, high strength and good ductility which together provides superior toughness compared to other types of steel [1]. TWIP steels have low stacking fault energy (SFE) between 15 and 40 mJ/m2 at room temperature through the high level of manganese, in the range of 15–30 wt% [2–6]. The high Mn content stabilizes the austenite at room temperature and produces a single phase face of centered cubic (fcc) steel with low stacking fault energy. It has been shown that in the case of a TWIP alloy with manganese content of 23 wt% and carbon content 0.65 wt%, as in the current research, the SFE is 23 mJ/m2. This SFE reduces the ease of cross slip upon deformation allowing twinning to become the favored deformation mechanism [7,8]. On the other hand, for this steel composition, the SFE was not low enough (γfcc r12 mJ/ m2) to promote the martensite transformation and, hence, this was suppressed in the current steel. The deformation behavior of TWIP steels has normally been examined by Barbier in interrupted tension [9] or compression tests [10] with different grain sizes in samples produced by the hot rolling and/or cold rolling followed by annealing [11,12]. In such n
Corresponding author. Tel.: +61 3 9905 9608; fax: +61 3 9905 4940. E-mail address: [email protected] (I.B. Timokhina).
0921-5093/$ - see front matter & 2013 Published by Elsevier B.V. http://dx.doi.org/10.1016/j.msea.2013.11.013
tests the strain that could be achieved was below 100% and grain size ranges varied from 1 to 60 mm and above [10–12]. It is known that the propensity for twinning decreases with a decrease in the grain size [13]. For example, there was no twinning observed in TWIP samples with a grain size of around 2 mm [12] and the strengthening mechanism in this case was attributed only to the dislocation interactions. Although the data on the effect of strain and grain size on the formation of deformation twins obtained in previous investigations is quite clear, the effect of severe plastic deformation, such as e.g. Equal Channel Angular Pressing (ECAP), which introduces a large shear strain above 100% into the material and refines the grains size to the submicron range, on the formation of deformation twins has not been investigated. Since the TWIP steel has high strength and ductility, it is expected that ECAP deformation could only be performed at elevated temperatures. However, an increase in temperature is known to lead to a decrease in the volume fraction of the twins, twin length and twin thickness [14,15]. Two of three Gibbs energy components of SFE are temperature dependent. As has been published in [5,16,17], the change in the temperature leads to changes in the SFE which promotes different deformation mechanisms. For example, an increase in the deformation temperature to 400 1C suppresses twinning and favours dislocation glide [18]. In addition, it also has been reported [19] that there is a critical strain for the onset of twinning, and this increases with the temperature. However, the behavior of ultrafine grained materials
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Fig. 1. Optical (a) and SEM (b–d) images of initial microstructure (a), after one ECAP pass at 400 1C (b), after four ECAP passes at 400 1C (c) and after one pass at 200 1C (d) of TWIP steel samples.
is quite often quite different from what is observed in course grained materials. In this paper for the first time, severe plastic deformation by ECAP has been applied to a TWIP steel. The formation of the ultrafine grained microstructure is discussed in relation to the strain history and temperature of material processing. Different temperatures and number of ECAP passes were used to manipulate the microstructure and mechanical properties.
Table 1 ECAP schedules.
2. Material and experimental procedure
study [5] to promote twinning deformation after one pass of ECAP. The changes in the mechanical properties and microstructures of the TWIP steel as a function of processing routes (Table 1) at elevated temperatures were studied. Sub-sized round tensile samples were machined out of the ECAP processed billets with a gage length of 10 mm. The tensile test was carried out using an Instron 5982 machine with a cross-head velocity of 1 mm/min corresponding to a nominal strain rate of about 0.002 s 1 for all conditions listed in Table 1. Materials after two conditions with the first and second passes at 200 1C were not mechanically tested. This was associated with the drastic increase in strength during deformation that led to only partial ECAP before the punch broke down. The lack of sufficient material volume processed by ECAP at 200 1C made it impossible to machine tensile samples but the microscopic study was been carried out. The grain size of the selected samples was measured using a FEI Quanta 3D FEG FIB scanning electron microscope with EBSD at a 15 kV accelerating voltage. The analysis of the
The composition of TWIP steel utilized in this study was Fe– 0.61C–22.3Mn–0.19Si–0.14Ni–0.27Cr (wt%). The 10 mm plates were supplied by Arcelor Mittal (France), which were processed by hot rolling from continuously cast ingot and annealed according to their industrial schedule. The initial microstructure was an equiaxed shape austenite grains with the size of 34721 mm, Fig. 1a. Cylinders of 10 mm diameter and 35 mm length were cut from the plate in the rolling direction and subjected to the different schedules of ECAP passes as given in Table 1. The isothermal ECAP rig used for processing is described elsewhere [20,21]. The deformation temperature of 400 1C was used to obtain grain refinement without triggering twinning. The transition temperature from twinning to slip deformation mode was chosen based on the data reported by others for a TWIP steel with similar composition [5]. The temperature of 200 oC was selected based on a previous
Number of ECAP passes
1 (1C)
1 (1C)
2 (1C)
2 (1C)
3 (1C)
4 (1C)
Temperature during ECAP passes
400
200
400 200
400 300
400 300 300
400 400 400 400
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dislocation substructure formed after different schedules of the ECAP passes was performed using a Philips CM 20 operation at 200 kV with a condenser aperture of 100 μm nominal diameter and a nominal beam diameter of 55 nm. The foils were cut from the center of the sliced samples parallel to the x and y planes. Each TEM foil was prepared by twin-jet polishing in a solution of 5% perchloric acid and 95% methanol using the Struers Tenupol 5 electropolisher. The operating voltage was 50 V at a temperature of 25 1C.
3. Results and discussion 3.1. Mechanical properties and strain hardening behavior of the samples after ECAP The engineering stress–strain curves in the initial hot rolled conditions showed an exceptional ductility of 123% and quite high strength of 930 7 5 MPa, Fig. 2b. After one ECAP pass at 400 1C the ductility decreased to 44%, while the strength increased to 1286 710 MPa, Fig. 2a. The second pass at 300 1C increased the strength further to 170278 MPa but decreased the ductility to 24%. However, the third pass at 300 1C had an opposite effect on strength, i.e. the strength dropped to 16247 10 MPa, while ductility also decreased to 21%, Fig. 2a. It should be noted that the samples between the passes were heated for about 5 min and some recovery might take place. Due to this recovery the strength after four passes of ECAP at 400 1C decreased to 15737 6 MPa, while the ductility was similar (25%) compared to those samples that were deformed at the temperature of 300 1C in the following ECAP passes, Fig. 2a. The properties of the TWIP steel in as-received and deformed conditions are summarized in Table 2 It can be seen that yield stress after processing by ECAP increases significantly from 312 MPa to 980–1480 MPa. The strain hardening rate normalized by shear modulus vs. difference between flow and yield stresses (true stresses) for all deformed states and the as-received condition is shown in Fig. 3. As can be seen in Fig. 3a, the character of strain hardening rate of the TWIP steel in the as-received condition, is qualitatively similar to that described by others [13] i.e. firstly, the curve showed a high overall strain hardening rate; secondly, the curve showed a local minima at intermediate strains and an increase in the strain hardening rate at higher deformation [13]. In particular, the initial value of the strain hardening rate equal to G/40, decreases to a minimum of G/50 when ðs sy Þ reaches 200 MPa and then starts to rise steadily to the value of G/28 and keeps constant for ðs sy Þ between 1200 and 1600 MPa. At ðs sy Þ higher than 1600 MPa the strain hardening rate drops rapidly. This curve behavior appeared to be associated with both dislocation accumulation and twin substructure formation [13], Fig. 3a. The strain hardening rate behavior for the samples processed by different ECAP schedules was different, Fig. 3b. The strain rate
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curves after one pass and four passes at 400 1C showed similar behaviors i.e. after the monotonic drop, the strain hardening rate stabilizes and remains constant for ðs sy Þ in the range of 200– 700 MPa and 400–600 MPa, respectively, Fig. 3b. However, the overall strain hardening rate was lower after one pass at 400 1C. The strain hardening rate curve after one pass at 400 1C followed by one pass at 300 1C initially revealed a steady strain hardening rate of G/5.5 up to 200 MPa with a continuous decrease at higher values of ðs sy Þ. The strain hardening rate curve after one pass at 400 1C followed by two passes at 300 1C demonstrated a rapid decrease from the initial value ðG=6⋯G=5Þ, Fig. 3b. 3.2. Microstructural characterization of the samples after ECAP The SEM of the samples after one pass at 400 1C showed the formation of elongated austenite grains with a thickness from 6 to 11 μm (Fig. 1b). The TEM of the samples from the similar deformation condition revealed the formation of parallel microbands separated by a distance of 260 7 37 nm, indicated by white arrows in Fig. 4a, as a representative microstructural feature. The microbands were thick and parallel within 21 to the {111}γ plane. A SAD pattern with [110]γ zone axis, obtained from the microbands, did not show extra diffraction sots that could be attributed to the twinning. Similar microbands were observed by others in the TWIP steels after high strain rate experiments [22,23] and were described as “first generation microbands” [22,23]. Dislocations with a specific cell configuration were observed inside the microbands, indicated by black arrows in Fig. 4b. As reported previously [24], these dislocations cells can play an important role in the strain hardening behavior due to the forest hardening associated with short range dislocation interactions [24]. This hardening mechanism results in a high critical stress when transferring plastic deformation across the microbands [24]. SEM of the microstructure after four passes at 400 1C showed the essential refinement of the austenite grains with grain size variation from 0.3 to 0.6 μm, Fig. 1c. The analysis of the microstructure using TEM after four passes at 400 1C revealed that the Table 2 Yield stress, strength and ductility (engineering) of TWIP steel at different conditions. ECAP schedules
sy (MPa)
su (MPa)
ε (%)
As-received 1 Pass (400 1C) 1 Pass (200 1C) 2 Passes (400 1C þ 200 1C) 2 Passes (400 1C þ 300 1C) 3 Passes (400 1C þ 3001C þ 3001C) 4 Passes (400 1C þ 400 1Cþ 400 1C þ400 1C)
312 980 – – 1250 1480 1290
930 1286 – – 1702 1624 1573
123 44 – – 24 21 25
Fig. 2. Engineering stress vs. engineering strain tensile curves for the samples processed by ECAP (a); sample in as-received condition (b).
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Fig. 3. Strain hardening rate (normalized by shear modulus) vs. difference between true and yield stresses for sample in as-received condition (a); samples processed by ECAP (b).
Fig. 4. TEM micrographs of the microstructure formed after one pass (a) and four passes (b–h) at 400 1C: (a) formation of microbands, [110]γ zone axis, white arrows indicate the microband walls and black arrows indicate the dislocations inside microbands, (b) formation of subgrains with [110]γ zone axis and free of dislocations, (c) and (d) subgrains free and containing parallel sessile dislocations, (e) and (f) formation of stacking faults in the subgrain, (g) bright and (h) dark field images of the twins formed in the microstructure.
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representative microstructural feature was the formation of subgrains with an average size of 507 10 nm within the austenite grains. The subgrains showed the low angle of grain boundary misorientation between them and high angle of grain boundary misorientation with respect to the matrix, Fig. 4b and c. The subgrains could roughly be divided into three groups: (i) free of dislocations, Fig. 4c; (ii) containing parallel sessile dislocations, which should be the result of high strain accommodation in the lattice, Fig. 4d; and (iii) containing stacking faults, Fig. 4e and f. It is believed that the microstructural recovery, which occurred at the temperature of 400 1C between each pass, should be taken into account in understanding the mechanism of the subgrains formation. The microbands, which were the representative dislocation substructure after one pass at 400 1C, were not observed in the microstructure of the samples after four passes at 400 1C. However, the microband walls could be the preferential nucleation sites for subgrain formation. The formation of the stacking faults, Fig. 4f (indicated by arrow), in 50 nm subgrains was not expected and appeared to be an intriguing feature. Despite the general understanding that twinning is absent at high temperatures such as 400 1C, the formation of fine twins crossing the austenite boundary was observed in the microstructure, Fig. 4g and e. The twins were surrounded by an area with high dislocation density, which could be associated with the stress propagation during the twin formation. The strain hardening rate curve behavior after four passes was similar to that after one pass at 400, although the average strain hardening rate was higher after four passes than after one pass. Moreover, a significant increase in yield and tensile strength after four passes (Table 2) was observed compared to the samples after one pass at 400 1C due to the formation of the more complex dislocation substructure with the formation of subgrains and thin twins. It should be noted that the formation of twins after four passes of ECAP at 400 1C could not be explained from a SFE point of view as this is proposed to be in the range of 12–45 mJ/m2 to favor the twinning mechanism. The calculated SFE for a similar TWIP steel (Fe–22 wt% Mn–0.6 wt% C at 300–400 1C) at similar temperature, carried out in [1], was much higher than the SFE for twinning. Twinning is normally not expected for coarse grained material for these temperatures. Moreover, a decrease in the grain size normally suppresses twinning [16] and, hence, the major reduction in the grain size after four passes should make twinning more difficult [25,26]. Therefore, both the temperature and grain size effects should lead to the suppression of twinning after four passes of ECAP at 400 1C. For coarse grained materials, both deformation mechanisms, dislocation slip and twinning, follow the Hall–Petch
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relationship with slope of the curve steeper for twinning [27], that makes twinning more difficult with grain size decreasing. However, as shown in several publications by Zhu et. al. [28,29] there is a transition of ultrafine grain size (UFG), which promotes the formation of twins in TWIP steels. Moreover, the formation of deformation twins was not only experimentally observed but also modeled in UFG materials, with several twinning mechanisms proposed [29]. Deformation twinning has been shown to be a major plastic deformation mechanism for materials with a grain size below 100 nm [28]. This is consistent with microstructural observations in the current research, when the several ECAP passes were performed at 400 1C. The representative microstructural features in the samples after one pass at 400 1C followed by one pass at 300 1C were microband formation along with the formation of twins, Fig. 5. The microbands had an average thickness of 20 75 nm and showed low grain boundary misorientation, Fig. 5b. The microbands were formed at an angle of 6575o with the respect to the matrix. It appeared that the microbands nucleated at the grain boundary and then propagated towards the grain interior, Fig. 5b. Straight dislocations were observed within the microbands, Fig. 5b. The diffraction pattern with the [110]γ zone axis did not show extra diffraction spots and the dislocation morphology did not change during sample tilting. This suggests that these were sessile dislocations that might have formed by the stair-rod cross slip mechanism [30]. The stair-rod dislocations are normally considered as the twinning precursors [30]. The micro-twins with an average spacing of 25 75 nm were also observed in the microstructure, Fig. 5a. The mechanical properties of the samples after one pass at 400 1C followed by one pass at 300 1C were similar to those after four passes at 400 1C, although the strain-hardening rate behavior was different. This was associated with the difference in the dislocation substructure formed after the two different deformation schedules. The addition of the third pass at 300 1C resulted in the formation of grains with high dislocation density, as can be seen in grain 1, Fig. 6a, and grains with twins, as observed in grain 2, Fig. 6a. Two types of micro-twin: (i) fine twins, Fig. 6b, c, and (ii) coarse twins, Fig. 6d, e, were detected The fine twins with an average distance between the twins of 207 5 nm were observed in the austenite grains with high dislocation density, Fig. 6b. However, in most of the austenite grains, two variants of fine deformation twins perpendicular to each other were observed: (i) the primary twins shown by the arrows in Fig. 6c with a spacing of 100 nm, and (ii) the secondary twin variant with a spacing of 50 nm, Fig. 6c. The formation of two variants of deformation
Fig. 5. TEM micrographs and diffraction patterns with [110]γ zone axis of twins formed in grains 2 and 3 (a), and microbands observed in grain 1 and microbands enlarge view (b) with zone axis [110]γ after two passes at 400 1C and 300 1C. Dot line represents the grain boundary.
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Fig. 6. TEM micrographs of the microstructure formed after passes at 400, 300 and 300 1C: (a) general view, 1 and 2 are the grains 1 and 2, respectively, (b) thin twins formed in the grain with high dislocation density, diffraction pattern with [110]γ zone axis, (c) two twins variants, diffraction pattern with [110]γ zone axis, (d) bright and (e) dark field images of thick twins with nano-twins inside, zone axis is [110]γ.
Fig. 7. Bright (a) and dark (b) field TEM images of the microstructure after one pass at 200 1C showing: (a) formation of the microbands after, [110]γ zone axis and (b) nanotwins formation (outlined area) within the microband; and the microstructure formed after two passes at 400 1C and 200 1C: (c) general view, (d) formation of cells, (e, f) formation of primary variant deformation twins and secondary variant twins indicated by arrows, zone axis for (e) is [110]γ and zone axis for (f) is [120]γ
I.B. Timokhina et al. / Materials Science & Engineering A 593 (2014) 163–169
twins played an important role in reducing the mean free path of dislocations. As published in [31], finer and denser twins provide more interfaces between the matrix and twins to accommodate more sessile dislocations leading to a composite effect and finally resulting in higher work hardening rates [31]. The coarse twins had a distance between them of 80 720 nm. It is important to highlight that the coarse twins were also characterized by the presence of {111}γ nano-twins inside the primary twins, Fig. 6d and e. It is suggested that the formation of the nano-twins inside the microtwins during the additional pass at 300 1C was due to temperature and the rotation of the shear plane to 451 anticlockwise at this. The work hardening rate after this deformation schedule was the highest and showed a rapid decrease associated with the dense deformation substructure, Fig. 3b, i.e. the microstructure showed its strain limit and was not able to produce new mobile dislocations during testing. Although the mechanical testing of the samples involving the first or second pass at 200 1C was not possible, it was possible to study the microstructures of steel after these two ECAP conditions. The SEM images showed elongated austenite grains with a thickness of 25 75 μm and the formation of twins in the grain interior, Fig. 1d. However, TEM revealed that microband formation and not twinning was the representative microstructural feature in the microstructure after one pass at 200 1C, Fig. 7a. The microband thickness was 50 720 nm. However, twin formation was also observed within a number of the austenite grains. Moreover, the diffraction pattern analysis and dark field TEM images enabled the observation of the nano-twins within the microbands, as well as micro-twins, Fig. 7b. The diffraction pattern with a zone axis of [110]γ from the nano-twins showed extra diffraction spots that could be evidence of twinning. It is important to emphasize that the nano-twins did not cross the microband boundary. The two passes, first at 400 1C followed by the pass at 200 1C led to the formation of a complex microstructure consisting of deformation cells, Fig. 7c, deformation twins and a second twin variant, Fig. 7. The distance between the primary twins varied from 60 to 270 nm, Fig. 3b and c. The two pass deformation induced the formation of a second variant of the nano-twin within the primary deformation twins, Fig. 7d. However, microbands were also observed in a number of grains. This high density of twinning explains the extraordinary hardening which prevented the possibility of continuous deformation. 4. Conclusions The microstructure–property relationship in the TWIP steel after ECAP tests at elevated temperatures has been studied for the
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first time. The deformation mechanism of the samples after different deformation schedules has been identified. The results show that there is the competition between the effect of the deformation temperature, grain size and accumulated strain on the mechanism of deformation. The wide range of the dislocation substructure such as parallel microbands, dislocation cells, subgrains, stacking faults etc. was formed as a function of ECAP schedule. Twinning was observed even at an elevated temperature of 400 1C. The formation of nano-twins within micro-twins was also found. The direct relationship between the dislocation substructure development after different ECAP schedules and strain hardening behavior was identified. References [1] S. Allain, J.-P. Chateau, O. Bouaziz, Mater. Sci. Eng. A 387–389 (2004) 143–147. [2] E. Mazancova, K. Mazanec, Mater. Sci. Eng. A 16 (2) (2009) 26–31. [3] O. Grassel, L. Kruger, G. Frommeyer, L.W. Meyer, Int. J. Plast. 16 (2000) 1391–1409. [4] J.G. Sevillano, Scr. Mater. 60 (2009) 336–339. [5] R. Ueji, N. Tsuchida, D. Terada, N. Tsuji, Y. Tanaka, A. Takemura, K. Kunishige, Scr. Mater. 59 (2008) 963–966. [6] S. Allain, J.P. Chateau, O. Bouaziz, S. Migot, N. Guelton, Mater. Sci. Eng. A 387– 389 (2004) 158–162. [7] S. Vercammen, B. Blanpain, B.C. De Cooman, P. Wollants, Acta Mater. 52 (2004) 2005–2012. [8] K. Sato, M. Ichinose, Y. Hirotsu, Y. Inoue, ISIJ Int. 29 (1989) 868–877. [9] D Barbier, N. Gey, S. Allain, N. Bozzolo, M. Humbert, Mater. Sci. Eng. A 500 (2009) 196–206. [10] Q. Lu, D. Li, S. Feng-shou, Y. Feng, K. Tong, Adv. Mater. Res. 194–196 (2011) 1235–1241. [11] Z. Yang, D. Zhu, W. Yi, S. Lin., C. Du, Adv. Mater. Res. 197–198 (2011) 655–661. [12] G. Dini, R. Ueji, A. Najafizadeh, Mater. Sci. Forum 654–656 (2010) 294–297. [13] I. Gutierrez-Urrutia, D. Raabe, Scr. Mater. 66 (2012) 992–996. [14] L. Remy, Acta Metall. 26 (3) (1978) 443–451. [15] Y. Dastur, W. Leslie, Metall. Mater. Trans. A 12 (5) (1981) 749–759. [16] S. Allain, J.P. Chateau, O. Bouaziz, Mater. Sci. Eng. A 387–389 (2004) 143–147. [17] S. Allain, et al., Mater. Sci. Eng. A 387–389 (2004) 158–162. [18] J.-L. Collet, F. Bley, A. Deschamps, C. Scott, Adv. Mater. Res. 15–17 (2007) 822–827. [19] X. Liang, Structure and Mechanical Properties of Fe–Mn Alloys (Ph.D.), McMaster University, 2008. [20] R. Lapovok, D. Tomus, B.C. Muddle, Mater. Sci. Eng. A 490 (1–2) (2008) 171–180. [21] R. Lapovok, Y. Estrin, M.V. Popov, T.G. Langdon, Adv. Eng. Mater. 10 (5) (2008) 429–433. [22] D.A. Hughes, W.D. Nix, Mater. Sci. Eng. A 122 (1989) 153. [23] D.A. Hughes, Acta Mater. 41 (1993) 1421. [24] I. Gutierrez-Urrutia, D. Raabe, Scr. Mater. 69 (2013) 53–56. [25] I. Karaman, H. Sehitoglu, H.J. Maier, Y.I. Chumlyakov, Acta Mater. 49 (2001) 3919–3933. [26] Z.B. Yang, D.Y. Zhu, W.F. Yi, S.M. Lin, C.M. Du, Adv. Mater. Res. 197–198 (2011) 655–661. [27] M.A. Meyers, O. Vöhringer, V.A. Lubarda, Acta Mater. 49 (2001) 4025–4039. [28] Y.T. Zhu, X.Z. Liao, X.L. Wu, J. Narayan, J. Mater. Sci. 48 (2013) 4467–4475. [29] Y.T. Zhu, X.Z. Liao, X.L. Wu, Prog. Mater. Sci. 57 (2012) 1–62. [30] I. Gutierrez-Urrutia, D. Raabe, Acta Mater. 59 (2011) 6449–6462. [31] H. Idrissi, K. Renard, D. Schryvers, P.J. Jacques, Scr. Mater. 63 (2010) 961–964.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Severe plastic deformation of a TWIP steel I.B. Timokhina a,n, A. Medvedev b, R. Lapovok b a b
Institute for Frontier Materials, Deakin University, Geelong, Victoria 3217, Australia CAHM, Materials Engineering Department, Monash University, Clayton, Victoria 3800, Australia
art ic l e i nf o
a b s t r a c t
Article history: Received 3 October 2013 Received in revised form 4 November 2013 Accepted 7 November 2013 Available online 17 November 2013
The severe plastic deformation of a Twinning Induced Plasticity (TWIP), 0.61C–22.3Mn–0.19Si–0.14Ni– 0.27Cr (wt%) steel by Equal Channel Angular Pressing (ECAP) at elevated temperatures was used to study the deformation mechanism as a function of accumulated strain and processing parameters. The relationship between the microstructures after different deformation schedules of ECAP at the temperatures of 200, 300 and 400 1C, strain hardening behavior and mechanical properties was studied. The best balance between strength and ductility (1702 MPa and 24%) was found after two passes at 400 1C and 300 1C of ECAP. It was due to the formation of deformation microbands and twins in the microstructure. The twinning was observed after all deformation schedules except after one pass at 400 1C. The important finding was the formation of twins in the ultrafine grains. Moreover, the stacking faults were observed in the subgrains with the size of 50 nm. It is also worth mentioning the formation of nano-twins within the micro-twins at the same time. It was found that the deformation schedule affects the dislocation substructure with formation of deformation bands, cells, subgrains, two variants of twins that, in turn, influence the strain-hardening behavior and mechanical properties. & 2013 Published by Elsevier B.V.
Keywords: Twinning Induced Plasticity steels Equal Channel Angular Pressing Mechanical properties Transmission electron microscopy Micro/nano-twins Dislocation substructure
1. Introduction The great interest in Twinning Induced Plasticity (TWIP) steels from the automotive industry is inspired by their high hardening rate during deformation, high strength and good ductility which together provides superior toughness compared to other types of steel [1]. TWIP steels have low stacking fault energy (SFE) between 15 and 40 mJ/m2 at room temperature through the high level of manganese, in the range of 15–30 wt% [2–6]. The high Mn content stabilizes the austenite at room temperature and produces a single phase face of centered cubic (fcc) steel with low stacking fault energy. It has been shown that in the case of a TWIP alloy with manganese content of 23 wt% and carbon content 0.65 wt%, as in the current research, the SFE is 23 mJ/m2. This SFE reduces the ease of cross slip upon deformation allowing twinning to become the favored deformation mechanism [7,8]. On the other hand, for this steel composition, the SFE was not low enough (γfcc r12 mJ/ m2) to promote the martensite transformation and, hence, this was suppressed in the current steel. The deformation behavior of TWIP steels has normally been examined by Barbier in interrupted tension [9] or compression tests [10] with different grain sizes in samples produced by the hot rolling and/or cold rolling followed by annealing [11,12]. In such n
Corresponding author. Tel.: +61 3 9905 9608; fax: +61 3 9905 4940. E-mail address: [email protected] (I.B. Timokhina).
0921-5093/$ - see front matter & 2013 Published by Elsevier B.V. http://dx.doi.org/10.1016/j.msea.2013.11.013
tests the strain that could be achieved was below 100% and grain size ranges varied from 1 to 60 mm and above [10–12]. It is known that the propensity for twinning decreases with a decrease in the grain size [13]. For example, there was no twinning observed in TWIP samples with a grain size of around 2 mm [12] and the strengthening mechanism in this case was attributed only to the dislocation interactions. Although the data on the effect of strain and grain size on the formation of deformation twins obtained in previous investigations is quite clear, the effect of severe plastic deformation, such as e.g. Equal Channel Angular Pressing (ECAP), which introduces a large shear strain above 100% into the material and refines the grains size to the submicron range, on the formation of deformation twins has not been investigated. Since the TWIP steel has high strength and ductility, it is expected that ECAP deformation could only be performed at elevated temperatures. However, an increase in temperature is known to lead to a decrease in the volume fraction of the twins, twin length and twin thickness [14,15]. Two of three Gibbs energy components of SFE are temperature dependent. As has been published in [5,16,17], the change in the temperature leads to changes in the SFE which promotes different deformation mechanisms. For example, an increase in the deformation temperature to 400 1C suppresses twinning and favours dislocation glide [18]. In addition, it also has been reported [19] that there is a critical strain for the onset of twinning, and this increases with the temperature. However, the behavior of ultrafine grained materials
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Fig. 1. Optical (a) and SEM (b–d) images of initial microstructure (a), after one ECAP pass at 400 1C (b), after four ECAP passes at 400 1C (c) and after one pass at 200 1C (d) of TWIP steel samples.
is quite often quite different from what is observed in course grained materials. In this paper for the first time, severe plastic deformation by ECAP has been applied to a TWIP steel. The formation of the ultrafine grained microstructure is discussed in relation to the strain history and temperature of material processing. Different temperatures and number of ECAP passes were used to manipulate the microstructure and mechanical properties.
Table 1 ECAP schedules.
2. Material and experimental procedure
study [5] to promote twinning deformation after one pass of ECAP. The changes in the mechanical properties and microstructures of the TWIP steel as a function of processing routes (Table 1) at elevated temperatures were studied. Sub-sized round tensile samples were machined out of the ECAP processed billets with a gage length of 10 mm. The tensile test was carried out using an Instron 5982 machine with a cross-head velocity of 1 mm/min corresponding to a nominal strain rate of about 0.002 s 1 for all conditions listed in Table 1. Materials after two conditions with the first and second passes at 200 1C were not mechanically tested. This was associated with the drastic increase in strength during deformation that led to only partial ECAP before the punch broke down. The lack of sufficient material volume processed by ECAP at 200 1C made it impossible to machine tensile samples but the microscopic study was been carried out. The grain size of the selected samples was measured using a FEI Quanta 3D FEG FIB scanning electron microscope with EBSD at a 15 kV accelerating voltage. The analysis of the
The composition of TWIP steel utilized in this study was Fe– 0.61C–22.3Mn–0.19Si–0.14Ni–0.27Cr (wt%). The 10 mm plates were supplied by Arcelor Mittal (France), which were processed by hot rolling from continuously cast ingot and annealed according to their industrial schedule. The initial microstructure was an equiaxed shape austenite grains with the size of 34721 mm, Fig. 1a. Cylinders of 10 mm diameter and 35 mm length were cut from the plate in the rolling direction and subjected to the different schedules of ECAP passes as given in Table 1. The isothermal ECAP rig used for processing is described elsewhere [20,21]. The deformation temperature of 400 1C was used to obtain grain refinement without triggering twinning. The transition temperature from twinning to slip deformation mode was chosen based on the data reported by others for a TWIP steel with similar composition [5]. The temperature of 200 oC was selected based on a previous
Number of ECAP passes
1 (1C)
1 (1C)
2 (1C)
2 (1C)
3 (1C)
4 (1C)
Temperature during ECAP passes
400
200
400 200
400 300
400 300 300
400 400 400 400
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dislocation substructure formed after different schedules of the ECAP passes was performed using a Philips CM 20 operation at 200 kV with a condenser aperture of 100 μm nominal diameter and a nominal beam diameter of 55 nm. The foils were cut from the center of the sliced samples parallel to the x and y planes. Each TEM foil was prepared by twin-jet polishing in a solution of 5% perchloric acid and 95% methanol using the Struers Tenupol 5 electropolisher. The operating voltage was 50 V at a temperature of 25 1C.
3. Results and discussion 3.1. Mechanical properties and strain hardening behavior of the samples after ECAP The engineering stress–strain curves in the initial hot rolled conditions showed an exceptional ductility of 123% and quite high strength of 930 7 5 MPa, Fig. 2b. After one ECAP pass at 400 1C the ductility decreased to 44%, while the strength increased to 1286 710 MPa, Fig. 2a. The second pass at 300 1C increased the strength further to 170278 MPa but decreased the ductility to 24%. However, the third pass at 300 1C had an opposite effect on strength, i.e. the strength dropped to 16247 10 MPa, while ductility also decreased to 21%, Fig. 2a. It should be noted that the samples between the passes were heated for about 5 min and some recovery might take place. Due to this recovery the strength after four passes of ECAP at 400 1C decreased to 15737 6 MPa, while the ductility was similar (25%) compared to those samples that were deformed at the temperature of 300 1C in the following ECAP passes, Fig. 2a. The properties of the TWIP steel in as-received and deformed conditions are summarized in Table 2 It can be seen that yield stress after processing by ECAP increases significantly from 312 MPa to 980–1480 MPa. The strain hardening rate normalized by shear modulus vs. difference between flow and yield stresses (true stresses) for all deformed states and the as-received condition is shown in Fig. 3. As can be seen in Fig. 3a, the character of strain hardening rate of the TWIP steel in the as-received condition, is qualitatively similar to that described by others [13] i.e. firstly, the curve showed a high overall strain hardening rate; secondly, the curve showed a local minima at intermediate strains and an increase in the strain hardening rate at higher deformation [13]. In particular, the initial value of the strain hardening rate equal to G/40, decreases to a minimum of G/50 when ðs sy Þ reaches 200 MPa and then starts to rise steadily to the value of G/28 and keeps constant for ðs sy Þ between 1200 and 1600 MPa. At ðs sy Þ higher than 1600 MPa the strain hardening rate drops rapidly. This curve behavior appeared to be associated with both dislocation accumulation and twin substructure formation [13], Fig. 3a. The strain hardening rate behavior for the samples processed by different ECAP schedules was different, Fig. 3b. The strain rate
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curves after one pass and four passes at 400 1C showed similar behaviors i.e. after the monotonic drop, the strain hardening rate stabilizes and remains constant for ðs sy Þ in the range of 200– 700 MPa and 400–600 MPa, respectively, Fig. 3b. However, the overall strain hardening rate was lower after one pass at 400 1C. The strain hardening rate curve after one pass at 400 1C followed by one pass at 300 1C initially revealed a steady strain hardening rate of G/5.5 up to 200 MPa with a continuous decrease at higher values of ðs sy Þ. The strain hardening rate curve after one pass at 400 1C followed by two passes at 300 1C demonstrated a rapid decrease from the initial value ðG=6⋯G=5Þ, Fig. 3b. 3.2. Microstructural characterization of the samples after ECAP The SEM of the samples after one pass at 400 1C showed the formation of elongated austenite grains with a thickness from 6 to 11 μm (Fig. 1b). The TEM of the samples from the similar deformation condition revealed the formation of parallel microbands separated by a distance of 260 7 37 nm, indicated by white arrows in Fig. 4a, as a representative microstructural feature. The microbands were thick and parallel within 21 to the {111}γ plane. A SAD pattern with [110]γ zone axis, obtained from the microbands, did not show extra diffraction sots that could be attributed to the twinning. Similar microbands were observed by others in the TWIP steels after high strain rate experiments [22,23] and were described as “first generation microbands” [22,23]. Dislocations with a specific cell configuration were observed inside the microbands, indicated by black arrows in Fig. 4b. As reported previously [24], these dislocations cells can play an important role in the strain hardening behavior due to the forest hardening associated with short range dislocation interactions [24]. This hardening mechanism results in a high critical stress when transferring plastic deformation across the microbands [24]. SEM of the microstructure after four passes at 400 1C showed the essential refinement of the austenite grains with grain size variation from 0.3 to 0.6 μm, Fig. 1c. The analysis of the microstructure using TEM after four passes at 400 1C revealed that the Table 2 Yield stress, strength and ductility (engineering) of TWIP steel at different conditions. ECAP schedules
sy (MPa)
su (MPa)
ε (%)
As-received 1 Pass (400 1C) 1 Pass (200 1C) 2 Passes (400 1C þ 200 1C) 2 Passes (400 1C þ 300 1C) 3 Passes (400 1C þ 3001C þ 3001C) 4 Passes (400 1C þ 400 1Cþ 400 1C þ400 1C)
312 980 – – 1250 1480 1290
930 1286 – – 1702 1624 1573
123 44 – – 24 21 25
Fig. 2. Engineering stress vs. engineering strain tensile curves for the samples processed by ECAP (a); sample in as-received condition (b).
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Fig. 3. Strain hardening rate (normalized by shear modulus) vs. difference between true and yield stresses for sample in as-received condition (a); samples processed by ECAP (b).
Fig. 4. TEM micrographs of the microstructure formed after one pass (a) and four passes (b–h) at 400 1C: (a) formation of microbands, [110]γ zone axis, white arrows indicate the microband walls and black arrows indicate the dislocations inside microbands, (b) formation of subgrains with [110]γ zone axis and free of dislocations, (c) and (d) subgrains free and containing parallel sessile dislocations, (e) and (f) formation of stacking faults in the subgrain, (g) bright and (h) dark field images of the twins formed in the microstructure.
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representative microstructural feature was the formation of subgrains with an average size of 507 10 nm within the austenite grains. The subgrains showed the low angle of grain boundary misorientation between them and high angle of grain boundary misorientation with respect to the matrix, Fig. 4b and c. The subgrains could roughly be divided into three groups: (i) free of dislocations, Fig. 4c; (ii) containing parallel sessile dislocations, which should be the result of high strain accommodation in the lattice, Fig. 4d; and (iii) containing stacking faults, Fig. 4e and f. It is believed that the microstructural recovery, which occurred at the temperature of 400 1C between each pass, should be taken into account in understanding the mechanism of the subgrains formation. The microbands, which were the representative dislocation substructure after one pass at 400 1C, were not observed in the microstructure of the samples after four passes at 400 1C. However, the microband walls could be the preferential nucleation sites for subgrain formation. The formation of the stacking faults, Fig. 4f (indicated by arrow), in 50 nm subgrains was not expected and appeared to be an intriguing feature. Despite the general understanding that twinning is absent at high temperatures such as 400 1C, the formation of fine twins crossing the austenite boundary was observed in the microstructure, Fig. 4g and e. The twins were surrounded by an area with high dislocation density, which could be associated with the stress propagation during the twin formation. The strain hardening rate curve behavior after four passes was similar to that after one pass at 400, although the average strain hardening rate was higher after four passes than after one pass. Moreover, a significant increase in yield and tensile strength after four passes (Table 2) was observed compared to the samples after one pass at 400 1C due to the formation of the more complex dislocation substructure with the formation of subgrains and thin twins. It should be noted that the formation of twins after four passes of ECAP at 400 1C could not be explained from a SFE point of view as this is proposed to be in the range of 12–45 mJ/m2 to favor the twinning mechanism. The calculated SFE for a similar TWIP steel (Fe–22 wt% Mn–0.6 wt% C at 300–400 1C) at similar temperature, carried out in [1], was much higher than the SFE for twinning. Twinning is normally not expected for coarse grained material for these temperatures. Moreover, a decrease in the grain size normally suppresses twinning [16] and, hence, the major reduction in the grain size after four passes should make twinning more difficult [25,26]. Therefore, both the temperature and grain size effects should lead to the suppression of twinning after four passes of ECAP at 400 1C. For coarse grained materials, both deformation mechanisms, dislocation slip and twinning, follow the Hall–Petch
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relationship with slope of the curve steeper for twinning [27], that makes twinning more difficult with grain size decreasing. However, as shown in several publications by Zhu et. al. [28,29] there is a transition of ultrafine grain size (UFG), which promotes the formation of twins in TWIP steels. Moreover, the formation of deformation twins was not only experimentally observed but also modeled in UFG materials, with several twinning mechanisms proposed [29]. Deformation twinning has been shown to be a major plastic deformation mechanism for materials with a grain size below 100 nm [28]. This is consistent with microstructural observations in the current research, when the several ECAP passes were performed at 400 1C. The representative microstructural features in the samples after one pass at 400 1C followed by one pass at 300 1C were microband formation along with the formation of twins, Fig. 5. The microbands had an average thickness of 20 75 nm and showed low grain boundary misorientation, Fig. 5b. The microbands were formed at an angle of 6575o with the respect to the matrix. It appeared that the microbands nucleated at the grain boundary and then propagated towards the grain interior, Fig. 5b. Straight dislocations were observed within the microbands, Fig. 5b. The diffraction pattern with the [110]γ zone axis did not show extra diffraction spots and the dislocation morphology did not change during sample tilting. This suggests that these were sessile dislocations that might have formed by the stair-rod cross slip mechanism [30]. The stair-rod dislocations are normally considered as the twinning precursors [30]. The micro-twins with an average spacing of 25 75 nm were also observed in the microstructure, Fig. 5a. The mechanical properties of the samples after one pass at 400 1C followed by one pass at 300 1C were similar to those after four passes at 400 1C, although the strain-hardening rate behavior was different. This was associated with the difference in the dislocation substructure formed after the two different deformation schedules. The addition of the third pass at 300 1C resulted in the formation of grains with high dislocation density, as can be seen in grain 1, Fig. 6a, and grains with twins, as observed in grain 2, Fig. 6a. Two types of micro-twin: (i) fine twins, Fig. 6b, c, and (ii) coarse twins, Fig. 6d, e, were detected The fine twins with an average distance between the twins of 207 5 nm were observed in the austenite grains with high dislocation density, Fig. 6b. However, in most of the austenite grains, two variants of fine deformation twins perpendicular to each other were observed: (i) the primary twins shown by the arrows in Fig. 6c with a spacing of 100 nm, and (ii) the secondary twin variant with a spacing of 50 nm, Fig. 6c. The formation of two variants of deformation
Fig. 5. TEM micrographs and diffraction patterns with [110]γ zone axis of twins formed in grains 2 and 3 (a), and microbands observed in grain 1 and microbands enlarge view (b) with zone axis [110]γ after two passes at 400 1C and 300 1C. Dot line represents the grain boundary.
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Fig. 6. TEM micrographs of the microstructure formed after passes at 400, 300 and 300 1C: (a) general view, 1 and 2 are the grains 1 and 2, respectively, (b) thin twins formed in the grain with high dislocation density, diffraction pattern with [110]γ zone axis, (c) two twins variants, diffraction pattern with [110]γ zone axis, (d) bright and (e) dark field images of thick twins with nano-twins inside, zone axis is [110]γ.
Fig. 7. Bright (a) and dark (b) field TEM images of the microstructure after one pass at 200 1C showing: (a) formation of the microbands after, [110]γ zone axis and (b) nanotwins formation (outlined area) within the microband; and the microstructure formed after two passes at 400 1C and 200 1C: (c) general view, (d) formation of cells, (e, f) formation of primary variant deformation twins and secondary variant twins indicated by arrows, zone axis for (e) is [110]γ and zone axis for (f) is [120]γ
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twins played an important role in reducing the mean free path of dislocations. As published in [31], finer and denser twins provide more interfaces between the matrix and twins to accommodate more sessile dislocations leading to a composite effect and finally resulting in higher work hardening rates [31]. The coarse twins had a distance between them of 80 720 nm. It is important to highlight that the coarse twins were also characterized by the presence of {111}γ nano-twins inside the primary twins, Fig. 6d and e. It is suggested that the formation of the nano-twins inside the microtwins during the additional pass at 300 1C was due to temperature and the rotation of the shear plane to 451 anticlockwise at this. The work hardening rate after this deformation schedule was the highest and showed a rapid decrease associated with the dense deformation substructure, Fig. 3b, i.e. the microstructure showed its strain limit and was not able to produce new mobile dislocations during testing. Although the mechanical testing of the samples involving the first or second pass at 200 1C was not possible, it was possible to study the microstructures of steel after these two ECAP conditions. The SEM images showed elongated austenite grains with a thickness of 25 75 μm and the formation of twins in the grain interior, Fig. 1d. However, TEM revealed that microband formation and not twinning was the representative microstructural feature in the microstructure after one pass at 200 1C, Fig. 7a. The microband thickness was 50 720 nm. However, twin formation was also observed within a number of the austenite grains. Moreover, the diffraction pattern analysis and dark field TEM images enabled the observation of the nano-twins within the microbands, as well as micro-twins, Fig. 7b. The diffraction pattern with a zone axis of [110]γ from the nano-twins showed extra diffraction spots that could be evidence of twinning. It is important to emphasize that the nano-twins did not cross the microband boundary. The two passes, first at 400 1C followed by the pass at 200 1C led to the formation of a complex microstructure consisting of deformation cells, Fig. 7c, deformation twins and a second twin variant, Fig. 7. The distance between the primary twins varied from 60 to 270 nm, Fig. 3b and c. The two pass deformation induced the formation of a second variant of the nano-twin within the primary deformation twins, Fig. 7d. However, microbands were also observed in a number of grains. This high density of twinning explains the extraordinary hardening which prevented the possibility of continuous deformation. 4. Conclusions The microstructure–property relationship in the TWIP steel after ECAP tests at elevated temperatures has been studied for the
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first time. The deformation mechanism of the samples after different deformation schedules has been identified. The results show that there is the competition between the effect of the deformation temperature, grain size and accumulated strain on the mechanism of deformation. The wide range of the dislocation substructure such as parallel microbands, dislocation cells, subgrains, stacking faults etc. was formed as a function of ECAP schedule. Twinning was observed even at an elevated temperature of 400 1C. The formation of nano-twins within micro-twins was also found. The direct relationship between the dislocation substructure development after different ECAP schedules and strain hardening behavior was identified. References [1] S. Allain, J.-P. Chateau, O. Bouaziz, Mater. Sci. Eng. A 387–389 (2004) 143–147. [2] E. Mazancova, K. Mazanec, Mater. Sci. Eng. A 16 (2) (2009) 26–31. [3] O. Grassel, L. Kruger, G. Frommeyer, L.W. Meyer, Int. J. Plast. 16 (2000) 1391–1409. [4] J.G. Sevillano, Scr. Mater. 60 (2009) 336–339. [5] R. Ueji, N. Tsuchida, D. Terada, N. Tsuji, Y. Tanaka, A. Takemura, K. Kunishige, Scr. Mater. 59 (2008) 963–966. [6] S. Allain, J.P. Chateau, O. Bouaziz, S. Migot, N. Guelton, Mater. Sci. Eng. A 387– 389 (2004) 158–162. [7] S. Vercammen, B. Blanpain, B.C. De Cooman, P. Wollants, Acta Mater. 52 (2004) 2005–2012. [8] K. Sato, M. Ichinose, Y. Hirotsu, Y. Inoue, ISIJ Int. 29 (1989) 868–877. [9] D Barbier, N. Gey, S. Allain, N. Bozzolo, M. Humbert, Mater. Sci. Eng. A 500 (2009) 196–206. [10] Q. Lu, D. Li, S. Feng-shou, Y. Feng, K. Tong, Adv. Mater. Res. 194–196 (2011) 1235–1241. [11] Z. Yang, D. Zhu, W. Yi, S. Lin., C. Du, Adv. Mater. Res. 197–198 (2011) 655–661. [12] G. Dini, R. Ueji, A. Najafizadeh, Mater. Sci. Forum 654–656 (2010) 294–297. [13] I. Gutierrez-Urrutia, D. Raabe, Scr. Mater. 66 (2012) 992–996. [14] L. Remy, Acta Metall. 26 (3) (1978) 443–451. [15] Y. Dastur, W. Leslie, Metall. Mater. Trans. A 12 (5) (1981) 749–759. [16] S. Allain, J.P. Chateau, O. Bouaziz, Mater. Sci. Eng. A 387–389 (2004) 143–147. [17] S. Allain, et al., Mater. Sci. Eng. A 387–389 (2004) 158–162. [18] J.-L. Collet, F. Bley, A. Deschamps, C. Scott, Adv. Mater. Res. 15–17 (2007) 822–827. [19] X. Liang, Structure and Mechanical Properties of Fe–Mn Alloys (Ph.D.), McMaster University, 2008. [20] R. Lapovok, D. Tomus, B.C. Muddle, Mater. Sci. Eng. A 490 (1–2) (2008) 171–180. [21] R. Lapovok, Y. Estrin, M.V. Popov, T.G. Langdon, Adv. Eng. Mater. 10 (5) (2008) 429–433. [22] D.A. Hughes, W.D. Nix, Mater. Sci. Eng. A 122 (1989) 153. [23] D.A. Hughes, Acta Mater. 41 (1993) 1421. [24] I. Gutierrez-Urrutia, D. Raabe, Scr. Mater. 69 (2013) 53–56. [25] I. Karaman, H. Sehitoglu, H.J. Maier, Y.I. Chumlyakov, Acta Mater. 49 (2001) 3919–3933. [26] Z.B. Yang, D.Y. Zhu, W.F. Yi, S.M. Lin, C.M. Du, Adv. Mater. Res. 197–198 (2011) 655–661. [27] M.A. Meyers, O. Vöhringer, V.A. Lubarda, Acta Mater. 49 (2001) 4025–4039. [28] Y.T. Zhu, X.Z. Liao, X.L. Wu, J. Narayan, J. Mater. Sci. 48 (2013) 4467–4475. [29] Y.T. Zhu, X.Z. Liao, X.L. Wu, Prog. Mater. Sci. 57 (2012) 1–62. [30] I. Gutierrez-Urrutia, D. Raabe, Acta Mater. 59 (2011) 6449–6462. [31] H. Idrissi, K. Renard, D. Schryvers, P.J. Jacques, Scr. Mater. 63 (2010) 961–964.