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Effect of boundaries on toughness in high-strength low-alloy steels from the view of crystallographic misorientation
Materials Letters 259 (2020) 126841
Contents lists available at ScienceDirect
Materials Letters journal homepage: www.elsevier.com/locate/mlblue
Effect of boundaries on toughness in high-strength low-alloy steels from the view of crystallographic misorientation Xiucheng Li a, Jingxiao Zhao a, Jingliang Wang a, Xuelin Wang a, Shilong Liu b,⇑, Chengjia Shang a,⇑ a b
Collaborative Innovation Center of Steel Technology, University of Science and Technology Beijing, Beijing 100083, China Institute of Advanced Steels and Materials, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China
a r t i c l e
i n f o
Article history: Received 30 September 2019 Received in revised form 16 October 2019 Accepted 18 October 2019 Available online 19 October 2019 Keywords: Grain boundaries Microstructure Toughness Ductile-to-brittle transition Misorientation
a b s t r a c t An abnormal variation of ductile-to-brittle transition temperature (DBTT) with density of high angle grain boundaries (HAGBs) was found in a high-strength low-alloy (HSLA) steel. Further analysis on special misorientation angle (SMA) indicates that HAGBs always have high SMA between {1 0 0}-planes, but do not definitely have high SMA between {1 1 0}-planes. When the boundaries were quantitatively recounted according to the SMA, the sample with lower DBTT exhibited higher density of boundaries with higher SMA between {1 1 0}-planes. Therefore, the density of HAGBs could not reflect the change of DBTT sufficiently, DBTT was essentially influenced by the distribution of SMA between {1 0 0}planes and SMA between {1 1 0}-planes. Ó 2019 Published by Elsevier B.V.
1. Introduction Toughness is very important in safety concern for high-strength low-alloy (HSLA) steels application, especially when the steels are used in low temperature environment. Because when the temperature decreases, the fracture strength of HSLA steels could be lower than the yield strength and thus little ductile plastic deformation takes place before fracture [1]. Normally, ductile-to-brittle transition temperature (DBTT) is an important consideration in steels application and steels with lower DBTT are always wanted. In the 1950s, Petch [2] proposed that DBTT was determined by grain size. Afterwards, grain refinement was demonstrated by extensive studies on toughness improvement for steels. However, the definition of ‘‘grain size” is not so straightforward for bainitic/martensitic steels due to their complex hierarchy in morphology and crystallography [3–5]. With the development of electron back-scattered diffraction (EBSD) technique, overall misorientation angle (OMA) was proposed as a criterion for characterizing grain boundaries [6]. It has been widely reported that high angle grain boundaries (HAGBs) contribute to the toughness greatly. However, there are still different viewpoints on the threshold for defining HAGBs, some researchers believe that OMA of HAGBs should be greater than 15° [7,8], while the others consider that only the HAGBs with OMA greater than 40° [9] or 45° [10–12] can con⇑ Corresponding authors. E-mail addresses: [email protected] (S. Liu), [email protected] (C. Shang). https://doi.org/10.1016/j.matlet.2019.126841 0167-577X/Ó 2019 Published by Elsevier B.V.
tribute greatly to the toughness of steels. Actually, OMA is too general to reflect the crystallographic features of polycrystalline. As pointed by Morris et al. [4], the angle between {1 0 0} crystallographic planes is more relevant than OMA when defining the effective grain size for cleavage fracture. Guo et al. [13] also stated the effective grain size for plastic deformation was related to the mean free path of dislocation glide along {1 1 0}-planes, but not directly controlled by OMA. Therefore, it is more reasonable to investigate specific misorientation angle (SMA), i.e. the misorientation of specific crystallographic planes, in building the relationship between crystallographic characteristics and properties. In this paper, we attempted to use SMA between the {1 0 0} cleavage planes ({1 0 0}-SMA), as well as SMA between {1 1 0} slip planes ({1 1 0}-SMA) of two neighboring crystallographic units as criteria to count the density of boundaries, and investigate the relevance between SMA and DBTT.
2. Experimental The chemical composition (weight percentage) of experimental steel [14,15] was 0.07%C, 0.23%Si, 1.06%Mn, 1.25%Cu, 1.74%Ni, 1.04%Cr, 0.49%Mo, 0.049%Nb, and balance Fe. Using Gleeble 3800 thermo-mechanical simulator, steel blanks were heated at 5 °C/s up to temperatures of 895/1000/1250 °C separately, held for 5 min, and followed by continuous cooling at 5 °C/s down to 500 °C and 2.5 °C/s down to room temperature. The samples corresponding to three austenitizing temperatures were denoted
2
X. Li et al. / Materials Letters 259 (2020) 126841
3. Results and discussion
Fig. 1. Ductile-to-brittle transition curves of three samples.
as A895, A1000 and A1250 respectively. Standard V-notch Charpy impact specimens were tested from 120 °C to room temperature. DBTT was calculated as the temperature with intermediate absorption energy of upper and lower shelf [16]. EBSD measurements were conducted employing a JSM-7001F field emission scanning electron microscope equipped with an EBSD camera, charactering area was 1.25 104 lm2 for each sample and step was 0.1 lm. A Python program was written to compute OMA and SMA of boundary.
Fig. 1 shows the impact absorption energy at different testing temperatures and ductile-to-brittle transition curves of three samples. All three samples exhibit classic ductile to brittle transition phenomenon with testing temperature decreasing. Three samples have different DBTTs and the change of DBTTs with austenitizing temperatures is not monotonous, A1000 has the lowest DBTT of 94 °C and A1250 has the highest DBTT of 52 °C. A similar result of toughness influenced by cooling rate was reported by You et al. [10] and it was explained with the difference in density of HAGBs (OMA > 45 °C). Fig. 2 shows the band contrast maps with HAGBs (Fig. 2 (a), (b) and (c)) and the statistic results of boundaries with OMA (Fig. 2 (d)) of three samples. Obviously, the density of boundaries exhibits a bimodal distribution showing two peaks at very low and high OMA. The total density of HAGBs (>15°) and HAGBs (45°) were counted separately. By either way to count HAGBs, A1250 shows the highest density and A895 shows the lowest. However, the sample has the lowest DBTT is A1000 but not A1250, therefore, it can be inferred that the relationship between OMA and DBTT is not stable. Fig. 3 shows the densities of boundaries characterized by {1 0 0}-SMA and {1 1 0}-SMA, as well as OMA. The distribution of {1 0 0}-SMA (Fig. 3(b)) shows two density peaks at low and high value, which is similar to the distribution of OMA (Fig. 3(a)). A1250 has higher density than the other two samples, thus this indicates A1250 should have had lower DBTT [4,10,11], however, the experiment results are just the opposite. Then, if the boundaries was counted according to {1 1 0}-SMA (Fig. 3(c)), there is only
Fig. 2. Band contrast maps of microstructure and the boundaries distribution (red line: 15° < OMA < 45°, white line: OMA > 45°) of (a)A895, (b)A1000, (c)A1250, and (d) density of boundaries with OMA. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
X. Li et al. / Materials Letters 259 (2020) 126841
Fig. 3. Boundary density as a function of (a) OMA, (b) {1 0 0}-SMA and (c) {1 1 0}-SMA of adjacent grains.
Fig. 4. Mechanism of toughness improvement by suppressing crack initiation.
3
4
X. Li et al. / Materials Letters 259 (2020) 126841
one boundary density peak for each sample and the peak is at low {1 1 0}-SMA range. Though the peak of A1250 is higher, A1000 and A895 have higher boundary density at high {1 1 0}-SMA range. Therefore, on one hand, high OMA does not mean high SMA, especially high {1 1 0}-SMA, and on the other hand, the difference of boundary density with {1 1 0}-SMA among three samples is more consistent with DBTT. The difference of OMA and {1 1 0}-SMA distribution among three samples is due to the proportion of special grain boundaries and general grain boundaries [17,18]. For bainitic/martensitic steels, variant boundaries are special boundaries which have specific misorientation and some of them have low R (reverse density of coincidence sites), while prior austenite grain boundaries are general grain boundaries. According to K-S relationship, {1 1 0}-SMAs of all variant boundaries within a packet are close to 0° [5]. Obviously, A1250 have higher density of variant boundaries comparing with A895 and A1000, and part of these variant boundaries have very low {1 1 0}-SMA. This could explain why more boundaries have high OMA but low {1 1 0}-SMA in A1250. The toughening mechanism of boundaries with high {1 1 0}SMA is through plasticity improvement of steels, and this effect is particularly significant in Charpy impact test. As illustrated in Fig. 4, in the crack initiation stage, when the local stress concentration at the notch reaches fracture initiation stress (ri), crack forms beneath the notch. There should be a plastic deformation zone, in which the stress is higher than yield strength (ry) and slip can occur. In crack initiation stage, boundaries with high {1 1 0}-SMA could hinder the dislocations gliding, thus decrease stress concentration at notch. The hindrance effect on dislocation gliding depends on the strength of boundaries [19,20], which is positively related to {1 1 0}-SMA of boundary [13]. Therefore, A1000 can form larger plastic deformation and absorb more impact energy, because it has higher boundary density at higher {1 1 0}-SMA range. Once the brittle crack has formed and started to propagate, boundaries with high {1 0 0}-SMA would become to the major influencing factor by hindering crack propagation. When crack meets a boundary with high {1 0 0}-SMA, the crack propagation path changes and brittle fracture strength increase from rf to rf0 but the yield strength changes little, thus more energy could be absorbed. However, as Li et al. [21] and Liu et al. [15] reported, large proportion of impact energy was absorbed in crack initiation stage in Charpy impact test. Therefore, though A1250 has higher boundary density with high {1 0 0}-SMA which means A1250 has stronger effect on hindering crack propagation, lower DBTT could still be achieved with A1000 because of the contribution of high {1 1 0}-SMA boundary density in crack initiation stage.
4. Conclusions In summary, for steels in this study, the boundary with high OMA does not definitely have high SMA, especially high {1 1 0}SMA. DBTT is more influenced by boundaries with high {1 0 0}SMA and high {1 1 0}-SMA at the same time, therefore sometimes distribution of OMA could not reflect the change of DBTT sufficiently. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements This work was supported by National Key Research and Development Plan of China (2017YFB0304700). The authors are grateful to Dr. Dongsheng Liu for experimental results sharing and discussions. References [1] J.W. Morris, Stronger, Science 320 (2008) 1022–1023. [2] N.J. Petch, Philo. Mag. 3 (1958) 1128–1136. [3] S. Morito, X. Huang, T. Furuhara, T. Maki, N. Hansen, Acta Mater. 54 (2006) 5323–5331. [4] J.W. Morris, C.S. Lee, Z. Guo, ISIJ Int. 43 (2003) 410–419. [5] H. Kitahara, R. Ueji, N. Tsuji, Y. Minamino, Acta Mater. 54 (2006) 1279–1288. [6] M.C. Kim, J.O. Yong, J.H. Hong, Scr. Mater. 43 (2000) 205–211. [7] B. Hwang, Y.G. Kim, S. Lee, Y.M. Kim, N.J. Kim, J.Y. Yoo, Metall. Mater. Trans. A 36 (2005) 2107–2114. [8] C. Miao, C. Shang, X. Wang, Acta Metall. Sin. 46 (2010) 541–546. [9] A. Lambert, X. Garat, T. Sturel, A.F. Gourgues, A. Gingell, Scr. Mater. 43 (2000) 161–166. [10] Y. You, C. Shang, W. Nie, S. Subramanian, Mater. Sci. Eng. A 558 (2012) 692– 701. [11] A.F. Gourgues, H.M. Flower, T.C. Lindley, Mater. Sci. Tech. 16 (2000) 26–40. [12] X. Li, X. Ma, S.V. Subramanian, C. Shang, R.D.K. Misra, Mater. Sci. Eng. A 616 (2014) 141–147. [13] Z. Guo, C.S. Lee, J.W. Morris, Acta Mater. 52 (2004) 5511–5518. [14] M. Luo, D. Liu, B. Cheng, R. Cao, J. Chen, J. Mater. Eng. Perform. 27 (2018) 4855– 4870. [15] D. Liu, M. Luo, B. Cheng, R. Cao, J. Chen, Metall. Mater. Trans. A 49 (2018) 4918– 4936. [16] Y. Tomita, K. Okabayashi, Metall. Trans. A 17 (1986) 1203–1209. [17] E.L. Maksimova, L.S. Shvindlerman, B.B. Straumal, Acta Metall. 36 (1988) 1573–1583. [18] B.B. Straumal, O.A. Kogtenkova, A.S. Gornakova, V.G. Sursaeva, B. Baretzky, J. Mater. Sci. 51 (2016) 382–404. [19] S. Yang, C. Shang, X. Wang, X. He, Acta Metall. Sin. 39 (2003) 579–584. [20] S. Liu, X. Li, H. Guo, S. Yang, X. Wang, C. Shang, R.D.K. Misra, Philo. Mag. 98 (2018) 934–958. [21] X. Li, X. Ma, S.V. Subramanian, C. Shang, Int. J. Fract. 193 (2015) 131–139.
Contents lists available at ScienceDirect
Materials Letters journal homepage: www.elsevier.com/locate/mlblue
Effect of boundaries on toughness in high-strength low-alloy steels from the view of crystallographic misorientation Xiucheng Li a, Jingxiao Zhao a, Jingliang Wang a, Xuelin Wang a, Shilong Liu b,⇑, Chengjia Shang a,⇑ a b
Collaborative Innovation Center of Steel Technology, University of Science and Technology Beijing, Beijing 100083, China Institute of Advanced Steels and Materials, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China
a r t i c l e
i n f o
Article history: Received 30 September 2019 Received in revised form 16 October 2019 Accepted 18 October 2019 Available online 19 October 2019 Keywords: Grain boundaries Microstructure Toughness Ductile-to-brittle transition Misorientation
a b s t r a c t An abnormal variation of ductile-to-brittle transition temperature (DBTT) with density of high angle grain boundaries (HAGBs) was found in a high-strength low-alloy (HSLA) steel. Further analysis on special misorientation angle (SMA) indicates that HAGBs always have high SMA between {1 0 0}-planes, but do not definitely have high SMA between {1 1 0}-planes. When the boundaries were quantitatively recounted according to the SMA, the sample with lower DBTT exhibited higher density of boundaries with higher SMA between {1 1 0}-planes. Therefore, the density of HAGBs could not reflect the change of DBTT sufficiently, DBTT was essentially influenced by the distribution of SMA between {1 0 0}planes and SMA between {1 1 0}-planes. Ó 2019 Published by Elsevier B.V.
1. Introduction Toughness is very important in safety concern for high-strength low-alloy (HSLA) steels application, especially when the steels are used in low temperature environment. Because when the temperature decreases, the fracture strength of HSLA steels could be lower than the yield strength and thus little ductile plastic deformation takes place before fracture [1]. Normally, ductile-to-brittle transition temperature (DBTT) is an important consideration in steels application and steels with lower DBTT are always wanted. In the 1950s, Petch [2] proposed that DBTT was determined by grain size. Afterwards, grain refinement was demonstrated by extensive studies on toughness improvement for steels. However, the definition of ‘‘grain size” is not so straightforward for bainitic/martensitic steels due to their complex hierarchy in morphology and crystallography [3–5]. With the development of electron back-scattered diffraction (EBSD) technique, overall misorientation angle (OMA) was proposed as a criterion for characterizing grain boundaries [6]. It has been widely reported that high angle grain boundaries (HAGBs) contribute to the toughness greatly. However, there are still different viewpoints on the threshold for defining HAGBs, some researchers believe that OMA of HAGBs should be greater than 15° [7,8], while the others consider that only the HAGBs with OMA greater than 40° [9] or 45° [10–12] can con⇑ Corresponding authors. E-mail addresses: [email protected] (S. Liu), [email protected] (C. Shang). https://doi.org/10.1016/j.matlet.2019.126841 0167-577X/Ó 2019 Published by Elsevier B.V.
tribute greatly to the toughness of steels. Actually, OMA is too general to reflect the crystallographic features of polycrystalline. As pointed by Morris et al. [4], the angle between {1 0 0} crystallographic planes is more relevant than OMA when defining the effective grain size for cleavage fracture. Guo et al. [13] also stated the effective grain size for plastic deformation was related to the mean free path of dislocation glide along {1 1 0}-planes, but not directly controlled by OMA. Therefore, it is more reasonable to investigate specific misorientation angle (SMA), i.e. the misorientation of specific crystallographic planes, in building the relationship between crystallographic characteristics and properties. In this paper, we attempted to use SMA between the {1 0 0} cleavage planes ({1 0 0}-SMA), as well as SMA between {1 1 0} slip planes ({1 1 0}-SMA) of two neighboring crystallographic units as criteria to count the density of boundaries, and investigate the relevance between SMA and DBTT.
2. Experimental The chemical composition (weight percentage) of experimental steel [14,15] was 0.07%C, 0.23%Si, 1.06%Mn, 1.25%Cu, 1.74%Ni, 1.04%Cr, 0.49%Mo, 0.049%Nb, and balance Fe. Using Gleeble 3800 thermo-mechanical simulator, steel blanks were heated at 5 °C/s up to temperatures of 895/1000/1250 °C separately, held for 5 min, and followed by continuous cooling at 5 °C/s down to 500 °C and 2.5 °C/s down to room temperature. The samples corresponding to three austenitizing temperatures were denoted
2
X. Li et al. / Materials Letters 259 (2020) 126841
3. Results and discussion
Fig. 1. Ductile-to-brittle transition curves of three samples.
as A895, A1000 and A1250 respectively. Standard V-notch Charpy impact specimens were tested from 120 °C to room temperature. DBTT was calculated as the temperature with intermediate absorption energy of upper and lower shelf [16]. EBSD measurements were conducted employing a JSM-7001F field emission scanning electron microscope equipped with an EBSD camera, charactering area was 1.25 104 lm2 for each sample and step was 0.1 lm. A Python program was written to compute OMA and SMA of boundary.
Fig. 1 shows the impact absorption energy at different testing temperatures and ductile-to-brittle transition curves of three samples. All three samples exhibit classic ductile to brittle transition phenomenon with testing temperature decreasing. Three samples have different DBTTs and the change of DBTTs with austenitizing temperatures is not monotonous, A1000 has the lowest DBTT of 94 °C and A1250 has the highest DBTT of 52 °C. A similar result of toughness influenced by cooling rate was reported by You et al. [10] and it was explained with the difference in density of HAGBs (OMA > 45 °C). Fig. 2 shows the band contrast maps with HAGBs (Fig. 2 (a), (b) and (c)) and the statistic results of boundaries with OMA (Fig. 2 (d)) of three samples. Obviously, the density of boundaries exhibits a bimodal distribution showing two peaks at very low and high OMA. The total density of HAGBs (>15°) and HAGBs (45°) were counted separately. By either way to count HAGBs, A1250 shows the highest density and A895 shows the lowest. However, the sample has the lowest DBTT is A1000 but not A1250, therefore, it can be inferred that the relationship between OMA and DBTT is not stable. Fig. 3 shows the densities of boundaries characterized by {1 0 0}-SMA and {1 1 0}-SMA, as well as OMA. The distribution of {1 0 0}-SMA (Fig. 3(b)) shows two density peaks at low and high value, which is similar to the distribution of OMA (Fig. 3(a)). A1250 has higher density than the other two samples, thus this indicates A1250 should have had lower DBTT [4,10,11], however, the experiment results are just the opposite. Then, if the boundaries was counted according to {1 1 0}-SMA (Fig. 3(c)), there is only
Fig. 2. Band contrast maps of microstructure and the boundaries distribution (red line: 15° < OMA < 45°, white line: OMA > 45°) of (a)A895, (b)A1000, (c)A1250, and (d) density of boundaries with OMA. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
X. Li et al. / Materials Letters 259 (2020) 126841
Fig. 3. Boundary density as a function of (a) OMA, (b) {1 0 0}-SMA and (c) {1 1 0}-SMA of adjacent grains.
Fig. 4. Mechanism of toughness improvement by suppressing crack initiation.
3
4
X. Li et al. / Materials Letters 259 (2020) 126841
one boundary density peak for each sample and the peak is at low {1 1 0}-SMA range. Though the peak of A1250 is higher, A1000 and A895 have higher boundary density at high {1 1 0}-SMA range. Therefore, on one hand, high OMA does not mean high SMA, especially high {1 1 0}-SMA, and on the other hand, the difference of boundary density with {1 1 0}-SMA among three samples is more consistent with DBTT. The difference of OMA and {1 1 0}-SMA distribution among three samples is due to the proportion of special grain boundaries and general grain boundaries [17,18]. For bainitic/martensitic steels, variant boundaries are special boundaries which have specific misorientation and some of them have low R (reverse density of coincidence sites), while prior austenite grain boundaries are general grain boundaries. According to K-S relationship, {1 1 0}-SMAs of all variant boundaries within a packet are close to 0° [5]. Obviously, A1250 have higher density of variant boundaries comparing with A895 and A1000, and part of these variant boundaries have very low {1 1 0}-SMA. This could explain why more boundaries have high OMA but low {1 1 0}-SMA in A1250. The toughening mechanism of boundaries with high {1 1 0}SMA is through plasticity improvement of steels, and this effect is particularly significant in Charpy impact test. As illustrated in Fig. 4, in the crack initiation stage, when the local stress concentration at the notch reaches fracture initiation stress (ri), crack forms beneath the notch. There should be a plastic deformation zone, in which the stress is higher than yield strength (ry) and slip can occur. In crack initiation stage, boundaries with high {1 1 0}-SMA could hinder the dislocations gliding, thus decrease stress concentration at notch. The hindrance effect on dislocation gliding depends on the strength of boundaries [19,20], which is positively related to {1 1 0}-SMA of boundary [13]. Therefore, A1000 can form larger plastic deformation and absorb more impact energy, because it has higher boundary density at higher {1 1 0}-SMA range. Once the brittle crack has formed and started to propagate, boundaries with high {1 0 0}-SMA would become to the major influencing factor by hindering crack propagation. When crack meets a boundary with high {1 0 0}-SMA, the crack propagation path changes and brittle fracture strength increase from rf to rf0 but the yield strength changes little, thus more energy could be absorbed. However, as Li et al. [21] and Liu et al. [15] reported, large proportion of impact energy was absorbed in crack initiation stage in Charpy impact test. Therefore, though A1250 has higher boundary density with high {1 0 0}-SMA which means A1250 has stronger effect on hindering crack propagation, lower DBTT could still be achieved with A1000 because of the contribution of high {1 1 0}-SMA boundary density in crack initiation stage.
4. Conclusions In summary, for steels in this study, the boundary with high OMA does not definitely have high SMA, especially high {1 1 0}SMA. DBTT is more influenced by boundaries with high {1 0 0}SMA and high {1 1 0}-SMA at the same time, therefore sometimes distribution of OMA could not reflect the change of DBTT sufficiently. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements This work was supported by National Key Research and Development Plan of China (2017YFB0304700). The authors are grateful to Dr. Dongsheng Liu for experimental results sharing and discussions. References [1] J.W. Morris, Stronger, Science 320 (2008) 1022–1023. [2] N.J. Petch, Philo. Mag. 3 (1958) 1128–1136. [3] S. Morito, X. Huang, T. Furuhara, T. Maki, N. Hansen, Acta Mater. 54 (2006) 5323–5331. [4] J.W. Morris, C.S. Lee, Z. Guo, ISIJ Int. 43 (2003) 410–419. [5] H. Kitahara, R. Ueji, N. Tsuji, Y. Minamino, Acta Mater. 54 (2006) 1279–1288. [6] M.C. Kim, J.O. Yong, J.H. Hong, Scr. Mater. 43 (2000) 205–211. [7] B. Hwang, Y.G. Kim, S. Lee, Y.M. Kim, N.J. Kim, J.Y. Yoo, Metall. Mater. Trans. A 36 (2005) 2107–2114. [8] C. Miao, C. Shang, X. Wang, Acta Metall. Sin. 46 (2010) 541–546. [9] A. Lambert, X. Garat, T. Sturel, A.F. Gourgues, A. Gingell, Scr. Mater. 43 (2000) 161–166. [10] Y. You, C. Shang, W. Nie, S. Subramanian, Mater. Sci. Eng. A 558 (2012) 692– 701. [11] A.F. Gourgues, H.M. Flower, T.C. Lindley, Mater. Sci. Tech. 16 (2000) 26–40. [12] X. Li, X. Ma, S.V. Subramanian, C. Shang, R.D.K. Misra, Mater. Sci. Eng. A 616 (2014) 141–147. [13] Z. Guo, C.S. Lee, J.W. Morris, Acta Mater. 52 (2004) 5511–5518. [14] M. Luo, D. Liu, B. Cheng, R. Cao, J. Chen, J. Mater. Eng. Perform. 27 (2018) 4855– 4870. [15] D. Liu, M. Luo, B. Cheng, R. Cao, J. Chen, Metall. Mater. Trans. A 49 (2018) 4918– 4936. [16] Y. Tomita, K. Okabayashi, Metall. Trans. A 17 (1986) 1203–1209. [17] E.L. Maksimova, L.S. Shvindlerman, B.B. Straumal, Acta Metall. 36 (1988) 1573–1583. [18] B.B. Straumal, O.A. Kogtenkova, A.S. Gornakova, V.G. Sursaeva, B. Baretzky, J. Mater. Sci. 51 (2016) 382–404. [19] S. Yang, C. Shang, X. Wang, X. He, Acta Metall. Sin. 39 (2003) 579–584. [20] S. Liu, X. Li, H. Guo, S. Yang, X. Wang, C. Shang, R.D.K. Misra, Philo. Mag. 98 (2018) 934–958. [21] X. Li, X. Ma, S.V. Subramanian, C. Shang, Int. J. Fract. 193 (2015) 131–139.