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Strain localization behavior in low-carbon martensitic steel during tensile deformation
Accepted Manuscript Strain localization behavior in low-carbon martensitic steel during tensile de‐ formation Hyuntaek Na, S. Nambu, M. Ojima, J. Inoue, T. Koseki PII: DOI: Reference:
S1359-6462(13)00435-1 http://dx.doi.org/10.1016/j.scriptamat.2013.08.030 SMM 10040
To appear in:
Scripta Materialia
Received Date: Revised Date: Accepted Date:
13 August 2013 30 August 2013 30 August 2013
Please cite this article as: H. Na, S. Nambu, M. Ojima, J. Inoue, T. Koseki, Strain localization behavior in lowcarbon martensitic steel during tensile deformation, Scripta Materialia (2013), doi: http://dx.doi.org/10.1016/ j.scriptamat.2013.08.030
This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Strain localization behavior in low-carbon martensitic steel during tensile deformation Hyuntaek Na, S. Nambu, M. Ojima, J. Inoue, and T. Koseki
Department of Materials Engineering, Graduate School of Engineering, The University of Tokyo, 7-3-1 Hongo, Bunkyo-ku, Tokyo 113-8656, Japan
Corresponding author: Tel. / Fax: +81-3-5841-7111; E-mail: [email protected]
Abstract The microscopic strain distribution of lath martensitic steel during tensile deformation up to a strain of 10% has been measured in situ. Strain localization, which indicates the grain interaction, is clearly observed in the vicinity of subblock boundaries, and it affects the inhomogeneous crystal rotation behavior within block (or subblock), and hence resultant rapid grain subdivision of the lath martensitic steel during elongation.
Keywords: Martensitic steels; Plastic deformation; Grain refining; Digital image correlation; Electron backscattering diffraction (EBSD)
Grain refinement in steel is a unique method of achieving a high strength and high toughness simultaneously [1-3]. Recently, Tsuji et al. have reported that ultrafine-grained steel can be easily fabricated using a low-carbon martensite steel (Fe-0.13 wt.% C) as a starting material by a conventional cold rolling process to reduce thickness by only 50% and a subsequent heat treatment at approximately 723 K [2]. They suggested that the deformation
during the cold rolling process results in the formation of an ultrafine-grained structure owing to the inhomogeneous deformation associated with its complex starting microstructures, such as packets, blocks, subblocks, and laths. In fact, Shanthraj and Zikry suggested from a numerical study that complex interactions between dislocations and crystallographic boundaries, such as lath and block boundaries, lead to an inhomogeneous deformation behavior in the vicinity of those interfaces during deformation [4]. In contrast, Ohmura et al. demonstrated by in situ observation during nanoindentation inside a transmission electron microscope (TEM) that the effect of interaction in a lath boundary is minor [5]. In our previous study [6], we demonstrated using an in situ tensile experiment in a field emission scanning electron microscope (FESEM) equipped with an electron backscattered diffraction (EBSD) analyzer that the inhomogeneity in the activation of the slip system within a single block actually starts to develop with a relatively small strain and causes the rapid grain refinement. However, it was not clarified which boundary was responsible for the inhomogeneity. In our present study, we aim to clarify the grain interaction behavior in lath martensite by investigating the changes in the strain distribution and activation behavior of the slip system using the digital image correlation (DIC) method [7] and EBSD analysis coupled with a crystal plasticity model, respectively. A multilayered structure [8-11] was employed as a tool to acquire sufficient plastic deformation in martensitic steel under uniform tensile loading. The employed multilayered structure is composed of two 0.2 wt.% C martensitic steel layers and one Type 316L austenitic steel layer. First, the 20 mm-thick Type 316L steel layer was sandwiched by the two 4 mm-thick 0.2 wt.% C steel layers. Then, the thickness of the as-sandwiched steel composite was reduced by hot rolling and subsequent cold rolling, so that the final thickness of the 0.2 wt.% C steel layer was approximately 130 µm. The chemical compositions of the components are shown in Table 1. The specimens were austenized at 1373 K for 1200 s and
subsequently quenched by Ar and 3% H2 gas to obtain a full lath martnesitic structure. The surface of the specimen was mechanically polished, and then lightly etched using nital (2% solution of nitric acid in ethanol) to reveal a high contrast microstructural pattern, which enables the effective tracking of the movement of the specimen surface by the DIC method during the subsequent in situ tensile experiments. The multilayered steel plate (30 mm × 5 mm × 1 mm) was uniformly elongated using a micro-tensile tester with an engineering strain rate of 1.0 × 10-3 s-1. During tensile test, the microstructure on the surface of the specimen was filmed using the FESEM (JSM7001F, JEOL). The distance between two indentation markers was measured on the surface of the martensitic steel layer, such that macroscopic tensile strains of 0, 3, 6, and 10% were applied in a stepwise manner to a single sample. Crystallographic features were analyzed using the EBSD analyzer (TexSEM Laboratory (TSL)) attached to the FESEM, with acceleration voltage 15 kV, beam spot diameter 20 nm and step size 0.5 µm. Scanning areas were approximately 150 × 200 µm2, 180 × 200 µm2, 200 × 250 µm2, and 250 × 300 µm2 for the cases of 0, 3, 6, and 10% elongations, respectively. The rolling (RD), transverse (TD), and normal (ND) directions of the specimen were used as the standard coordinate system in the crystallographic orientation analysis. The tensile direction corresponds to RD in this study. Accordingly, all the inverse pole figures and color coding in the inverse pole figure maps in this study plot the crystal orientation parallel to the RD. Strain fields were derived using the DIC software based on Image J [12] developed in house on the basis of a classical algorithm [13], by keeping the average distance between the points of interest at 0.5 µm. Figures 1(a) and 1(b) show the microstructure of the martensitic steel before and after 10% elongation, respectively, as revealed by EBSD analysis. The black lines in Fig. 1 indicate the boundaries where the misorientation between adjacent points is more than 5˚, and it is demonstrated that, after 10% elongation, many new boundaries with misorientations of
more than 5˚ are formed. Figure 2 shows the change in distribution of the maximum shear strain in the martensitic steel (highlighted area in Fig 1) during deformation obtained by DIC analysis. The grain structure derived by EBSD analysis is superimposed in the figure, and the black lines indicate boundaries with misorientations of more than 15˚. Increment in maximum shear strain for macroscopic tensile strain intervals of 0%-3%, 3%-6%, and 6%10% are shown in Figs 2(a), 2(b), and 2(c), respectively. As shown in Fig. 2(a), the deformation behavior of the martensitic steel is inhomogeneous even at a strain less than 3%. From 3% to 6% elongation (see Fig. 2(b)), a significant strain localization is found in some of the blocks. Afterward, the intensity of strain accumulation in the strain localized area steadily increases up to 10% elongation, as shown in Fig. 2(c). Increment in maximum shear strain for macroscopic tensile strain intervals of 6%-10% and the grain boundary map before deformation within a selected block of the martensitic steel are presented in Figs. 3(a) and 3(b), respectively. The profiles of increment of maximum shear strain along line A (depicted in Fig. 3(a)) and the misorientation angles along line B (depicted in Fig. 3(b)) are shown in Figs. 3(c) and 3(d), respectively. A total of four subblock boundaries, as indicated in Fig. 3(d) by B1, B2, B3, and B4, are found within the block. The misorientation angles at the subblock boundaries are respectively 4.9˚, 8.2˚, 7.4˚, and 5.6˚, which are in good agreement with those obtained in the previous study by Kitahara et al. [14]. As shown in Fig. 3(c), the peak positions of the increment of maximum shear strain indicated as P1, P2, P3, and P4 correspond well to the positions of the subblock boundaries. This result indicates that a subblock boundary works as a barrier that interacts with mobile dislocation during deformation [15]. Noticeable strain localization at the P4 that corresponds to the B4 subblock boundary seems to be attributed to grain interaction in the vicinity of not only the B4 subblock boundary but also pre-existed (such as inner microstructure, block, and packet etc.) and newly generated boundaries. Figure 4(a) shows the series of grain boundary distribution
within a selected block during elongation. New boundaries with misorientations ranging from 5˚ to 10˚ (low angle boundary #1, LAB #1), from 10˚ to 15˚ (LAB #2), and above 15˚ (high angle boundary, HAB) are presented in red, blue, and black, respectively. After 3% elongation, subblock boundaries can be clearly observed in some areas (see highlighted area with red in Fig. 4(a)). This indicates that crystal rotation behavior of each subblock varies owing to the difference in initial crystal orientation (variant). It is also confirmed that new grain boundaries were formed adjacent to the subblock boundaries after 10% elongation. New boundaries are observed around the aforementioned subblock boundaries as well as block boundaries, most of which fall to LAB #1 and LAB #2. The boundaries with misorientations of more than 5˚ after 10% elongation can be classified into three types, block boundary (BB), subblock boundary (SBB), and newly generated boundary (NGB)), the spatial distributions of which are presented in black, red, and green lines, respectively, in Fig. 4(b) around one of the subblock boundaries. Figure 4(c) shows the profile of the misorientation angle along the line parallel to the subblock boundary where a severe strain accumulation (indicated by a gray line in Fig. 4(b)) was observed after 10% elongation. Interestingly enough, the misorientation from the reference points is almost constant except for the discontinuity found at the NGB. As shown in Fig. 4(d), even though the crystal orientations at all observation points basically move to the <110> pole, there exists a clear difference among them, which is natural because the grain interaction differs from point to point. It should be noted, however, that the crystal rotation behaviors along the gray line clearly fall into three definite groups, which correspond to the regions depicted as G1, G2, and G3 in Fig. 4(c). The slip systems activated during the 10% elongation in, G1, G2, and G3, were identified and are shown in Table 2. An algorithm of the identification of the activated slip systems is described in detail in our previous publication [11]. The activation behavior of slip systems was basically found to follow the lath constraint mechanism discussed in the
previous studies [10, 11]. That is, only one or two in-lath type slip systems were activated. That limitation of activated slip systems presumably leads to a stepwise shift in crystal rotation behavior along the subblock boundary, instead of a smooth gradual shift. In summary, in contrast to the case of ferritic steel, crystallographic refinement becomes significant in 0.2 wt.% C martensitic steel with only approximately 10% elongation. From the in situ measurement of the plastic strain distribution using the DIC technique, a clear indication of grain interaction is evidenced in the vicinity of subblock boundaries during deformation, which is sufficiently effective for inducing an inhomogeneous crystal rotation behavior within a block (or subblock). Along the subblock boundary, a stepwise shift in crystal rotation trajectory is observed. The limitation of activated slip systems (the lath constraint mechanism) is also found to be associated with the unique local deformation behavior of lath martensitic steel.
The present research was conducted as part of the ISIJ Research Promotion Grant and ALCA project funded by the Iron and Steel Institute of Japan (ISIJ) and Japan Science and Technology Agency, respectively.
References [1] R.Z. Valiev, Mater. Sci. Eng. A, 234 (1997) 59. [2] N. Tsuji, R. Ueji, Y. Minamino, Y. Saito, Scripta Mater. 46 (2002) 305. [3] J.K. Choi, D.H. Seo, J.S. Lee, K.K Um, W.Y. Choo, ISIJ Int. 43 (2003) 764. [4] P. Shanthraj and M.A. Zikry, J. Mater. Res. 27 (2012) 1598. [5] T. Ohmura, A.M. Minor, E.A. Stach, J.W. Morris, Jr., J. Mater. Res. 19 (2004) 3626. [6] H. Na, S. Nambu, M. Ojima, J. Inoue, T. Koseki, submitted for publication. [7] M.A. Sutton, W.J. Wolters, W.H. Peters, W.F. Ranson, S.R. McNeill, Image Vision Comput. 1 (1983) 133. [8] J. Inoue, S. Nambu, Y. Ishimoto, T. Koseki, Scripta Mater. 59 (2008) 1055. [9] S. Nambu, M. Michiuchi, J. Inoue, T. Koseki, Comp. Sci. Tech. 69 (2009) 1936. [10] S. Nambu, M. Michiuchi, Y. Ishimoto, K. Asakura, J. Inoue, T. Koseki, Scripta Mater. 60 (2009) 221. [11] M. Michiuchi, S. Nambu, Y. Ishimoto, J. Inoue, T. Koseki, Acta Mater. 57 (2009) 5283. [12] M.D. Abramoff, P.J. Magelhaes, S.J. Ram, Biophotonics Int. 11 (2004) 36. [13] S. Bergonnier, F. Hild, S. Roux, J. Strain Anal. Eng. 40 (2005) 185. [14] H. Kitahara, R. Ueji, N. Tsuji, Y. Minamino, Acta Mater. 54 (2006) 1279. [15] M.R. Stoudta, L.E. Levinea, A. Creuzigera, J.B. Hubbarda, Mater. Sci. Eng. A 530 (2011) 107.
Figure captions Fig. 1 Crystal orientation maps of 0.2 wt.% C martensitic steel before (a) and after 10% elongation (b).
Fig. 2 Series of increment in maximum shear strain for macroscopic tensile strain intervals of (a) from 0% to 3%, (b) from 3% to 6%, and (c) from 6% to 10% in 0.2 wt.% C martensitic steel.
Fig. 3 Maximum shear strain distribution map (increment for macroscopic tensile strain interval of 6%-10%) (a), grain boundary(map (before elongation) (b) of selected block, plastic strain profile along line A (c), and point-to-point misorientation profile along line B (d).
Fig. 4 Deformation of selected block during the in-situ observation (a), schematic illustration of boundary distribution in local area within the selected block after 10% elongation (b), misorientation profile (point-to-origin) in the vicinity of typical subblock boundary (c), and standard stereographic projection of changes in tensile directions of observation points after 10% elongation (d).
Fig. 1
Fig. 2
Fig. 3
Fig. 4
Table 1 Chemical compositions of constituent materials (wt. %) Type of steel C Si Mn Ni Martensitic steel layer 0.2 0.25 0.25 14 Austenitic steel layer 0.02 0.63 0.84 12.09
Cr 17.76
Mo 2.12
Fe Bal. Bal.
Table 2 Activated slip systems at observation points during 10% elongation (four slip systems with highest Schmid factor among twenty-four slip systems are shown, and activated slip systems are highlighted in gray) (SP, SD, and SF are slip plane, slip direction, and Schmid factor, respectively) P1 P2 P3 SP SD SF SP SD SF SP SD SF (110) B2 0.498 (110) B2 0.494 (110) B2 0.496 (110) B4 0.469 (112) B4 0.483 (110) B4 0.468 (112) B2 0.447 (110) B2 0.435 (112) B2 0.451 (112) B4 0.415 (112) B4 0.420 (112) B4 0.407
S1359-6462(13)00435-1 http://dx.doi.org/10.1016/j.scriptamat.2013.08.030 SMM 10040
To appear in:
Scripta Materialia
Received Date: Revised Date: Accepted Date:
13 August 2013 30 August 2013 30 August 2013
Please cite this article as: H. Na, S. Nambu, M. Ojima, J. Inoue, T. Koseki, Strain localization behavior in lowcarbon martensitic steel during tensile deformation, Scripta Materialia (2013), doi: http://dx.doi.org/10.1016/ j.scriptamat.2013.08.030
This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Strain localization behavior in low-carbon martensitic steel during tensile deformation Hyuntaek Na, S. Nambu, M. Ojima, J. Inoue, and T. Koseki
Department of Materials Engineering, Graduate School of Engineering, The University of Tokyo, 7-3-1 Hongo, Bunkyo-ku, Tokyo 113-8656, Japan
Corresponding author: Tel. / Fax: +81-3-5841-7111; E-mail: [email protected]
Abstract The microscopic strain distribution of lath martensitic steel during tensile deformation up to a strain of 10% has been measured in situ. Strain localization, which indicates the grain interaction, is clearly observed in the vicinity of subblock boundaries, and it affects the inhomogeneous crystal rotation behavior within block (or subblock), and hence resultant rapid grain subdivision of the lath martensitic steel during elongation.
Keywords: Martensitic steels; Plastic deformation; Grain refining; Digital image correlation; Electron backscattering diffraction (EBSD)
Grain refinement in steel is a unique method of achieving a high strength and high toughness simultaneously [1-3]. Recently, Tsuji et al. have reported that ultrafine-grained steel can be easily fabricated using a low-carbon martensite steel (Fe-0.13 wt.% C) as a starting material by a conventional cold rolling process to reduce thickness by only 50% and a subsequent heat treatment at approximately 723 K [2]. They suggested that the deformation
during the cold rolling process results in the formation of an ultrafine-grained structure owing to the inhomogeneous deformation associated with its complex starting microstructures, such as packets, blocks, subblocks, and laths. In fact, Shanthraj and Zikry suggested from a numerical study that complex interactions between dislocations and crystallographic boundaries, such as lath and block boundaries, lead to an inhomogeneous deformation behavior in the vicinity of those interfaces during deformation [4]. In contrast, Ohmura et al. demonstrated by in situ observation during nanoindentation inside a transmission electron microscope (TEM) that the effect of interaction in a lath boundary is minor [5]. In our previous study [6], we demonstrated using an in situ tensile experiment in a field emission scanning electron microscope (FESEM) equipped with an electron backscattered diffraction (EBSD) analyzer that the inhomogeneity in the activation of the slip system within a single block actually starts to develop with a relatively small strain and causes the rapid grain refinement. However, it was not clarified which boundary was responsible for the inhomogeneity. In our present study, we aim to clarify the grain interaction behavior in lath martensite by investigating the changes in the strain distribution and activation behavior of the slip system using the digital image correlation (DIC) method [7] and EBSD analysis coupled with a crystal plasticity model, respectively. A multilayered structure [8-11] was employed as a tool to acquire sufficient plastic deformation in martensitic steel under uniform tensile loading. The employed multilayered structure is composed of two 0.2 wt.% C martensitic steel layers and one Type 316L austenitic steel layer. First, the 20 mm-thick Type 316L steel layer was sandwiched by the two 4 mm-thick 0.2 wt.% C steel layers. Then, the thickness of the as-sandwiched steel composite was reduced by hot rolling and subsequent cold rolling, so that the final thickness of the 0.2 wt.% C steel layer was approximately 130 µm. The chemical compositions of the components are shown in Table 1. The specimens were austenized at 1373 K for 1200 s and
subsequently quenched by Ar and 3% H2 gas to obtain a full lath martnesitic structure. The surface of the specimen was mechanically polished, and then lightly etched using nital (2% solution of nitric acid in ethanol) to reveal a high contrast microstructural pattern, which enables the effective tracking of the movement of the specimen surface by the DIC method during the subsequent in situ tensile experiments. The multilayered steel plate (30 mm × 5 mm × 1 mm) was uniformly elongated using a micro-tensile tester with an engineering strain rate of 1.0 × 10-3 s-1. During tensile test, the microstructure on the surface of the specimen was filmed using the FESEM (JSM7001F, JEOL). The distance between two indentation markers was measured on the surface of the martensitic steel layer, such that macroscopic tensile strains of 0, 3, 6, and 10% were applied in a stepwise manner to a single sample. Crystallographic features were analyzed using the EBSD analyzer (TexSEM Laboratory (TSL)) attached to the FESEM, with acceleration voltage 15 kV, beam spot diameter 20 nm and step size 0.5 µm. Scanning areas were approximately 150 × 200 µm2, 180 × 200 µm2, 200 × 250 µm2, and 250 × 300 µm2 for the cases of 0, 3, 6, and 10% elongations, respectively. The rolling (RD), transverse (TD), and normal (ND) directions of the specimen were used as the standard coordinate system in the crystallographic orientation analysis. The tensile direction corresponds to RD in this study. Accordingly, all the inverse pole figures and color coding in the inverse pole figure maps in this study plot the crystal orientation parallel to the RD. Strain fields were derived using the DIC software based on Image J [12] developed in house on the basis of a classical algorithm [13], by keeping the average distance between the points of interest at 0.5 µm. Figures 1(a) and 1(b) show the microstructure of the martensitic steel before and after 10% elongation, respectively, as revealed by EBSD analysis. The black lines in Fig. 1 indicate the boundaries where the misorientation between adjacent points is more than 5˚, and it is demonstrated that, after 10% elongation, many new boundaries with misorientations of
more than 5˚ are formed. Figure 2 shows the change in distribution of the maximum shear strain in the martensitic steel (highlighted area in Fig 1) during deformation obtained by DIC analysis. The grain structure derived by EBSD analysis is superimposed in the figure, and the black lines indicate boundaries with misorientations of more than 15˚. Increment in maximum shear strain for macroscopic tensile strain intervals of 0%-3%, 3%-6%, and 6%10% are shown in Figs 2(a), 2(b), and 2(c), respectively. As shown in Fig. 2(a), the deformation behavior of the martensitic steel is inhomogeneous even at a strain less than 3%. From 3% to 6% elongation (see Fig. 2(b)), a significant strain localization is found in some of the blocks. Afterward, the intensity of strain accumulation in the strain localized area steadily increases up to 10% elongation, as shown in Fig. 2(c). Increment in maximum shear strain for macroscopic tensile strain intervals of 6%-10% and the grain boundary map before deformation within a selected block of the martensitic steel are presented in Figs. 3(a) and 3(b), respectively. The profiles of increment of maximum shear strain along line A (depicted in Fig. 3(a)) and the misorientation angles along line B (depicted in Fig. 3(b)) are shown in Figs. 3(c) and 3(d), respectively. A total of four subblock boundaries, as indicated in Fig. 3(d) by B1, B2, B3, and B4, are found within the block. The misorientation angles at the subblock boundaries are respectively 4.9˚, 8.2˚, 7.4˚, and 5.6˚, which are in good agreement with those obtained in the previous study by Kitahara et al. [14]. As shown in Fig. 3(c), the peak positions of the increment of maximum shear strain indicated as P1, P2, P3, and P4 correspond well to the positions of the subblock boundaries. This result indicates that a subblock boundary works as a barrier that interacts with mobile dislocation during deformation [15]. Noticeable strain localization at the P4 that corresponds to the B4 subblock boundary seems to be attributed to grain interaction in the vicinity of not only the B4 subblock boundary but also pre-existed (such as inner microstructure, block, and packet etc.) and newly generated boundaries. Figure 4(a) shows the series of grain boundary distribution
within a selected block during elongation. New boundaries with misorientations ranging from 5˚ to 10˚ (low angle boundary #1, LAB #1), from 10˚ to 15˚ (LAB #2), and above 15˚ (high angle boundary, HAB) are presented in red, blue, and black, respectively. After 3% elongation, subblock boundaries can be clearly observed in some areas (see highlighted area with red in Fig. 4(a)). This indicates that crystal rotation behavior of each subblock varies owing to the difference in initial crystal orientation (variant). It is also confirmed that new grain boundaries were formed adjacent to the subblock boundaries after 10% elongation. New boundaries are observed around the aforementioned subblock boundaries as well as block boundaries, most of which fall to LAB #1 and LAB #2. The boundaries with misorientations of more than 5˚ after 10% elongation can be classified into three types, block boundary (BB), subblock boundary (SBB), and newly generated boundary (NGB)), the spatial distributions of which are presented in black, red, and green lines, respectively, in Fig. 4(b) around one of the subblock boundaries. Figure 4(c) shows the profile of the misorientation angle along the line parallel to the subblock boundary where a severe strain accumulation (indicated by a gray line in Fig. 4(b)) was observed after 10% elongation. Interestingly enough, the misorientation from the reference points is almost constant except for the discontinuity found at the NGB. As shown in Fig. 4(d), even though the crystal orientations at all observation points basically move to the <110> pole, there exists a clear difference among them, which is natural because the grain interaction differs from point to point. It should be noted, however, that the crystal rotation behaviors along the gray line clearly fall into three definite groups, which correspond to the regions depicted as G1, G2, and G3 in Fig. 4(c). The slip systems activated during the 10% elongation in, G1, G2, and G3, were identified and are shown in Table 2. An algorithm of the identification of the activated slip systems is described in detail in our previous publication [11]. The activation behavior of slip systems was basically found to follow the lath constraint mechanism discussed in the
previous studies [10, 11]. That is, only one or two in-lath type slip systems were activated. That limitation of activated slip systems presumably leads to a stepwise shift in crystal rotation behavior along the subblock boundary, instead of a smooth gradual shift. In summary, in contrast to the case of ferritic steel, crystallographic refinement becomes significant in 0.2 wt.% C martensitic steel with only approximately 10% elongation. From the in situ measurement of the plastic strain distribution using the DIC technique, a clear indication of grain interaction is evidenced in the vicinity of subblock boundaries during deformation, which is sufficiently effective for inducing an inhomogeneous crystal rotation behavior within a block (or subblock). Along the subblock boundary, a stepwise shift in crystal rotation trajectory is observed. The limitation of activated slip systems (the lath constraint mechanism) is also found to be associated with the unique local deformation behavior of lath martensitic steel.
The present research was conducted as part of the ISIJ Research Promotion Grant and ALCA project funded by the Iron and Steel Institute of Japan (ISIJ) and Japan Science and Technology Agency, respectively.
References [1] R.Z. Valiev, Mater. Sci. Eng. A, 234 (1997) 59. [2] N. Tsuji, R. Ueji, Y. Minamino, Y. Saito, Scripta Mater. 46 (2002) 305. [3] J.K. Choi, D.H. Seo, J.S. Lee, K.K Um, W.Y. Choo, ISIJ Int. 43 (2003) 764. [4] P. Shanthraj and M.A. Zikry, J. Mater. Res. 27 (2012) 1598. [5] T. Ohmura, A.M. Minor, E.A. Stach, J.W. Morris, Jr., J. Mater. Res. 19 (2004) 3626. [6] H. Na, S. Nambu, M. Ojima, J. Inoue, T. Koseki, submitted for publication. [7] M.A. Sutton, W.J. Wolters, W.H. Peters, W.F. Ranson, S.R. McNeill, Image Vision Comput. 1 (1983) 133. [8] J. Inoue, S. Nambu, Y. Ishimoto, T. Koseki, Scripta Mater. 59 (2008) 1055. [9] S. Nambu, M. Michiuchi, J. Inoue, T. Koseki, Comp. Sci. Tech. 69 (2009) 1936. [10] S. Nambu, M. Michiuchi, Y. Ishimoto, K. Asakura, J. Inoue, T. Koseki, Scripta Mater. 60 (2009) 221. [11] M. Michiuchi, S. Nambu, Y. Ishimoto, J. Inoue, T. Koseki, Acta Mater. 57 (2009) 5283. [12] M.D. Abramoff, P.J. Magelhaes, S.J. Ram, Biophotonics Int. 11 (2004) 36. [13] S. Bergonnier, F. Hild, S. Roux, J. Strain Anal. Eng. 40 (2005) 185. [14] H. Kitahara, R. Ueji, N. Tsuji, Y. Minamino, Acta Mater. 54 (2006) 1279. [15] M.R. Stoudta, L.E. Levinea, A. Creuzigera, J.B. Hubbarda, Mater. Sci. Eng. A 530 (2011) 107.
Figure captions Fig. 1 Crystal orientation maps of 0.2 wt.% C martensitic steel before (a) and after 10% elongation (b).
Fig. 2 Series of increment in maximum shear strain for macroscopic tensile strain intervals of (a) from 0% to 3%, (b) from 3% to 6%, and (c) from 6% to 10% in 0.2 wt.% C martensitic steel.
Fig. 3 Maximum shear strain distribution map (increment for macroscopic tensile strain interval of 6%-10%) (a), grain boundary(map (before elongation) (b) of selected block, plastic strain profile along line A (c), and point-to-point misorientation profile along line B (d).
Fig. 4 Deformation of selected block during the in-situ observation (a), schematic illustration of boundary distribution in local area within the selected block after 10% elongation (b), misorientation profile (point-to-origin) in the vicinity of typical subblock boundary (c), and standard stereographic projection of changes in tensile directions of observation points after 10% elongation (d).
Fig. 1
Fig. 2
Fig. 3
Fig. 4
Table 1 Chemical compositions of constituent materials (wt. %) Type of steel C Si Mn Ni Martensitic steel layer 0.2 0.25 0.25 14 Austenitic steel layer 0.02 0.63 0.84 12.09
Cr 17.76
Mo 2.12
Fe Bal. Bal.
Table 2 Activated slip systems at observation points during 10% elongation (four slip systems with highest Schmid factor among twenty-four slip systems are shown, and activated slip systems are highlighted in gray) (SP, SD, and SF are slip plane, slip direction, and Schmid factor, respectively) P1 P2 P3 SP SD SF SP SD SF SP SD SF (110) B2 0.498 (110) B2 0.494 (110) B2 0.496 (110) B4 0.469 (112) B4 0.483 (110) B4 0.468 (112) B2 0.447 (110) B2 0.435 (112) B2 0.451 (112) B4 0.415 (112) B4 0.420 (112) B4 0.407