Indentation-induced deformation behavior in martensitic steel observed through in-situ nanoindentation in a transmission electron microscopy

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1 Materials Science Forum Vols (2006) pp online at (2006) Trans Tech Publications, Switzerland No. 114 Indentation-induced deformation behavior in martensitic steel observed through in-situ nanoindentation in a transmission electron microscopy T. Ohmura 1,a, A. Minor 2,b, K. Tsuzaki 1,c, J.W. Morris, Jr. 3,d 1 Steel Research Center, National Institute for Materials Science, Tsukuba, Japan 2 National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, California USA 3 Department of Materials Science and Engineering, University of California, Berkeley, California USA a OHMURA.Takahito@nims.go.jp, b AMinor@lbl.gov c TSUZAKI.Kaneaki@nims.go.jp, d jwmorris@uclink4.berkeley.edu Keywords: transmission electron microscopy, in-situ nanoindentation, Fe-C martensite, grain boundary, dislocation Abstract. Deformation behavior in the vicinity of grain boundary in Fe-0.4wt%C tempered martensitic steel were studied through in-situ nanoindentation in a TEM. Two types of boundaries were imaged in the dislocated martensitic structure: a low-angle lath boundary and a high-angle block boundary. In the case of a low-angle grain boundary, the dislocations induced by the indenter piled up against the boundary. As the indenter penetrated further, a critical stress appears to have been reached and a high density of dislocations was suddenly emitted on the far side of the grain boundary into the adjacent grain. In the case of the high-angle grain boundary, the numerous dislocations that were produced by the indentation were simply absorbed into the boundary, with no indication of pile-up or the transmission of strain. Introduction The interaction of dislocations with grain boundaries is of fundamental importance to engineering the strength of polycrystalline materials especially for ultra-fine grain metals produced by severe plastic deformation. The flow stress of a material increases as the density of grain boundaries increases, usually in accordance with the well-known Hall-Petch relation. A variety of models have been proposed to interpret this relation 1-5, and a number of experimental investigations have studied dislocation behavior in the vicinity of grain boundaries However, in situ straining TEM studies of dislocation activity at grain boundaries have observed individual dislocations to pile up or pass through boundaries 9,10. These studies show interactions between pre-existing lattice dislocations and grain boundaries, and hence relate only to the initial stages of plastic deformation. The flow stress after yielding is strongly affected by dislocation interactions under conditions of high dislocation density, and deformation behavior in this regime must also be understood. Fe-C based martensitic steels are among the most important high-strength materials with fine microstructure. Their structures are often dislocated lath martensite, which has a high dislocation density in the as-quenched condition. The microstructure has a characteristic form that includes at least three distinct kinds of structural units: lath, block or packet, and prior austenite grain 11,12. The prior austenite grain consists of some packets or blocks, and they are divided into the lath structure. Only the lath boundaries are low angle. While Ohmura et al have shown through conventional nanoindentation techniques that grain boundary strengthening is a significant factor in martensitic steels, the relative roles of the three kinds of boundaries remain unclear. Morris and co-workers 17,18 have suggested on theoretical grounds that the prior austenite grain boundaries should dominate, while the lath and packet boundaries should be relatively unimportant. The low-angle lath boundaries are not expected to pose more than nominal barriers to dislocation All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of the publisher: Trans Tech Publications Ltd, Switzerland, (ID: /10/06,22:37:54)

2 240 Nanomaterials by Severe Plastic Deformation motion while the high-angle block boundaries allow continuity of five of the six {110} glide planes of the bcc martensite with only small-angle deflections. The present study was undertaken to use the newly available technique of in situ nanoindentation in TEM to study the interactions between the dense dislocation distribution in Fe-C lath martensitic steel with the different internal boundaries. Experimental The sample used in this study was a high-purity Fe-C binary martensite with a carbon content of 0.4 wt%. The specimen was austenitized at 1323K for 900s, ice-brine quenched and tempered at 723K for 5400s. Following heat treatment the specimen surface was mechanically polished and then electropolished. A FEI Strata 235 dual-beam focused ion beam (FIB) apparatus was then used to thin the specimen so that it would be transparent to the electron beam in a geometry suitable for in situ nanoindentation in a unique TEM nanoindentation stage that is described elsewhere 19,20. The electropolished surface was coated with two Pt films before FIB milling in order to avoid damage to the surface from the incident Ga + ions. The Pt coating was first deposited with an electron beam, and the covered by ion beam-deposited Pt. The nanoindenter tip was maneuvered in the microscope to crack the brittle Pt films, and the Pt was then peeled off of the surface with the indenter to reveal a clean electropolished surface. Indentations were made into the surface such that a specific low angle or high angle boundary would be visible beneath the indenter, and the deformation at that boundary was followed in real time. Results and Discussion TEM micrographs of the Fe-C martensite before nanoindentation are shown in Fig. 1. Two suitably oriented boundaries were found and indented. The two boundaries are highlighted with arrowheads in Fig. 1. The first boundary chosen was the low-angle boundary shown in Fig. 1(a). The selected-area diffraction pattern (SADP) included in the figure shows that the grains on either side of the boundary have the same [111] α zone axis, indicating that the grain boundary is low-angle. Both its low-angle character and its visual appearance strongly suggest that this boundary is a lath boundary 21,22. The second boundary chosen was the high-angle boundary shown in Fig. 1(b). In this case, the SADP from the bounding grain nearest the surface shows a [100] α zone axis pattern while that from the adjacent grain has a [111] α zone axis pattern. Using three-dimensional orientation analysis software (TexSEM Laboratories, Inc., TOCA), the misorientation across the grain boundary was determined to be ~ 56º, with a rotation axis near [0.16, 0.64, 0.76]. Morito et al. 23 have calculated the 23 possible combinations of misorientation angle and rotation axis among the 24 variants of the Kurdjumov-Sachs (K-S) relationship that results from the martensitic phase transformation ({111} γ {110} α, [110] γ [111] α ). These are the possible misorientations across block boundaries in martensite. The misorientation angle and the rotation axis of the boundary in Fig. 1(b) are close to those of one of those possible block boundaries, which has a (a) Fig. 1 TEM micrographs of the Fe-C martensite before nanoindentation including (a) low-angle grain boundary and (b) high-angle grain boundary. (b)

3 Materials Science Forum Vols misorientation angle of and a rotation axis of (a) [-0.246, 0.628, 0], which strongly suggests that this high-angle boundary is a block boundary. An alternate, perhaps simpler picture is given by Guo and Morris 18,24 ; before the shear that is necessary to complete the transformation the two K-S variants that share the same {111} γ plane and <110> γ direction are rotated by 60º about a <110> α direction, in close agreement with the experimental measurement. This suggest that the boundary under study is a block boundary between K-S (b) variants with a shared {110} α close-packed plane in the boundary. Fig. 2 represents a series of video frames from the in situ TEM nanoindentation of the low-angle boundary. Fig. 2(a) shows the diamond indenter prior to indentation, as it approached the surface from the upper-right corner of the figure. The grain boundary is indicated by arrow-heads. As the indenter penetrated into the grain (Fig. 2(b)), the deformation was (c) accommodated by the motion of dislocations away from the indenter contact point toward the grain boundary. The low-angle boundary offered some resistance to the dislocation motion, resulting in a pile-up at the grain boundary that is shown in Fig. 2(c) (a frame taken at an indenter penetration depth of 46 nm). As the indenter penetrated further, a large number of dislocations were emitted on the far side from the indenter tip (left and lower side on the micrograph) of the grain boundary into (d) the adjacent grain. A dense and tangled dislocation structure is seen on the far side of the grain boundary in Fig. 2(d), which was taken at an 84 nm penetration depth. When the indenter was withdrawn a high density of dislocations was retained near the far side of the boundary. The dislocation interactions with the high-angle grain boundary (block boundary) were dramatically different, as documented in Fig. 3. Fig. 3(a) shows the Fig. 2 In situ TEM micrographs of the initial state before indentation, with the high-angle grain Fe-C martensite, including low-angle boundary: (a) before indentation, (b) at boundary indicated by arrows. During the early stages 21 nm, (c) at 46 nm and (d) at 84 nm of the indentation dislocations can be seen to sweep penetration depth. across the grain from right to left, shown part-way in Fig. 3(b). Note that the dislocations in the left area of the indented grain are not seen in the Fig. 3 (b) any more, while they are clearly shown in the Fig. 3 (a). This could be due to a change of diffraction condition into an invisible condition of the dislocations because of the deformation of the sample. At approximately 60 nm penetration depth, a dense cloud of dislocations reached the high-angle boundary. However, as shown in Figs. 3(c), on reaching the boundary the dislocations simply disappeared; there is no indication of a pile-up or significant penetration into the adjacent grain. In fact, as illustrated in Fig. 3(d), there is virtually no change in the dislocation configuration on the far side of the boundary. The grain beyond the boundary is essentially unaffected. Fig. 4 shows the configuration of the high angle and low angle grain boundaries after

4 242 Nanomaterials by Severe Plastic Deformation deformation, with the outline of the original grain boundary positions schematically drawn with broken lines. In the case of low-angle boundary shown in Fig. 4(a), significant plastic deformation occurs in the adjvacent grain as well as in the indented grain, resulting in a large shift in the location of the grain boundary. In the case of the high-angle boundary shown in Fig. 4(b), the indented grain was heavily deformed while the adjacent grain showed little change in shape. In particular, the grain boundary remained in its original position. Given that the indented grain in Fig. 4(b) clearly was deformed while the adjacent grain across the block boundary was not, and given that many dislocations were annihilated at the interface along with the plastic slip they carried, we must ask where all that deformation went. While further investigation is needed, and is underway, the simplest plausible answer is that almost all of the deformation was, ultimately, perpendicular to the field of view, and was accommodated by sliding on the block boundary. To elaborate on this mechanism, note that, as discussed above, the grains across the block boundary are K-S related to the same parent austenite, and appear to share a boundary that is a {110} plane in both crystals. Since {110} is the dominant slip plane in bcc iron, the adjacent grains can slip in any parallel direction at the boundary; hence grain boundary sliding is a fairly simple process. At the same time, the sample being indented is thin in the direction perpendicular to the foil and can deform perpendicular to the foil to accommodate the plastic impression made by the indenter. Prior work suggests that it is not particularly difficult to transmit strain across block or packet boundaries in dislocated martensite. Experiment shows that the Hall-Petch effect of refining block size is small 5,17,24, and theory shows that five of the six available {110} slip planes have no more than small deflections at Fig. 3 In situ TEM micrographs of the Fe-C martensite, including high-angle boundary: (a) before indentation, (b) at 70 nm, (c) at 83 nm and (d) 92 nm penetration depth. a block boundary 17,18,24. However, there is some resistance, and the present data suggests that the resistance is sufficient to deflect the primary deformation into the direction perpendicular to the foil, with the dislocations that approach the boundary being accommodated and annihilated there by grain boundary sliding on the common {110} slip plane in the boundary. It follows that the deformation behavior observed at the high angle boundary is not well described by conventional models. The conventional models state that a grain boundary acts as both a barrier against dislocation motion toward the grain boundary and as a source of new dislocations. The behavior of the indentation-induced dislocations at the high-angle boundary showed neither of these two effects, but instead showed only absorption of dislocations by the boundary. In the absence of dislocation pile-ups at a grain boundary, no stress concentration arises; hence no dislocation source is activated in the adjacent grain. This is the reason for that no (a) (b) (c) (d)

5 Materials Science Forum Vols specific plastic deformation occurs in the adjacent grain and, as shown in Fig. 4(b), the high-angle boundary retains its original position with respect to the rest of the microstructure. In the case of the low-angle lath boundary, the slip planes of the dislocations that approach the boundary are essentially continuous across it. The dislocations that approach the boundary are resisted by dislocation-dislocation interactions with the array of dislocation in the boundary, but pass through with relative ease. Since it is, apparently, easier to transmit strain across the boundary than to deflect it in the through-thickness direction by sliding the boundary, the adjacent grain joins in the deformation. (a) (b) Fig. 4 TEM micrographs of the Fe-C martensite after nanoindentation, with outlines of the original positions before indentation. (a) low-angle grain boundary and (b) high-angle grain boundary. Conclusion Deformation behavior in the vicinity of interface in Fe-0.4wt%C tempered martensitic steel were studied through in-situ nanoindentation in a TEM. Two types of boundaries were imaged in the dislocated martensitic structure: a low-angle lath boundary and a high-angle block boundary. In the case of a low-angle grain boundary, the dislocations induced by the indenter piled up against the boundary. As the indenter penetrated further, a critical stress appears to have been reached and a high density of dislocations was suddenly emitted on the far side of the grain boundary into the adjacent grain. Since slip planes are essentially continuous across the boundary, the resistance to dislocation glide is presumably due to dislocation-dislocation interactions. In the case of the high-angle grain boundary, the numerous dislocations that were produced by the indentation were simply absorbed into the boundary, with no indication of pile-up or the transmission of strain. The plastic strain is accommodated by deformation perpendicular to the foil. The apparent reason for this behavior is that the interface between the two K-S variants across the boundary is a {110} α slip plane that is common to both variants, with the consequence that slip in the boundary plane (grain boundary sliding) is relatively easy. Acnowledgement The work done at LBNL and the University of California was supported by the Director, Office of Energy Research, Office of Basic Energy Sciences, U. S. Department of Energy, under Contract No. DE-AC03-76SF References [1] E.O. Hall: The deformation and ageing of mild steel. Proc. R. Soc.B64, 747 (1951). [2] N.J. Petch: The cleavage strength of polycrystals. J. Iron Steel Inst.174, 25 (1953). [3] J.C.M. Li: Petch relation and grain boundary sources. Trans. AIME 227, 239 (1963). [4] M.F. Ashby: The deformation of plastically non-homogeneous materials. Phil. Mag.21, 399 (1970). [5] J.W. Morris, Jr.: in Proc. Int. Symposium on Ultrafine Grained Steels, edited by S. Takaki and T. Maki (Iron and Steel Inst., Tokyo, Japan, 2001), p.34. [6] J.J. Hauser, and B. Chalmers: The plastic deformation of bicrystals of f.c.c. metals. Acta

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