Uniaxial tensile plastic deformation of a bulk
APPLIED PHYSICS LETTERS 89, 101918 共2006兲
Uniaxial tensile plastic deformation of a bulk nanocrystalline alloy studied by a high-energy x-ray diffraction technique G. J. Fana兲 Department of Materials Science and Engineering, The University of Tennessee, Knoxville, Tennessee 37996
L. F. Fu Department of Chemical Engineering and Materials Science, University of California, Davis, California 95616
Y. D. Wang Department of Materials Science and Engineering, The University of Tennessee, Knoxville, Tennessee 37996 and School of Materials and Metallurgy, Northeastern University, Shenyang 110004, China
Y. Ren Experimental Facilities Division, Advanced Phonon Source, Argonne National Laboratory, Argonne, Illinois 60439
H. Choo, P. K. Liaw, and G. Y. Wang Department of Materials Science and Engineering, The University of Tennessee, Knoxville, Tennessee 37996
N. D. Browning National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, California 94720
共Received 27 June 2006; accepted 19 July 2006; published online 8 September 2006兲 By employing a high-energy x-ray diffraction technique, the authors report that uniaxial tensile plastic deformation induced the grain growth and texture development in a bulk nanocrystalline Ni–Fe alloy. The effects become more pronounced with increasing the plastic strain 共closer to the fracture surface兲. The texture development accompanying the grain rotation indicates that dislocation motion contributed to the observed plasticity in the nanocrystalline Ni–Fe alloy. The quantitative experimental data suggest that the dislocation storage was absent in the uniformly deforming region; whereas the dislocation storage was present in the necking region, where the grain growth was substantial. © 2006 American Institute of Physics. 关DOI: 10.1063/1.2348783兴 Nanocrystalline 共nc兲 materials with average grain sizes less than 100 nm exhibit unusual deformation mechanisms. Molecular dynamics simulations have suggested that the plasticity in the nc metals and alloys is mainly mediated by grain-boundary 共GB兲 activities,1–3 in contrast to the dislocation pileups at the GBs in their coarse-grained polycrystalline counterparts.4 However, experimental verification of deformation mechanisms in nc metals and alloys is complicated by various competing mechanisms.5–14 A complete understanding of these competing mechanisms poses a great challenge experimentally, particularly in the threedimensional 共3D兲 bulk nc specimens. Previous investigations of deformation mechanisms of nc metals and alloys were primarily based on the in situ straining transmission electron microscopy 共TEM兲 observations in the two-dimensional thin foils,5,7,10,13,14 where artifacts often exist. The study of deformation mechanisms in a 3D bulk nc specimen was rare in the literature due to the limitations of the available experimental techniques, i.e., TEM, and to the intrinsic brittleness often observed in the bulk nc specimens. A good tensile ductility was recently realized in the bulk nc metals and alloys prepared by severe plastic deformation techniques 共i.e., mechanical milling and cold rolling兲.14–17 In the present investigation, the bulk nc a兲
Electronic mail: [email protected]
Ni-18 wt % Fe 共weight: wt %兲 sheets, which exhibit a good combination of the very high strength and a large tensile plastic strain, were produced using a pulsed electrodeposition technique. A high-energy x-ray diffraction 共HEXRD兲 technique was employed to uncover the underlying deformation mechanisms in the bulk nc alloy. This advanced nondestructive technique allows full penetration of the high-energy x rays into the bulk specimen, which, therefore, provides good statistics. Due to the small beam size of 50⫻ 50 m2, HEXRD can detect the area very close to the fracture surface, where the sample experienced a large plastic strain. This information is valuable in understanding the fracture mechanisms of nc materials. The bulk nc Ni–Fe alloy was synthesized by a pulsed electrodeposition technique. For the HEXRD measurements, an x-ray beam with an energy of 115 keV was used to provide diffraction patterns 共Debye-Scherrer rings兲 collected by a two-dimensional detector 共Mar345兲 in transmission geometry. Diffraction peaks were fitted by the Cauchy peak shape, and the physical peak shape was obtained by deconvolution of the measured peaks from the instrumental peak shape 共calibrated by Si powders兲. The Williamson-Hall integral breadth method was employed to determine the crystallite domain size and the microstrain from the 共111兲 and 共222兲 peaks.18 The inverse pole figures were determined directly from the Debye-Scherrer rings.19 TEM observations were carried out using a Schottky field-emission gun FEI Tecnai
0003-6951/2006/89共10兲/101918/3/$23.00 89, 101918-1 © 2006 American Institute of Physics Downloaded 23 Sep 2006 to 169.237.215.179. Redistribution subject to AIP license or copyright, see http://apl.aip.org/apl/copyright.jsp
101918-2
Fan et al.
FIG. 1. 共Color online兲 Engineering stress-strain curve for the as-deposited bulk nc Ni–Fe alloy 共inset兲 at a strain rate, ˙ , of 10−2 s−1. The alloy exhibited a combination of high strength and good tensile ductility. The elastic strain, e, the uniform plastic strain, p, and the postnecking plastic strain, n, are illustrated.
F20 UT microscope with a spatial resolution of 0.14 nm operating at 200 kV. The grain size of the as-prepared bulk nc Ni–Fe plate ranges from about 2 to 50 nm with an average value of about 23 nm. Figure 1 displays a typical engineering stress-strain curve at a strain rate, , of 10−2 s−1 for a dog-bone specimen with a gauge section of 4 ⫻ 3 mm2 and a gauge length of 25 mm 共included in Fig. 1兲. The bulk nc specimen has an ultimate tensile strength of 2.0 GPa, and an elongation to failure of f = e + p + n = 7.2%, where e is the elastic strain, p is the uniform plastic strain, and n is the postnecking plastic strain 共Fig. 1兲. The sample after the tensile tests was examined by HEXRD. By transmitting the sample from the facture surface
Appl. Phys. Lett. 89, 101918 共2006兲
共necking area兲 to the uniformly deforming region with a step of 75 m, the distributions of the crystallite domain size and microstrain as a function of the distance from the fracture surface, L, were mapped with a spatial resolution of about 50 m 关Fig. 2共a兲兴. The average crystallite domain size, determined by HEXRD, is about 21.8 nm, which is in good agreement with 23 nm measured by TEM. The crystalline domain size increases from 21.8 nm in the as-deposited state to about 32.5 nm during the uniform plastic deformation. Moreover, the crystallite domain size increases rapidly from 32.5 nm to a maximum of 307 nm 共L = 0.075 mm兲 when L 艋 0.9 mm. This abrupt increase in the crystallite domain size is due to the large localized plastic strain in the necking region. Assuming that the sample fractured in the middle of the necking area, the total length of the necking area along the loading direction is, therefore, about 1.8 mm. Thus, the average plastic strain in the necking region was estimated by p + n ⫻ Lgauge / 1.8= 20.3%, where Lgauge is the gauge length of the tensile specimen 共25 mm兲 and n = 1.1%. This value is much larger than the observed uniform plastic strain, p = 5% 共Fig. 1兲. As shown in Fig. 2共a兲, the microstrain, which is microscopically related to the dislocation density of the materials, did not change after the plastic deformation in the uniformly deforming region 共L ⬎ 0.9 mm兲 and fluctuates around the value of 0.255% established from the as-deposited specimen. A slight increase in the microstrain was noticed with L 艋 0.9 mm. These results indicate that the bulk nc Ni–Fe alloy, if deformed uniformly, does not exhibit dislocation storage, an observation that agrees well with the results obtained by Budrovic et al.12 The increase in the crystalline domain sizes 关Fig. 2共a兲兴 indicates that grain coarsening occurs during the plastic de-
FIG. 2. 共Color online兲 Crystallite domain size and the microstrain 共a兲, measured by HEXRD, as a function of the distance from the fracture surface, L. The red arrows in 共a兲 point to the place where the necking starts in the sample. The inverse pole figures for the sample before the deformation 共b兲 and after deformation at point I 共c兲 and point II 共d兲, as marked in 共a兲, indicate a texture development after the plastic deformation. Downloaded 23 Sep 2006 to 169.237.215.179. Redistribution subject to AIP license or copyright, see http://apl.aip.org/apl/copyright.jsp
101918-3
Appl. Phys. Lett. 89, 101918 共2006兲
Fan et al.
the dislocation motion becomes more active in the necking region, leading to a strong texture and a slight increase in the microstrain of the sample. The slight increase in the microstrain suggests the presence of dislocation storage, which may be due to the larger grain sizes in the necking region as compared with the relatively small grain sizes in the uniformly deforming region.
FIG. 3. Bright-field TEM image 共a兲 for the bulk nc Ni–Fe alloy after the plastic deformation shown in the inset of Fig. 1. The grain-size distribution 共b兲 indicates that the average grain size after the plastic deformation is about 73 nm, approximately three times greater than that of the as-deposited state 共dashed line兲.
formation. Molecular dynamics simulations suggest that grain coarsening caused by plastic deformation in nc metals and alloys can be due to GB migration20 and/or grain rotation.21 The grain coarsening due to the grain rotation was theoretically analyzed by Li four decades ago.22 It is suggested that grain coarsening measured by HEXRD is due to GB migration, leading to an increase in the subgrain size. To confirm the existence of the grain rotation during plastic deformation, the texture of the bulk nc specimen before and after deformation was quantitatively measured by the inverse pole figures along the loading direction 共LD兲. Both GB sliding 共accommodated by the GB diffusion兲 and dislocation glide in the preferred active slip systems can induce grain rotation. Grain rotation due to the GB sliding will not introduce a texture, whereas grain rotation due to a dislocation mechanism will introduce a texture in the sample.23 Figures 2共b兲–2共d兲 are the inverse pole figures measured in the asdeposited nc specimen and in the specimen after the plastic deformation 关I and II as marked in Fig. 2共a兲兴, respectively. Before the deformation, no obvious crystallographic anisotropy was detected in the plane 关Fig. 2共b兲兴. After the deformation, the texture components, i.e., LD 储 具111典 and LD 储 具001典, were developed at L = 1.35 mm 关Fig. 2共c兲兴. These texture components were typically observed in the fccstructured polycrystalline materials under uniaxial deformation.24 The volume fraction of these texture components increases at L = 0.15 mm 关Fig. 2共d兲兴. Therefore, the grain rotation due to the dislocation motion takes place, particularly during the postnecking deformation. To confirm that grain coarsening occurs, the deformed sample was further examined by TEM. Figure 3共a兲 shows a typical bright-field TEM image after the uniaxial tensile plastic deformation at ˙ = 10−2 s−1. Compared with the asdeposited sample, the grain size increases significantly after the plastic deformation. The average grain size, as measured by the grain-size distribution 关Fig. 3共b兲兴, is about 73 nm. This stress-induced grain coarsening observed by HEXRD and TEM experiments during uniaxial tensile tests is consistent with the previous reports using the indentation technique.25,26 Our HEXRD results indicate that dislocation motion, which causes a texture development in the deformed samples, plays a role during the uniaxial tensile deformation of the bulk nc Ni–Fe alloy. The dislocation density remains the same in the uniformly deforming region since the microstrain measured by HEXRD is almost unchanged. However,
This work was supported by the National Science Foundation 共NSF兲 International Materials Institutes 共IMI兲 Program 共DMR-0231320兲. The microscopy was performed at the National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, supported by the Director, Office of Science, of the U.S. Department of Energy 共DOE兲 under contract No. DE-AC02-05CH11231 and by Grant No. DEFG02-03ER46057. One of the authors 共Y.D.W.兲 was also supported by the National Natural Science Foundation of China 共Grant No. 50471026兲. The use of the Advanced Photon Source was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Science, under Contract No. W-31-109-ENG-38. J. Schiotz, F. D. Di Tolla, and K. W. Jacobsen, Nature 共London兲 391, 561 共1998兲. H. Van Swygenhoven and P. M. Derlet, Phys. Rev. B 64, 224105 共2001兲. 3 D. Wolf, V. Yamakov, S. R. Phillpot, A. Mukherjee, and H. Gleiter, Acta Mater. 53, 1 共2005兲. 4 H. Mughrabi, Plastic Deformation and Fracture of Materials 共WileyVCH, Weinheim, Germany, 1991兲, Vol. 6, p. 1. 5 M. Ke, S. A. Hackney, W. W. Milligan, and E. C. Aifantis, Nanostruct. Mater. 5, 689 共1995兲. 6 R. Z. Valiev, R. K. Islamgaliev, and I. V. Alexandrov, Prog. Mater. Sci. 45, 103 共2000兲. 7 R. C. Hugo, H. Kung, J. R. Weertman, R. Mitra, J. A. Knapp, and D. M. Follstaedt, Acta Mater. 51, 1937 共2003兲. 8 M. W. Chen, E. Ma, K. J. Hemker, H. W. Sheng, Y. M. Wang, and X. M. Cheng, Science 300, 1275 共2003兲. 9 X. Z. Liao, F. Zhou, E. J. Lavernia, S. G. Srinivasan, M. I. Baskes, D. W. He, and Y. T. Zhu, Appl. Phys. Lett. 83, 632 共2003兲. 10 K. S. Kumar, S. Suresh, M. F. Chisholm, J. A. Horton, and P. Wang, Acta Mater. 51, 387 共2003兲. 11 C. A. Schuh, T. G. Nieh, and H. Iwasaki, Acta Mater. 51, 431 共2003兲. 12 Z. Budrovic, H. Van Swygenhoven, P. M. Derlet, S. Van Petegem, and B. Schmitt, Science 304, 273 共2004兲. 13 Z. W. Shan, E. A. Stach, J. M. K. Wiezorek, J. A. Knapp, D. M. Follstaedt, and S. X. Mao, Science 305, 654 共2004兲. 14 K. M. Youssef, R. O. Scattergood, K. L. Murty, J. A. Horton, and C. C. Koch, Appl. Phys. Lett. 87, 091904 共2005兲. 15 R. Z. Valiev, I. V. Alexandrov, Y. T. Zhu, and T. C. Lowe, J. Mater. Res. 17, 5 共2002兲. 16 Y. M. Wang, M. W. Chen, F. H. Zhou, and E. Ma, Nature 共London兲 419, 912 共2002兲. 17 G. J. Fan, H. Choo, P. K. Liaw, and E. J. Lavernia, Acta Mater. 54, 1759 共2006兲. 18 H. P. Klug and L. E. Alexander, X-ray Diffraction Procedures, 2nd ed. 共Wiley, New York, 1974兲, p. 505. 19 Y. D. Wang, R. L. Peng, J. Almer, M. Oden, Y. D. Liu, J. N. Deng, C. S. He, L. Chen, Q. L. Li, and L. Zuo, Adv. Mater. 共Weinheim, Ger.兲 17, 1221 共2005兲. 20 J. Schiotz, Mater. Sci. Eng., A 375-377, 975 共2004兲. 21 D. Moldovan, V. Yamakov, D. Wolf, and S. R. Phillpot, Phys. Rev. Lett. 89, 206101 共2002兲. 22 J. C. M. Li, J. Appl. Phys. 33, 2958 共1962兲. 23 E. Ma, Science 305, 623 共2004兲. 24 F. J. Humphreys and M. Hatherly, Recrystallization and Related Annealing Phenomena, 2nd ed. 共Elsevier, Oxford, 2004兲, p. 67. 25 K. Zhang, J. R. Weertman, and J. A. Eastman, Appl. Phys. Lett. 85, 5197 共2004兲. 26 M. Jin, A. M. Minor, E. A. Stach, and J. W. Morris, Acta Mater. 52, 5381 共2004兲. 1
2
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Uniaxial tensile plastic deformation of a bulk nanocrystalline alloy studied by a high-energy x-ray diffraction technique G. J. Fana兲 Department of Materials Science and Engineering, The University of Tennessee, Knoxville, Tennessee 37996
L. F. Fu Department of Chemical Engineering and Materials Science, University of California, Davis, California 95616
Y. D. Wang Department of Materials Science and Engineering, The University of Tennessee, Knoxville, Tennessee 37996 and School of Materials and Metallurgy, Northeastern University, Shenyang 110004, China
Y. Ren Experimental Facilities Division, Advanced Phonon Source, Argonne National Laboratory, Argonne, Illinois 60439
H. Choo, P. K. Liaw, and G. Y. Wang Department of Materials Science and Engineering, The University of Tennessee, Knoxville, Tennessee 37996
N. D. Browning National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, California 94720
共Received 27 June 2006; accepted 19 July 2006; published online 8 September 2006兲 By employing a high-energy x-ray diffraction technique, the authors report that uniaxial tensile plastic deformation induced the grain growth and texture development in a bulk nanocrystalline Ni–Fe alloy. The effects become more pronounced with increasing the plastic strain 共closer to the fracture surface兲. The texture development accompanying the grain rotation indicates that dislocation motion contributed to the observed plasticity in the nanocrystalline Ni–Fe alloy. The quantitative experimental data suggest that the dislocation storage was absent in the uniformly deforming region; whereas the dislocation storage was present in the necking region, where the grain growth was substantial. © 2006 American Institute of Physics. 关DOI: 10.1063/1.2348783兴 Nanocrystalline 共nc兲 materials with average grain sizes less than 100 nm exhibit unusual deformation mechanisms. Molecular dynamics simulations have suggested that the plasticity in the nc metals and alloys is mainly mediated by grain-boundary 共GB兲 activities,1–3 in contrast to the dislocation pileups at the GBs in their coarse-grained polycrystalline counterparts.4 However, experimental verification of deformation mechanisms in nc metals and alloys is complicated by various competing mechanisms.5–14 A complete understanding of these competing mechanisms poses a great challenge experimentally, particularly in the threedimensional 共3D兲 bulk nc specimens. Previous investigations of deformation mechanisms of nc metals and alloys were primarily based on the in situ straining transmission electron microscopy 共TEM兲 observations in the two-dimensional thin foils,5,7,10,13,14 where artifacts often exist. The study of deformation mechanisms in a 3D bulk nc specimen was rare in the literature due to the limitations of the available experimental techniques, i.e., TEM, and to the intrinsic brittleness often observed in the bulk nc specimens. A good tensile ductility was recently realized in the bulk nc metals and alloys prepared by severe plastic deformation techniques 共i.e., mechanical milling and cold rolling兲.14–17 In the present investigation, the bulk nc a兲
Electronic mail: [email protected]
Ni-18 wt % Fe 共weight: wt %兲 sheets, which exhibit a good combination of the very high strength and a large tensile plastic strain, were produced using a pulsed electrodeposition technique. A high-energy x-ray diffraction 共HEXRD兲 technique was employed to uncover the underlying deformation mechanisms in the bulk nc alloy. This advanced nondestructive technique allows full penetration of the high-energy x rays into the bulk specimen, which, therefore, provides good statistics. Due to the small beam size of 50⫻ 50 m2, HEXRD can detect the area very close to the fracture surface, where the sample experienced a large plastic strain. This information is valuable in understanding the fracture mechanisms of nc materials. The bulk nc Ni–Fe alloy was synthesized by a pulsed electrodeposition technique. For the HEXRD measurements, an x-ray beam with an energy of 115 keV was used to provide diffraction patterns 共Debye-Scherrer rings兲 collected by a two-dimensional detector 共Mar345兲 in transmission geometry. Diffraction peaks were fitted by the Cauchy peak shape, and the physical peak shape was obtained by deconvolution of the measured peaks from the instrumental peak shape 共calibrated by Si powders兲. The Williamson-Hall integral breadth method was employed to determine the crystallite domain size and the microstrain from the 共111兲 and 共222兲 peaks.18 The inverse pole figures were determined directly from the Debye-Scherrer rings.19 TEM observations were carried out using a Schottky field-emission gun FEI Tecnai
0003-6951/2006/89共10兲/101918/3/$23.00 89, 101918-1 © 2006 American Institute of Physics Downloaded 23 Sep 2006 to 169.237.215.179. Redistribution subject to AIP license or copyright, see http://apl.aip.org/apl/copyright.jsp
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Fan et al.
FIG. 1. 共Color online兲 Engineering stress-strain curve for the as-deposited bulk nc Ni–Fe alloy 共inset兲 at a strain rate, ˙ , of 10−2 s−1. The alloy exhibited a combination of high strength and good tensile ductility. The elastic strain, e, the uniform plastic strain, p, and the postnecking plastic strain, n, are illustrated.
F20 UT microscope with a spatial resolution of 0.14 nm operating at 200 kV. The grain size of the as-prepared bulk nc Ni–Fe plate ranges from about 2 to 50 nm with an average value of about 23 nm. Figure 1 displays a typical engineering stress-strain curve at a strain rate, , of 10−2 s−1 for a dog-bone specimen with a gauge section of 4 ⫻ 3 mm2 and a gauge length of 25 mm 共included in Fig. 1兲. The bulk nc specimen has an ultimate tensile strength of 2.0 GPa, and an elongation to failure of f = e + p + n = 7.2%, where e is the elastic strain, p is the uniform plastic strain, and n is the postnecking plastic strain 共Fig. 1兲. The sample after the tensile tests was examined by HEXRD. By transmitting the sample from the facture surface
Appl. Phys. Lett. 89, 101918 共2006兲
共necking area兲 to the uniformly deforming region with a step of 75 m, the distributions of the crystallite domain size and microstrain as a function of the distance from the fracture surface, L, were mapped with a spatial resolution of about 50 m 关Fig. 2共a兲兴. The average crystallite domain size, determined by HEXRD, is about 21.8 nm, which is in good agreement with 23 nm measured by TEM. The crystalline domain size increases from 21.8 nm in the as-deposited state to about 32.5 nm during the uniform plastic deformation. Moreover, the crystallite domain size increases rapidly from 32.5 nm to a maximum of 307 nm 共L = 0.075 mm兲 when L 艋 0.9 mm. This abrupt increase in the crystallite domain size is due to the large localized plastic strain in the necking region. Assuming that the sample fractured in the middle of the necking area, the total length of the necking area along the loading direction is, therefore, about 1.8 mm. Thus, the average plastic strain in the necking region was estimated by p + n ⫻ Lgauge / 1.8= 20.3%, where Lgauge is the gauge length of the tensile specimen 共25 mm兲 and n = 1.1%. This value is much larger than the observed uniform plastic strain, p = 5% 共Fig. 1兲. As shown in Fig. 2共a兲, the microstrain, which is microscopically related to the dislocation density of the materials, did not change after the plastic deformation in the uniformly deforming region 共L ⬎ 0.9 mm兲 and fluctuates around the value of 0.255% established from the as-deposited specimen. A slight increase in the microstrain was noticed with L 艋 0.9 mm. These results indicate that the bulk nc Ni–Fe alloy, if deformed uniformly, does not exhibit dislocation storage, an observation that agrees well with the results obtained by Budrovic et al.12 The increase in the crystalline domain sizes 关Fig. 2共a兲兴 indicates that grain coarsening occurs during the plastic de-
FIG. 2. 共Color online兲 Crystallite domain size and the microstrain 共a兲, measured by HEXRD, as a function of the distance from the fracture surface, L. The red arrows in 共a兲 point to the place where the necking starts in the sample. The inverse pole figures for the sample before the deformation 共b兲 and after deformation at point I 共c兲 and point II 共d兲, as marked in 共a兲, indicate a texture development after the plastic deformation. Downloaded 23 Sep 2006 to 169.237.215.179. Redistribution subject to AIP license or copyright, see http://apl.aip.org/apl/copyright.jsp
101918-3
Appl. Phys. Lett. 89, 101918 共2006兲
Fan et al.
the dislocation motion becomes more active in the necking region, leading to a strong texture and a slight increase in the microstrain of the sample. The slight increase in the microstrain suggests the presence of dislocation storage, which may be due to the larger grain sizes in the necking region as compared with the relatively small grain sizes in the uniformly deforming region.
FIG. 3. Bright-field TEM image 共a兲 for the bulk nc Ni–Fe alloy after the plastic deformation shown in the inset of Fig. 1. The grain-size distribution 共b兲 indicates that the average grain size after the plastic deformation is about 73 nm, approximately three times greater than that of the as-deposited state 共dashed line兲.
formation. Molecular dynamics simulations suggest that grain coarsening caused by plastic deformation in nc metals and alloys can be due to GB migration20 and/or grain rotation.21 The grain coarsening due to the grain rotation was theoretically analyzed by Li four decades ago.22 It is suggested that grain coarsening measured by HEXRD is due to GB migration, leading to an increase in the subgrain size. To confirm the existence of the grain rotation during plastic deformation, the texture of the bulk nc specimen before and after deformation was quantitatively measured by the inverse pole figures along the loading direction 共LD兲. Both GB sliding 共accommodated by the GB diffusion兲 and dislocation glide in the preferred active slip systems can induce grain rotation. Grain rotation due to the GB sliding will not introduce a texture, whereas grain rotation due to a dislocation mechanism will introduce a texture in the sample.23 Figures 2共b兲–2共d兲 are the inverse pole figures measured in the asdeposited nc specimen and in the specimen after the plastic deformation 关I and II as marked in Fig. 2共a兲兴, respectively. Before the deformation, no obvious crystallographic anisotropy was detected in the plane 关Fig. 2共b兲兴. After the deformation, the texture components, i.e., LD 储 具111典 and LD 储 具001典, were developed at L = 1.35 mm 关Fig. 2共c兲兴. These texture components were typically observed in the fccstructured polycrystalline materials under uniaxial deformation.24 The volume fraction of these texture components increases at L = 0.15 mm 关Fig. 2共d兲兴. Therefore, the grain rotation due to the dislocation motion takes place, particularly during the postnecking deformation. To confirm that grain coarsening occurs, the deformed sample was further examined by TEM. Figure 3共a兲 shows a typical bright-field TEM image after the uniaxial tensile plastic deformation at ˙ = 10−2 s−1. Compared with the asdeposited sample, the grain size increases significantly after the plastic deformation. The average grain size, as measured by the grain-size distribution 关Fig. 3共b兲兴, is about 73 nm. This stress-induced grain coarsening observed by HEXRD and TEM experiments during uniaxial tensile tests is consistent with the previous reports using the indentation technique.25,26 Our HEXRD results indicate that dislocation motion, which causes a texture development in the deformed samples, plays a role during the uniaxial tensile deformation of the bulk nc Ni–Fe alloy. The dislocation density remains the same in the uniformly deforming region since the microstrain measured by HEXRD is almost unchanged. However,
This work was supported by the National Science Foundation 共NSF兲 International Materials Institutes 共IMI兲 Program 共DMR-0231320兲. The microscopy was performed at the National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, supported by the Director, Office of Science, of the U.S. Department of Energy 共DOE兲 under contract No. DE-AC02-05CH11231 and by Grant No. DEFG02-03ER46057. One of the authors 共Y.D.W.兲 was also supported by the National Natural Science Foundation of China 共Grant No. 50471026兲. The use of the Advanced Photon Source was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Science, under Contract No. W-31-109-ENG-38. J. Schiotz, F. D. Di Tolla, and K. W. Jacobsen, Nature 共London兲 391, 561 共1998兲. H. Van Swygenhoven and P. M. Derlet, Phys. Rev. B 64, 224105 共2001兲. 3 D. Wolf, V. Yamakov, S. R. Phillpot, A. Mukherjee, and H. Gleiter, Acta Mater. 53, 1 共2005兲. 4 H. Mughrabi, Plastic Deformation and Fracture of Materials 共WileyVCH, Weinheim, Germany, 1991兲, Vol. 6, p. 1. 5 M. Ke, S. A. Hackney, W. W. Milligan, and E. C. Aifantis, Nanostruct. Mater. 5, 689 共1995兲. 6 R. Z. Valiev, R. K. Islamgaliev, and I. V. Alexandrov, Prog. Mater. Sci. 45, 103 共2000兲. 7 R. C. Hugo, H. Kung, J. R. Weertman, R. Mitra, J. A. Knapp, and D. M. Follstaedt, Acta Mater. 51, 1937 共2003兲. 8 M. W. Chen, E. Ma, K. J. Hemker, H. W. Sheng, Y. M. Wang, and X. M. Cheng, Science 300, 1275 共2003兲. 9 X. Z. Liao, F. Zhou, E. J. Lavernia, S. G. Srinivasan, M. I. Baskes, D. W. He, and Y. T. Zhu, Appl. Phys. Lett. 83, 632 共2003兲. 10 K. S. Kumar, S. Suresh, M. F. Chisholm, J. A. Horton, and P. Wang, Acta Mater. 51, 387 共2003兲. 11 C. A. Schuh, T. G. Nieh, and H. Iwasaki, Acta Mater. 51, 431 共2003兲. 12 Z. Budrovic, H. Van Swygenhoven, P. M. Derlet, S. Van Petegem, and B. Schmitt, Science 304, 273 共2004兲. 13 Z. W. Shan, E. A. Stach, J. M. K. Wiezorek, J. A. Knapp, D. M. Follstaedt, and S. X. Mao, Science 305, 654 共2004兲. 14 K. M. Youssef, R. O. Scattergood, K. L. Murty, J. A. Horton, and C. C. Koch, Appl. Phys. Lett. 87, 091904 共2005兲. 15 R. Z. Valiev, I. V. Alexandrov, Y. T. Zhu, and T. C. Lowe, J. Mater. Res. 17, 5 共2002兲. 16 Y. M. Wang, M. W. Chen, F. H. Zhou, and E. Ma, Nature 共London兲 419, 912 共2002兲. 17 G. J. Fan, H. Choo, P. K. Liaw, and E. J. Lavernia, Acta Mater. 54, 1759 共2006兲. 18 H. P. Klug and L. E. Alexander, X-ray Diffraction Procedures, 2nd ed. 共Wiley, New York, 1974兲, p. 505. 19 Y. D. Wang, R. L. Peng, J. Almer, M. Oden, Y. D. Liu, J. N. Deng, C. S. He, L. Chen, Q. L. Li, and L. Zuo, Adv. Mater. 共Weinheim, Ger.兲 17, 1221 共2005兲. 20 J. Schiotz, Mater. Sci. Eng., A 375-377, 975 共2004兲. 21 D. Moldovan, V. Yamakov, D. Wolf, and S. R. Phillpot, Phys. Rev. Lett. 89, 206101 共2002兲. 22 J. C. M. Li, J. Appl. Phys. 33, 2958 共1962兲. 23 E. Ma, Science 305, 623 共2004兲. 24 F. J. Humphreys and M. Hatherly, Recrystallization and Related Annealing Phenomena, 2nd ed. 共Elsevier, Oxford, 2004兲, p. 67. 25 K. Zhang, J. R. Weertman, and J. A. Eastman, Appl. Phys. Lett. 85, 5197 共2004兲. 26 M. Jin, A. M. Minor, E. A. Stach, and J. W. Morris, Acta Mater. 52, 5381 共2004兲. 1
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