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Surface nanocrystallization of Ti–45Al–7Nb–0.3W intermetallics induced by surface mechanical grinding treatment
Materials Letters 166 (2016) 59–62
Contents lists available at ScienceDirect
Materials Letters journal homepage: www.elsevier.com/locate/matlet
Surface nanocrystallization of Ti–45Al–7Nb–0.3W intermetallics induced by surface mechanical grinding treatment Kun Zhao, Yong Liu n, Tianhang Yao, Bin Liu, Yuehui He State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 24 July 2015 Received in revised form 30 November 2015 Accepted 4 December 2015 Available online 8 December 2015
Surface mechanical grinding treatment was conducted on a Ti–45Al–7Nb–0.3W intermetallics. The results show that the penetration depth plays an important role in the microstructural evolution. The γα2 phase transformation is observed at a penetration depth of 100 μm. Nanocrystallized α2 grains (20 nm) are found in the treated surface at a penetration depth of 200 μm. After the surface mechanical grinding treatment, the micro-hardness of the top surface increases significantly, which is 50% higher than that of the substrate. It is also found that the γ-α2 phase transformation is induced by the localized strain in the surface, while the mechanism for the nanocrystallization is proposed to be the twin–twin intersection. & 2015 Published by Elsevier B.V.
Keywords: Intermetallic alloys and compounds TiAl Surface mechanical grinding treatment Nanocrystallization Phase transformation
1. Introduction
2. Experimental
The surface nanocrystallization has been widely developed to synthesize a gradient grain structure, in order to improve the mechanical properties of metals and alloys. As a new method for the surface modification, the surface mechanical grinding treatment (SMGT) induces a very high rate shear deformation and a high strain gradient in the top surface of specimens. The shear strain rates can be 102–104 s 1 and the corresponding accumulative equivalent strains are 15–30 [1]. A gradient nanostructure can be formed after SMGT. Lu et al. has fabricated a gradient nanograined layer on pure copper, which has an extraordinary intrinsic tensile plasticity [2]. Besides, the grain refinement and the plastic deformation at large strains also result in disordering of the alloys and even the phase transformation [3,4]. It was found that γ-TiAl alloy powder was partially disordered and transformed to nanometer-sized hcp phase during the ball milling [5]. Similar results were reported in bulk TiAl alloys after the high pressure torsion (HPT), but the nanocrystallization mechanism was not clarified [6]. In this work, SMGT was conducted to form a nanostructured layer in a Ti–45Al–7Nb–0.3W alloy. The microstructural evolution and the mechanism of nanocrystallization were discussed.
An as-HIPed Ti–45Al–7Nb–0.3W (at%) alloy was used in the present study. The original SMGT was conducted on the circumference of cylinders [2]. A modified SMGT was performed on plates at room temperature. A hemispherical WC–Co tip (r ¼6 mm) was penetrated into the surface of specimen, and then moved on the plate from one side to the other. No material removal or contamination was found. The plastic deformation was uniform in the surface layer with a small roughness (Ra ¼0.35 μm). The preset penetration depth of the WC–Co tip into the sample per pass was 50 μm. The specimens were treated with the same operating parameters for 2 and 4 times. The total penetration depths (ap) were 50 μm, 100 μm and 200 μm, separately. The phase characterization of the treated surface was conducted on a Rigaku D/max 2550VB þ XRD equipment with a Cu radiation. A detailed description of the long range order (LRO) parameter and the mass fractions of phases can be referred to Ref. [6]. The mass fraction of γ phase is calculated using the matrixflushing method. For the determination of the mass fractions of phases, the intensity of the (200)γ was compared to the intensity of the (201)α2. The LRO parameters of the γ-TiAl phase were calculated from the intensities of the (100)γ superlattice and (200)γ peaks, with reference to the fully ordered alloys. The microstructural characterization of the specimens was carried out on a Leica DM4500P optical microscope and a transmission electron microscope (JEOL JEM 2010F). The TEM foils at different depths were obtained by mechanically polishing the
n
Corresponding author. E-mail address: [email protected] (Y. Liu).
http://dx.doi.org/10.1016/j.matlet.2015.12.025 0167-577X/& 2015 Published by Elsevier B.V.
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K. Zhao et al. / Materials Letters 166 (2016) 59–62
sample from the untreated side to a thickness of 80 μm, and then by electro-polishing through the conventional twin-jet technique with an electrolyte consisting of 60% methanol, 35% normal butanol and 5% perchloric acid. Micro-hardness measurements were conducted on a BUEHLER OmniMet MHT automatic testing system with a load of 10 g and a loading time of 10–15 s.
3. Results Fig. 1(a) shows XRD patterns of the Ti–45Al–7Nb–0.3W alloy treated with different penetration depths. The substrate alloy exhibited two intermetallic phases, mainly γ phase and a small amount of α2 phase. The intensities of the superlattice peaks, such as (110)γ and (001)γ decreased continuously with the total penetration depth increasing. In the meantime, the intensities of the fundamental peaks, such as (002)γ and (200)γ, decreased continuously. At a total penetration depth of 200 μm, the peaks were broadened significantly, indicating a reduction in the grain size. The LRO parameter and the mass fraction of the γ phase are displayed in Table 1. The LRO parameter did not change at a total penetration depth of 50 μm. With the increase of the total penetration depth from 50 μm to 200 μm, the LRO parameter of γ phase decreased from 100% to 20%. When the penetration depth increased from 0 to 200 μm, the mass fraction of γ phase decreased from 74% to 14%. Therefore, the SMGT induced the disordering of γ phase and γ-α2 transformation in the surface. Fig. 1(b) indicates that the substrate alloy was a near-γ microstructure with an average grain size of 8 μm. No significant microstructural change occurred at a penetration depth of 50 μm
Table 1 Structure and properties of TiAl after SMGT. Penetration depth (ap) (μm)
LRO parameter of γ phase (%)
Mass fraction of γ phase (wt%)
0 50 100 200
100 100 94 20
74 48 44 14
(not shown). When the penetration depth increased to 100 μm, a bright layer of a thickness of 18 μm formed in the treated surface (Fig. 1(c)). According to the XRD patterns in Fig. 1(a), the bright layer consisted of α2 phase. It indicates that γ phase transformed to α2 phase in the treated surface. In Fig. 1(d), the thickness of the α2 rich layer increased to 26 μm at a penetration depth of 200 μm. In order to analyze the phase transformation and the grain refinement, the microstructures in different depths were observed by using TEM. The bright field (BF) images and corresponding selected area electron diffraction (SAED) patterns are shown in Fig. 2. Fig. 2(a) and (b) indicates that the surface layer consisted of nanocrystalline α2 phase (about 20 nm). In Fig. 2(c), the twin–twin intersection occurred at a depth of 80 μm from the surface. The density of twins was very high, and the spacing between neighboring twin boundaries was 30–90 nm. The corresponding SAED patterns in Fig. 2(d) indicated large misorientations. Twin–twin intersections were also found at a depth of 170 μm from the surface (Fig. 2(e)). Compared with the SAED patterns in Fig. 2(d), the spots in the Fig. 2(f) were not elongated, showing small misorientations. No twins or dislocations were found in γ grains of the
Fig. 1. (a) XRD patterns of as-SMGTed TiAl alloys and substrate, (b) microstructure of substrate, (c) microstructure of as-SMGTed specimens at penetration depth of 100 μm, and (d) microstructure of as-SMGTed specimens at penetration depth of 200 μm.
K. Zhao et al. / Materials Letters 166 (2016) 59–62
61
Fig. 2. BF images and corresponding SAED patterns of as-SMGTed TiAl alloys (ap ¼ 200 μm) in different depths: (a),(b) the top surface, nanocrystallized α2 grains; (c),(d) in a depth of 80 μm, (c) shows twin–twin intersection, (d) shows large misorientations; (e),(f) in a depth of 170 μm, (e) shows twin–twin intersection, (f) shows small misorientations; and (g),(h) the substrate alloy.
substrate alloy (Fig. 2(g)). The SAED patterns in Fig. 2(h) indicated that there was a perfect lattice structure. The micro-hardness changed with the depth from the top surface (in Fig. 3). In the as-SMGTed alloy, the highest value of
hardness was 685HV0.01 in the top layer, which is 50% higher than that of the substrate (453HV0.01). In the subsurface layer, which consisted of nanometer-sized grains, the micro-hardness was around 600HV0.01. With the increase of the depth, the micro-
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K. Zhao et al. / Materials Letters 166 (2016) 59–62
considered as a pseudo f.c.c structure. The standard α2 phase is a D019 structure, which is an ordered h.c.p structure. The structural change from f.c.c to h.c.p can be formed by the gliding of 1/6 < 112¯ ] Shockley partial dislocations on {111} planes [9]. When one 1/6 < 112¯ ] Shockley partial dislocation sweeps on a {111} plane, two layers of h.c.p structure will be converted from the f.c.c structure. After that, if another 1/6 < 112¯ ] Shockley partial dislocation glides on every other {111} plane, another two layers of h. c.p structure will be obtained. The γ-α2 phase transformation can be treated as a stress-induced transformation. The stacking sequence changes from cubic “ABCABC” to hexagonal “ABABAB”. The spacing between {111} close-packed planes in the γ phase is 0.231 nm. The distance between the nearest atoms in the γ phase is 0.281 nm. For the stress-induced α2, a ¼0.562 nm and c¼0.462 nm, which is close to the cell parameters for standard α2 phase (a ¼0.580 nm, c ¼0.465 nm). Fig. 3. Micro-hardness of as-SMGTed alloys (ap ¼ 200 μm), the hardness decreases with the distance from the surface.
5. Conclusions
hardness sharply decreased and became stable in the substrate.
4. Discussion In TiAl alloy, twins nucleate from the superposition of extended stacking faults on alternate {111} planes of γ phase. The addition of Nb lowers the intrinsic stacking fault energy, and promotes the formation of the mechanical twinning in high-Nb TiAl alloys. Due to the difficult deformability of α2 phase, during SMGT, the plastic deformation mainly occurs in γ phase, and 1/6 < 112¯ ]{111} twining is the deformation mechanism [7]. When twinning takes place, the γ grain is divided into the twin/matrix lamellar structure with a specific orientation. The deformation twinning can be regarded as the first step for the nanocrystallization [8]. The second step is the refinement of twin/ matrix lamellae into equiaxed nanometer-sized grains through the twin–twin intersection [8]. When the driving force for the deformation twinning is large enough to overcome the barriers of the encountered twin boundaries, twin–twin intersection occurs. Twin–twin intersections can divide twin/matrix lamellae into rhombic blocks. The intersected volume undergoes a rotation to change its orientation. When the shear strain increases, the multiplication of deformation twinning leads to more twin–twin intersections, refining the structure into equiaxed grains with large misorientations in nanometer scale [8]. The stress induces the phase transformations between γ and α2 phase in TiAl alloys [9]. Moreover, γ-α2 transformations is possible in the twin–twin intersection region of deformed TiAl alloys [9]. γ phase is L10 structure (a ¼0.398 nm and c¼ 0.405 nm). The c/ a ratio of γ phase (1.03) is close to 1, and thus γ phase can be
(1) Through SMGT, a gradient grain structure can be formed in the surface of TiAl alloy. In the top surface, nanocrystallized α2 grains (∼20 nm) can be obtained. Both the γ-α2 phase transformation and twin–twin intersections occur with the increase of the penetration depth. (2) After SMGT, the micro-hardness of the top surface increases significantly, which is 50% higher than that of the substrate. (3) The γ-α2 phase transformation is induced by the localized strain in the surface, while the mechanism for the nanocrystallization is proposed to be the twin–twin intersection.
Acknowledgments The work was financially supported by the Project of Innovation-driven Plan in Central South University (2015CX004) and the National Natural Science Foundation of China (51301203).
Reference [1] [2] [3] [4] [5] [6]
X.C. Liu, H.W. Zhang, K. Lu, Science 342 (2013) 337–340. T.H. Fang, W.L. Li, N.R. Tao, K. Lu, Science 331 (2011) 1587–1590. H.W. Zhang, Z.K. Hei, G. Liu, J. Lu, K. Lu, Acta Mater. 51 (2003) 1871–1881. X. Wu, N. Tao, Y. Hong, G. Liu, B. Xu, J. Lu, K. Lu, Acta Mater. 53 (2005) 681–691. T. Klassen, M. Oehring, R. Bormann, Acta Mater. 45 (1997) 3935–3948. A.V. Korznikov, O. Dimitrov, G.F. Korznikova, J.P. Dallas, A. Quivy, R.Z. Valiev, et al., Nanostruct. Mater. 11 (1999) 17–23. [7] T. Sawai, A. Hishinuma, J. Phys. Chem. Solids 66 (2005) 335–338. [8] N.R. Tao, K. Lu, Scr. Mater. 60 (2009) 1039–1043. [9] C.L. Chen, W. Lu, L.L. He, H.Q. Ye, J. Mater. Res. 22 (2007) 2416–2422.
Contents lists available at ScienceDirect
Materials Letters journal homepage: www.elsevier.com/locate/matlet
Surface nanocrystallization of Ti–45Al–7Nb–0.3W intermetallics induced by surface mechanical grinding treatment Kun Zhao, Yong Liu n, Tianhang Yao, Bin Liu, Yuehui He State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 24 July 2015 Received in revised form 30 November 2015 Accepted 4 December 2015 Available online 8 December 2015
Surface mechanical grinding treatment was conducted on a Ti–45Al–7Nb–0.3W intermetallics. The results show that the penetration depth plays an important role in the microstructural evolution. The γα2 phase transformation is observed at a penetration depth of 100 μm. Nanocrystallized α2 grains (20 nm) are found in the treated surface at a penetration depth of 200 μm. After the surface mechanical grinding treatment, the micro-hardness of the top surface increases significantly, which is 50% higher than that of the substrate. It is also found that the γ-α2 phase transformation is induced by the localized strain in the surface, while the mechanism for the nanocrystallization is proposed to be the twin–twin intersection. & 2015 Published by Elsevier B.V.
Keywords: Intermetallic alloys and compounds TiAl Surface mechanical grinding treatment Nanocrystallization Phase transformation
1. Introduction
2. Experimental
The surface nanocrystallization has been widely developed to synthesize a gradient grain structure, in order to improve the mechanical properties of metals and alloys. As a new method for the surface modification, the surface mechanical grinding treatment (SMGT) induces a very high rate shear deformation and a high strain gradient in the top surface of specimens. The shear strain rates can be 102–104 s 1 and the corresponding accumulative equivalent strains are 15–30 [1]. A gradient nanostructure can be formed after SMGT. Lu et al. has fabricated a gradient nanograined layer on pure copper, which has an extraordinary intrinsic tensile plasticity [2]. Besides, the grain refinement and the plastic deformation at large strains also result in disordering of the alloys and even the phase transformation [3,4]. It was found that γ-TiAl alloy powder was partially disordered and transformed to nanometer-sized hcp phase during the ball milling [5]. Similar results were reported in bulk TiAl alloys after the high pressure torsion (HPT), but the nanocrystallization mechanism was not clarified [6]. In this work, SMGT was conducted to form a nanostructured layer in a Ti–45Al–7Nb–0.3W alloy. The microstructural evolution and the mechanism of nanocrystallization were discussed.
An as-HIPed Ti–45Al–7Nb–0.3W (at%) alloy was used in the present study. The original SMGT was conducted on the circumference of cylinders [2]. A modified SMGT was performed on plates at room temperature. A hemispherical WC–Co tip (r ¼6 mm) was penetrated into the surface of specimen, and then moved on the plate from one side to the other. No material removal or contamination was found. The plastic deformation was uniform in the surface layer with a small roughness (Ra ¼0.35 μm). The preset penetration depth of the WC–Co tip into the sample per pass was 50 μm. The specimens were treated with the same operating parameters for 2 and 4 times. The total penetration depths (ap) were 50 μm, 100 μm and 200 μm, separately. The phase characterization of the treated surface was conducted on a Rigaku D/max 2550VB þ XRD equipment with a Cu radiation. A detailed description of the long range order (LRO) parameter and the mass fractions of phases can be referred to Ref. [6]. The mass fraction of γ phase is calculated using the matrixflushing method. For the determination of the mass fractions of phases, the intensity of the (200)γ was compared to the intensity of the (201)α2. The LRO parameters of the γ-TiAl phase were calculated from the intensities of the (100)γ superlattice and (200)γ peaks, with reference to the fully ordered alloys. The microstructural characterization of the specimens was carried out on a Leica DM4500P optical microscope and a transmission electron microscope (JEOL JEM 2010F). The TEM foils at different depths were obtained by mechanically polishing the
n
Corresponding author. E-mail address: [email protected] (Y. Liu).
http://dx.doi.org/10.1016/j.matlet.2015.12.025 0167-577X/& 2015 Published by Elsevier B.V.
60
K. Zhao et al. / Materials Letters 166 (2016) 59–62
sample from the untreated side to a thickness of 80 μm, and then by electro-polishing through the conventional twin-jet technique with an electrolyte consisting of 60% methanol, 35% normal butanol and 5% perchloric acid. Micro-hardness measurements were conducted on a BUEHLER OmniMet MHT automatic testing system with a load of 10 g and a loading time of 10–15 s.
3. Results Fig. 1(a) shows XRD patterns of the Ti–45Al–7Nb–0.3W alloy treated with different penetration depths. The substrate alloy exhibited two intermetallic phases, mainly γ phase and a small amount of α2 phase. The intensities of the superlattice peaks, such as (110)γ and (001)γ decreased continuously with the total penetration depth increasing. In the meantime, the intensities of the fundamental peaks, such as (002)γ and (200)γ, decreased continuously. At a total penetration depth of 200 μm, the peaks were broadened significantly, indicating a reduction in the grain size. The LRO parameter and the mass fraction of the γ phase are displayed in Table 1. The LRO parameter did not change at a total penetration depth of 50 μm. With the increase of the total penetration depth from 50 μm to 200 μm, the LRO parameter of γ phase decreased from 100% to 20%. When the penetration depth increased from 0 to 200 μm, the mass fraction of γ phase decreased from 74% to 14%. Therefore, the SMGT induced the disordering of γ phase and γ-α2 transformation in the surface. Fig. 1(b) indicates that the substrate alloy was a near-γ microstructure with an average grain size of 8 μm. No significant microstructural change occurred at a penetration depth of 50 μm
Table 1 Structure and properties of TiAl after SMGT. Penetration depth (ap) (μm)
LRO parameter of γ phase (%)
Mass fraction of γ phase (wt%)
0 50 100 200
100 100 94 20
74 48 44 14
(not shown). When the penetration depth increased to 100 μm, a bright layer of a thickness of 18 μm formed in the treated surface (Fig. 1(c)). According to the XRD patterns in Fig. 1(a), the bright layer consisted of α2 phase. It indicates that γ phase transformed to α2 phase in the treated surface. In Fig. 1(d), the thickness of the α2 rich layer increased to 26 μm at a penetration depth of 200 μm. In order to analyze the phase transformation and the grain refinement, the microstructures in different depths were observed by using TEM. The bright field (BF) images and corresponding selected area electron diffraction (SAED) patterns are shown in Fig. 2. Fig. 2(a) and (b) indicates that the surface layer consisted of nanocrystalline α2 phase (about 20 nm). In Fig. 2(c), the twin–twin intersection occurred at a depth of 80 μm from the surface. The density of twins was very high, and the spacing between neighboring twin boundaries was 30–90 nm. The corresponding SAED patterns in Fig. 2(d) indicated large misorientations. Twin–twin intersections were also found at a depth of 170 μm from the surface (Fig. 2(e)). Compared with the SAED patterns in Fig. 2(d), the spots in the Fig. 2(f) were not elongated, showing small misorientations. No twins or dislocations were found in γ grains of the
Fig. 1. (a) XRD patterns of as-SMGTed TiAl alloys and substrate, (b) microstructure of substrate, (c) microstructure of as-SMGTed specimens at penetration depth of 100 μm, and (d) microstructure of as-SMGTed specimens at penetration depth of 200 μm.
K. Zhao et al. / Materials Letters 166 (2016) 59–62
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Fig. 2. BF images and corresponding SAED patterns of as-SMGTed TiAl alloys (ap ¼ 200 μm) in different depths: (a),(b) the top surface, nanocrystallized α2 grains; (c),(d) in a depth of 80 μm, (c) shows twin–twin intersection, (d) shows large misorientations; (e),(f) in a depth of 170 μm, (e) shows twin–twin intersection, (f) shows small misorientations; and (g),(h) the substrate alloy.
substrate alloy (Fig. 2(g)). The SAED patterns in Fig. 2(h) indicated that there was a perfect lattice structure. The micro-hardness changed with the depth from the top surface (in Fig. 3). In the as-SMGTed alloy, the highest value of
hardness was 685HV0.01 in the top layer, which is 50% higher than that of the substrate (453HV0.01). In the subsurface layer, which consisted of nanometer-sized grains, the micro-hardness was around 600HV0.01. With the increase of the depth, the micro-
62
K. Zhao et al. / Materials Letters 166 (2016) 59–62
considered as a pseudo f.c.c structure. The standard α2 phase is a D019 structure, which is an ordered h.c.p structure. The structural change from f.c.c to h.c.p can be formed by the gliding of 1/6 < 112¯ ] Shockley partial dislocations on {111} planes [9]. When one 1/6 < 112¯ ] Shockley partial dislocation sweeps on a {111} plane, two layers of h.c.p structure will be converted from the f.c.c structure. After that, if another 1/6 < 112¯ ] Shockley partial dislocation glides on every other {111} plane, another two layers of h. c.p structure will be obtained. The γ-α2 phase transformation can be treated as a stress-induced transformation. The stacking sequence changes from cubic “ABCABC” to hexagonal “ABABAB”. The spacing between {111} close-packed planes in the γ phase is 0.231 nm. The distance between the nearest atoms in the γ phase is 0.281 nm. For the stress-induced α2, a ¼0.562 nm and c¼0.462 nm, which is close to the cell parameters for standard α2 phase (a ¼0.580 nm, c ¼0.465 nm). Fig. 3. Micro-hardness of as-SMGTed alloys (ap ¼ 200 μm), the hardness decreases with the distance from the surface.
5. Conclusions
hardness sharply decreased and became stable in the substrate.
4. Discussion In TiAl alloy, twins nucleate from the superposition of extended stacking faults on alternate {111} planes of γ phase. The addition of Nb lowers the intrinsic stacking fault energy, and promotes the formation of the mechanical twinning in high-Nb TiAl alloys. Due to the difficult deformability of α2 phase, during SMGT, the plastic deformation mainly occurs in γ phase, and 1/6 < 112¯ ]{111} twining is the deformation mechanism [7]. When twinning takes place, the γ grain is divided into the twin/matrix lamellar structure with a specific orientation. The deformation twinning can be regarded as the first step for the nanocrystallization [8]. The second step is the refinement of twin/ matrix lamellae into equiaxed nanometer-sized grains through the twin–twin intersection [8]. When the driving force for the deformation twinning is large enough to overcome the barriers of the encountered twin boundaries, twin–twin intersection occurs. Twin–twin intersections can divide twin/matrix lamellae into rhombic blocks. The intersected volume undergoes a rotation to change its orientation. When the shear strain increases, the multiplication of deformation twinning leads to more twin–twin intersections, refining the structure into equiaxed grains with large misorientations in nanometer scale [8]. The stress induces the phase transformations between γ and α2 phase in TiAl alloys [9]. Moreover, γ-α2 transformations is possible in the twin–twin intersection region of deformed TiAl alloys [9]. γ phase is L10 structure (a ¼0.398 nm and c¼ 0.405 nm). The c/ a ratio of γ phase (1.03) is close to 1, and thus γ phase can be
(1) Through SMGT, a gradient grain structure can be formed in the surface of TiAl alloy. In the top surface, nanocrystallized α2 grains (∼20 nm) can be obtained. Both the γ-α2 phase transformation and twin–twin intersections occur with the increase of the penetration depth. (2) After SMGT, the micro-hardness of the top surface increases significantly, which is 50% higher than that of the substrate. (3) The γ-α2 phase transformation is induced by the localized strain in the surface, while the mechanism for the nanocrystallization is proposed to be the twin–twin intersection.
Acknowledgments The work was financially supported by the Project of Innovation-driven Plan in Central South University (2015CX004) and the National Natural Science Foundation of China (51301203).
Reference [1] [2] [3] [4] [5] [6]
X.C. Liu, H.W. Zhang, K. Lu, Science 342 (2013) 337–340. T.H. Fang, W.L. Li, N.R. Tao, K. Lu, Science 331 (2011) 1587–1590. H.W. Zhang, Z.K. Hei, G. Liu, J. Lu, K. Lu, Acta Mater. 51 (2003) 1871–1881. X. Wu, N. Tao, Y. Hong, G. Liu, B. Xu, J. Lu, K. Lu, Acta Mater. 53 (2005) 681–691. T. Klassen, M. Oehring, R. Bormann, Acta Mater. 45 (1997) 3935–3948. A.V. Korznikov, O. Dimitrov, G.F. Korznikova, J.P. Dallas, A. Quivy, R.Z. Valiev, et al., Nanostruct. Mater. 11 (1999) 17–23. [7] T. Sawai, A. Hishinuma, J. Phys. Chem. Solids 66 (2005) 335–338. [8] N.R. Tao, K. Lu, Scr. Mater. 60 (2009) 1039–1043. [9] C.L. Chen, W. Lu, L.L. He, H.Q. Ye, J. Mater. Res. 22 (2007) 2416–2422.