- Email: [email protected]
nanocrystal composite with high ductility and strain-hardening
Materials Science & Engineering A 560 (2013) 339–342
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Ti-based amorphous/nanocrystal composite with high ductility and strain-hardening D.K. Yang a,n, J.T. Wang a, D. Fabijanic a, P. Cizek a, B.S. Li a, J.Z. Lu b, P.D. Hodgson a a b
Institute for Frontier Materials, Deakin University, Waurn Ponds, Victoria 3217, Australia School of Mechanical Engineering, Jiangsu University, Zhenjiang 212013, PR China
a r t i c l e i n f o
abstract
Article history: Received 21 August 2012 Received in revised form 17 September 2012 Accepted 20 September 2012 Available online 27 September 2012
The mechanical behavior of a Ti-based amorphous/nanocrystal composite was studied by microcompression testing. The results show that this composite has high ductility and strain-hardening. The deformability improvement is attributed to the nucleation, multiplication and interaction of localized shear flows and shear bands during deformation. Microscopic observations indicate that compressive loading facilitates nanocrystal precipitation in the amorphous matrix, which interact with crack propagation through the amorphous matrix, resulting in remarkable plasticity. & 2012 Elsevier B.V. All rights reserved.
Keywords: Amorphous materials Nanocrystal Microcompression Composite Mechanical properties
1. Introduction Amorphous metallic alloys, or metallic glasses (MGs) represent a new class of advanced materials and have attracted significant interest for structural applications [1,2]. In comparison with common engineering materials, most MGs exhibit high strength and substantial fracture toughness. On the other hand, MGs have limited plasticity and fail in an apparently brittle manner at room temperature [2,3]. The macroscopic plastic deformation of MGs has been known to be accomplished through the formation and evolution of shear bands, in which a high amount of plastic flow is localized within a rather narrow region [4]. Finally, failure occurs typically along one dominant shear band that cuts across the sample at a direction of maximum resolved shear stress due to highly localized shearing and thermal softening [4–6]. Ductility improvement of MGs largely depends on the suppression of the localized strain softening caused by shear bands. Therefore, methods for improving the plasticity of MGs revolve around the question of how to activate multiple shear bands and how to suppress their highly catastrophic propagation [7,8]. In order to improve the plastic deformability of MGs, a straightforward strategy is to introduce ductile crystalline phases into bulk MGs to synthesize the MG matrix composites, in which the potential rapid catastrophic failure along a single shear band is blocked by the interaction of the shear bands with ductile crystalline phases [4,5,9].
In our previous work, we applied cryorolling in conjunction with surface mechanical attrition treatment (SMAT) [10] to commercialpurity Ti as a strategy to engineer a multilayered hierarchical structure (MHS) with a graded microstructural transition between successive layers [11]. The microstructural study has revealed that the outer layer of MHS Ti (thickness 30 mm) is composed of an apparent amorphous phase and crystallites. After compressive loading of MHS Ti sample (5 5 10 mm3), multiple shear bands along various directions were observed in the amorphous/nanocrystallite (A/NC) layer and the intersections and interactions of shear bands in three dimensions were numerous. This hints A/NC layer of MHS Ti may possess extended ductility. MGs have great potential for applications in micro-electromechanical systems because of their high strength and hardness. Thus, it is of great interest to study the properties of microscale specimen. The micro-compression test developed by Uchic and Dimiduk [12] provides opportunity to understand the deformation mechanisms of materials at the microscale. In this study, we investigate the mechanical properties of Ti-based A/NC composite produced by cryogenic rolling and subsequent SMAT using microcompression testing. SEM analysis of the deformed micropillar and TEM analysis of the deformed microstructure were used to provide fundamental insights into the mechanism underlying the high ductility in Ti-based A/NC composite.
2. Experimental procedures n
Corresponding author. Tel.: þ61 3 5227 1283; fax: þ 61 3 5227 1103. E-mail addresses: [email protected], [email protected] (D.K. Yang). 0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.09.076
A commercial Ti plate (Grade 2) with a mean grain size of 60 mm was used in this study. Table 1 lists the chemical
340
D.K. Yang et al. / Materials Science & Engineering A 560 (2013) 339–342
compositions of the as received Ti. The as-received Ti was cryogenic rolled from 36 to 6 mm in thickness with a reduction of 2 mm per pass. Cryorolling was performed by immersing the Ti plate into liquid nitrogen for 15 min before each rolling pass. After cryorolling, the workpiece was cut along the rolled direction (RD) and polished to a rectangular bar with dimensions of 5 5 90 mm3. Subsequently, four lateral surfaces of the rectangular bar were subjected to SMAT in sequence. The SMAT process was performed in a low vacuum using 8 mm diameter hardened stainless steel balls at a vibration frequency of 50 Hz for 60 min. To avoid a temperature rise on the workpiece surface from mechanical energy, every 10 min of operation was followed by a dwell period of 15 min. The produced Ti has a multilayered hierarchical structure consisting of three distinct layers: an outer amorphous/nanocrystallite (A/NC) layer ( 30 mm thick), an underneath nanograined (NG) layer ( 60 mm thick) and a UFG core [11]. The cross-section of MHS Ti was mechanically polished to produce a parallel-sided and flat surface, which was subsequently polished to a final sub-micron finish to remove any residual damage from mechanical polishing. Micropillars with diameter of 2 mm and aspect ratio of 2:1 were fabricated in the center region of A/NC layer using focused ion beam (FIB) in a dual beam SEM (FEI Quanta 3D). Firstly, a Ga þ beam operated at 30 kV and 5 nA was faced to the A/NC layer to mill a ring-shaped crater, creating a pillar of 5 mm in diameter. Then, at the same voltage but reduced current (50 pA), the premilled pillar was polished Table 1 Chemical compositions of the as-received Ti. N
C
H
Fe
O
Ti
0.0030
0.0130
0.0013
0.0300
0.1250
Balance
Fig. 1. Elemental depth profiling of the MHS Ti and as-received Ti.
to the desired diameter and height. The compression tests were performed using a flat punch indenter of 5 mm diameter in a nanoindenter (UMIS/CSIRO). The flat punch was driven down in a force-controlled mode to compress the pillars. The 4 mm high pillars were compressed at a displacement rate of 1 nm/s, producing a strain rate of 2.5 10 4 s 1. Microstructural observations were conducted using a field-emission gun scanning electron microscope (SEM, Zeiss Supra 55VP) operated at 5 kV and a transmission electron microscope (TEM, Jeol JEM 2100) operated at 200 kV.
3. Results and discussion Elemental depth profiling was collected using a Leco GDS850A glow discharge optical emission spectrometer (GDOES). The GDOES results presented in Fig. 1 did not show substantial elemental composition variation between the as-received Ti and MHS Ti in the depth up to 30 mm. Therefore, the MHS process did not alter the surface chemical composition. TEM analysis has revealed that the 30 mm thick top layer was composed of a bright phase matrix and a discreet darker nanostructure (Fig. 2a). The upper right inset taken from the bright matrix region exhibited a broad diffuse halo in a selected-area diffraction pattern (SAD), which is typically a fully amorphous phase. The SAD pattern in the lower left inset taken from the interface between the bright and dark phases clearly demonstrates the presence of a nanocrystalline phase together with the amorphous phase. Highresolution TEM images (Fig. 2b and c), taken from the marked areas in Fig. 2a further confirm the respective amorphous and nanocrystalline character of the phases. Quantitative analysis of the TEM images showed that the fraction of the amorphous phase is 78%. The size of the crystallites was less than 20 nm and the crystallites congregate into islands surrounded by the amorphous matrix. The formation of amorphous phase by the cryorolling and subsequent SMAT process appears to bear some similarity with that observed in other SPD processes. Presumably, it is caused by the crystal structure collapse, which was induced by the high internal lattice strain and the dissolution of small atoms into the large lattice [13–15]. Fig. 3 shows the typical engineering stress–strain curves for the Ti-based A/NC composite under micro-compressive loading at a strain rate of 2.5 10 4 s 1. In present study, the microcompression specimen axially was compressed to a strain of 28%. The composite exhibited an elastic strain limit up to 2.5%, followed by plastic deformation. A strain hardening behavior was evident when deformed beyond yielding to a total compressive ductility up to 18% strain. When the strain is 420%, the sharp increase of the stress is most possibly caused by the friction occurring at the contacts surfaces of the specimen and the indenter [8]. In contrast to most amorphous alloys where the
b
c 100nm
5nm
2nm
Fig. 2. (a) TEM bright-field (BF) image of the microstructure 20 mm below the top surface, revealing that the top layer is composed of an amorphous phase (bright region) and nanocrystallites (dark region). The upper right inset is the SAD pattern of the amorphous phase and the lower left inset is the SAD pattern of the nanocrystallite; (b) and (c) high-resolution (HR) TEM images obtained from the regions marked in (a).
D.K. Yang et al. / Materials Science & Engineering A 560 (2013) 339–342
elastic–plastic transition is abrupt and the deformation is characterized by strain ‘pop-in’ events under micro-compression [16–20], the current A/NC composite displays a continuing elastic–plastic transition and gradual strain. SEM images representing the same pillar before and after deformation are shown in the inset of Fig. 3. The increase in lateral dimensions and reduction in height of the pillar suggest a high deformability of this A/NC composite. This behavior is very similar to some highly ductile Zr- or Al-based amorphous materials, e.g. [8,20]. Due to the continuous elastic–plastic deformation transition, we take the 0.2% offset yield strength as the onset of plasticity. The 0.2% offset
Bulk MHS Ti Loading stop
Fig. 3. Engineering stress–strain curves for microscale compression specimen of Ti-based A/NC composite. Inset shows SEM images of a micropillar before and after compression to strain of 28%.
a
341
yield strength of A/NC composite micropillar is 1.8 GPa, an increase of 5.9% over the Ti-based bulk MGs ( 1.7 GPa) [6]. The slight increase in the strength may be caused by the geometrical taper of the pillar, a strengthening artefact caused by the FIB damaged layer or a decrease in the defect population in the microscale specimen [17–20]. The elastic modulus was determined using the method developed in [16] and then corrected to eliminate the effect of the base compliance [21]. The measured elastic modulus is 97.3 GPa, which is comparable to the quoted modulus of bulk Ti based MG specimens ( 90 GPa) [6]. This demonstrates that the misalignment effect in the current study is negligible. Compression loading was stopped at the yielding, strainhardening and stress-drop deformation stages to understand how plastic strains were accommodated and the origin of the strain-hardening during deformation (Fig. 3). Typical surface morphologies of these deformed micropillars are shown in Fig. 4. Uniform localized flows with 10–50 nm wide formed approximately vertical to the compression direction on the surface of the pillar immediately after yielding ( 2% strain) (Fig. 4a). Localized flow increased with further deformation into the strain hardening region (Fig. 4b). These localized flows combined to form a rough surface morphology to support the increasing deformation. This indicates that a homogenous nucleation and distribution of a high density of localized flows accommodate the applied strain rather than accumulation and cause catastrophic failure along one dominant shear band. When the strain increased to 18%, a large number of shear bands with space of 30–80 nm distributed on the surface of the specimen (Fig. 4c). The multiplication, interaction and intersection of the shear bands create a
b
c
200nm
200nm
200nm
Fig. 4. Surface morphology change of micropillar at different deformation stages obtained by high resolution SEM: (a) right after yielding ( 2% strain), localized flows originated on the surface; (b) ‘strain-hardening’ stage ( 8% strain), localized flows linked to form a rough surface morphology and (c) further strain to 18%, the multiplication, interaction and intersection of the shear bands on the surface.
1nm
10nm
10nm
Fig. 5. (a) TEM BF images of the microstructures of the A/NC layer after compression deformation to 5% strain, showing abundant nanocrystallites distributed homogeneously in the amorphous matrix. The upper right inset is the SAD pattern corresponding to the area and the low left inset is the HRTEM image recorded from the black speck, confirming the presence of nanocrystallites and (b) BF TEM images showing the microstructure of the A/NC layer after strained to 18%. Small nanograins are clearly observed with a typical size of 5–20 nm. The inset shows the corresponding SAD pattern.
342
D.K. Yang et al. / Materials Science & Engineering A 560 (2013) 339–342
microcrack extension was effectively arrested when it propagated through the amorphous matrix containing the precipitated nanocrystallites and the tip of the suppressed microcrack propagated in a jagged manner inside the amorphous matrix. The precipitated nanocrystallites thus act as obstacles to the direct crack propagation during loading.
2nm 4. Conclusions
10nm Fig. 6. TEM observations showing that the microcrack extension becomes arrested when it propagates across the amorphous matrix containing the precipitated nanocrystallites and the tip of the suppressed microcrack propagates in a jagged pattern. The upper right inset is the SAD pattern corresponding to the area and the lower left inset is the HRTEM image recorded from a black speck.
local restriction effect, which arrests the further propagation of these bands, thus restricting the potential catastrophic failure, similar to that shown in [5,8] for bulk MGs. To further explore the deformation micro-mechanisms, bulk MHS Ti rectangular specimens with dimensions of 5 5 10 mm3 were compressed at a strain rate of 2.5 10 4 s 1. The compression testing result for bulk MHS Ti was given in Fig. 3. It can be seen that the yield strength of A/NC composite is much higher than bulk MHT Ti ( 1.1 GPa) which is composed by A/NC layer, NG layer and UFG core. This reflects that A/NC structure has higher strength than NG or UFG structure. The A/NC layer of the tested specimen was investigated using TEM. Bright field images of the A/NC layer microstructure after 5% compressive strain show abundant dark regions with an average size of 5 nm distributed homogeneously in the amorphous matrix (Fig. 5a). The HR TEM image and the SAD pattern (inset in Fig. 5a) of the corresponding area verify the crystalline nature of the dark regions. The homogenous distribution of the precipitated crystallites were absent from the undeformed specimen (Fig. 2a), indicating that this is not an artefact from TEM sample preparation. Small nanograins are clearly observed with a typical size of 5–20 nm when further strained to 18% (Fig. 5b). The corresponding SAD pattern shows the typical character of nanograined polycrystalline materials. It appears that compressive deformation induced the formation of nanocrystallites, similar to deformation-induced nanocrystallite formation in amorphous matrices in other studies [4,7,22–26], with [4,8,26] showing that nanocrystal precipitation not only promotes the nucleation of shear bands but also alters shear band propagating directions and assists their branching and multiplication. The nanocrystal precipitations shown in Fig. 5 should play a similar role in hindering, branching, and deflecting the propagation of shear bands, resulting in intersection, interaction and multiplication of shear bands, as shown in Fig. 4. TEM observations have provided evidence that the A/NC layer (after strain to 35%) is able to disperse the strain localization upon compression by the interaction of the precipitated nanocrystallites with the propagation of the shear bands and microcracks. As shown in Fig. 6, the
In summary, a Ti-based A/NC composite was produced by sequentially combining cryorolling and SMAT. This composite showed high ductility and strain-hardening, which is attributed to the homogeneous nucleation of shear localized flows and continuous multiplication and interaction of shear bands during deformation. Microscopic observation of the deformed composite shows the compressive loading facilities nanocrystallite precipitation in the amorphous matrix. The precipitates appear to arrest and deflect crack propagation through the amorphous matrix, contributing to the high plasticity.
Acknowledgments DKY acknowledges the financial support through Alfred Deakin Postdoctoral Fellowship. PH acknowledges the support from Australian Research Council through Australia Laureate Fellowship. References [1] M.F. Ashby, A.L. Greer, Scripta Mater. 54 (2006) 321. [2] D.C. Hofmann, J.Y. Suh, A. Wiest, G. Duan, M.L. Lind, M.D. Demetriou, W.L. Johnson, Nature 451 (2008) 1085. [3] L. Tian, Y.Q. Cheng, Z.W. Shan, C.C. Wang, X.D. Han, J. Sun, E. Ma, Nature Commun.. 3 (2012) 609. [4] M.W. Chen, A. Inoue, W. Zhang, T. Sakurai, Phys. Rev. Lett. 96 (2006) 245502. [5] H.K. Kim, K.B. Lee, J.C. Lee, Mater. Sci. Eng. A 552 (2012) 399. [6] X.J. Gu, S.J. Poon, G.J. Shiflet, J.J. Lewandowski, Acta Mater. 58 (2010) 1708. [7] J.C. Lee, Y.C. Kim, J.P. Ahn, H.S. Kim, Acta Mater. 53 (2005) 129. [8] Y.H. Liu, G. Wang, R.J. Wang, D.Q. Zhao, M.X. Pan, W.H. Wang, Science 315 (2007) 1385. [9] H.B. Yu, J. Hu, X.X. Xia, B.A. Sun, X.X. Li, W.H. Wang, H.Y. Bai, Scripta Mater. 61 (2009) 640. [10] K. Lu, J. Lu, Mater. Sci. Eng. A 375–377 (2004) 38. [11] D.K. Yang, P. Cizek, D. Fabijanic, J.T. Wang, B.S. Li, P.D. Hodgson, Work Hardening in Ultrafine-Grained Titanium: Multilayering and Grading, Submitted for publication. [12] M.D. Uchic, D.M. Dimiduk, Mater. Sci. Eng. A 400–401 (2005) 268. [13] F. Sun, P. Rojas, A. Zuniga, E.J. Lavernia, Mater. Sci. Eng. A 430 (2006) 90. [14] A. Sagel, H. Sieber, H.J. Fecht, J.H. Ferepezko, Acta Mater. 46 (1998) 4233. [15] S. Ohsaki, S. Kato, N. Tsuji, T. Ohkubo, K. Hono, Acta Mater. 55 (2007) 2885. [16] C.J. Lee, J.C. Huang, T.G. Nieh, Appl. Phys. Lett. 91 (2007) 161913. [17] Y.H. Lee, Y.T. Cheng, H.S. Chou, H.M. Chen, X.H. Du, C.I. Chang, J.C. Huang, S.R. Jian, J.S.C. Jang, T.G. Nieh, Scripta Mater. 58 (2008) 890. [18] B.E. Schuster, Q. Wei, M.H. Ervin, S.O. Hruszkewycz, M.K. Miller, T.C. Hufnagel, K.T. Ramesh, Scripta Mater. 57 (2007) 517. [19] L.C. Zhang, F. Jiang, Y.L. Zhao, J.F. Zhang, L. He, J. Sun, Mater. Sci. Eng. A 527 (2010) 4122. [20] A. Bharathula, S.W. Lee, W.J. Wright, K.M. Flores, Acta Mater. 58 (2010) 5789. [21] H. Zhang, B.E. Schuster, Q. Wei, K.T. Ramesh, Scripta Mater. 54 (2006) 181. [22] H. Chen, Y. He, G.J. Shiflet, S.J. Poon, Nature 367 (1994) 541. [23] W.H. Jiang, M. Atzmon, Acta Mater. 51 (2003) 4095. [24] J.J. Kim, Y. Choi, S. Suresh, A.S. Argon, Science 295 (2002) 654. [25] J.C. Lee, Y.C. Kim, J.P. Ahn, H.S. Kim, S.H. Lee, B.J. Lee, Acta Mater. 52 (2003) 1525. [26] S.W. Lee, M.Y. Huh, E. Fleury, J.C. Lee, Acta Mater. 54 (2006) 349.
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Ti-based amorphous/nanocrystal composite with high ductility and strain-hardening D.K. Yang a,n, J.T. Wang a, D. Fabijanic a, P. Cizek a, B.S. Li a, J.Z. Lu b, P.D. Hodgson a a b
Institute for Frontier Materials, Deakin University, Waurn Ponds, Victoria 3217, Australia School of Mechanical Engineering, Jiangsu University, Zhenjiang 212013, PR China
a r t i c l e i n f o
abstract
Article history: Received 21 August 2012 Received in revised form 17 September 2012 Accepted 20 September 2012 Available online 27 September 2012
The mechanical behavior of a Ti-based amorphous/nanocrystal composite was studied by microcompression testing. The results show that this composite has high ductility and strain-hardening. The deformability improvement is attributed to the nucleation, multiplication and interaction of localized shear flows and shear bands during deformation. Microscopic observations indicate that compressive loading facilitates nanocrystal precipitation in the amorphous matrix, which interact with crack propagation through the amorphous matrix, resulting in remarkable plasticity. & 2012 Elsevier B.V. All rights reserved.
Keywords: Amorphous materials Nanocrystal Microcompression Composite Mechanical properties
1. Introduction Amorphous metallic alloys, or metallic glasses (MGs) represent a new class of advanced materials and have attracted significant interest for structural applications [1,2]. In comparison with common engineering materials, most MGs exhibit high strength and substantial fracture toughness. On the other hand, MGs have limited plasticity and fail in an apparently brittle manner at room temperature [2,3]. The macroscopic plastic deformation of MGs has been known to be accomplished through the formation and evolution of shear bands, in which a high amount of plastic flow is localized within a rather narrow region [4]. Finally, failure occurs typically along one dominant shear band that cuts across the sample at a direction of maximum resolved shear stress due to highly localized shearing and thermal softening [4–6]. Ductility improvement of MGs largely depends on the suppression of the localized strain softening caused by shear bands. Therefore, methods for improving the plasticity of MGs revolve around the question of how to activate multiple shear bands and how to suppress their highly catastrophic propagation [7,8]. In order to improve the plastic deformability of MGs, a straightforward strategy is to introduce ductile crystalline phases into bulk MGs to synthesize the MG matrix composites, in which the potential rapid catastrophic failure along a single shear band is blocked by the interaction of the shear bands with ductile crystalline phases [4,5,9].
In our previous work, we applied cryorolling in conjunction with surface mechanical attrition treatment (SMAT) [10] to commercialpurity Ti as a strategy to engineer a multilayered hierarchical structure (MHS) with a graded microstructural transition between successive layers [11]. The microstructural study has revealed that the outer layer of MHS Ti (thickness 30 mm) is composed of an apparent amorphous phase and crystallites. After compressive loading of MHS Ti sample (5 5 10 mm3), multiple shear bands along various directions were observed in the amorphous/nanocrystallite (A/NC) layer and the intersections and interactions of shear bands in three dimensions were numerous. This hints A/NC layer of MHS Ti may possess extended ductility. MGs have great potential for applications in micro-electromechanical systems because of their high strength and hardness. Thus, it is of great interest to study the properties of microscale specimen. The micro-compression test developed by Uchic and Dimiduk [12] provides opportunity to understand the deformation mechanisms of materials at the microscale. In this study, we investigate the mechanical properties of Ti-based A/NC composite produced by cryogenic rolling and subsequent SMAT using microcompression testing. SEM analysis of the deformed micropillar and TEM analysis of the deformed microstructure were used to provide fundamental insights into the mechanism underlying the high ductility in Ti-based A/NC composite.
2. Experimental procedures n
Corresponding author. Tel.: þ61 3 5227 1283; fax: þ 61 3 5227 1103. E-mail addresses: [email protected], [email protected] (D.K. Yang). 0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.09.076
A commercial Ti plate (Grade 2) with a mean grain size of 60 mm was used in this study. Table 1 lists the chemical
340
D.K. Yang et al. / Materials Science & Engineering A 560 (2013) 339–342
compositions of the as received Ti. The as-received Ti was cryogenic rolled from 36 to 6 mm in thickness with a reduction of 2 mm per pass. Cryorolling was performed by immersing the Ti plate into liquid nitrogen for 15 min before each rolling pass. After cryorolling, the workpiece was cut along the rolled direction (RD) and polished to a rectangular bar with dimensions of 5 5 90 mm3. Subsequently, four lateral surfaces of the rectangular bar were subjected to SMAT in sequence. The SMAT process was performed in a low vacuum using 8 mm diameter hardened stainless steel balls at a vibration frequency of 50 Hz for 60 min. To avoid a temperature rise on the workpiece surface from mechanical energy, every 10 min of operation was followed by a dwell period of 15 min. The produced Ti has a multilayered hierarchical structure consisting of three distinct layers: an outer amorphous/nanocrystallite (A/NC) layer ( 30 mm thick), an underneath nanograined (NG) layer ( 60 mm thick) and a UFG core [11]. The cross-section of MHS Ti was mechanically polished to produce a parallel-sided and flat surface, which was subsequently polished to a final sub-micron finish to remove any residual damage from mechanical polishing. Micropillars with diameter of 2 mm and aspect ratio of 2:1 were fabricated in the center region of A/NC layer using focused ion beam (FIB) in a dual beam SEM (FEI Quanta 3D). Firstly, a Ga þ beam operated at 30 kV and 5 nA was faced to the A/NC layer to mill a ring-shaped crater, creating a pillar of 5 mm in diameter. Then, at the same voltage but reduced current (50 pA), the premilled pillar was polished Table 1 Chemical compositions of the as-received Ti. N
C
H
Fe
O
Ti
0.0030
0.0130
0.0013
0.0300
0.1250
Balance
Fig. 1. Elemental depth profiling of the MHS Ti and as-received Ti.
to the desired diameter and height. The compression tests were performed using a flat punch indenter of 5 mm diameter in a nanoindenter (UMIS/CSIRO). The flat punch was driven down in a force-controlled mode to compress the pillars. The 4 mm high pillars were compressed at a displacement rate of 1 nm/s, producing a strain rate of 2.5 10 4 s 1. Microstructural observations were conducted using a field-emission gun scanning electron microscope (SEM, Zeiss Supra 55VP) operated at 5 kV and a transmission electron microscope (TEM, Jeol JEM 2100) operated at 200 kV.
3. Results and discussion Elemental depth profiling was collected using a Leco GDS850A glow discharge optical emission spectrometer (GDOES). The GDOES results presented in Fig. 1 did not show substantial elemental composition variation between the as-received Ti and MHS Ti in the depth up to 30 mm. Therefore, the MHS process did not alter the surface chemical composition. TEM analysis has revealed that the 30 mm thick top layer was composed of a bright phase matrix and a discreet darker nanostructure (Fig. 2a). The upper right inset taken from the bright matrix region exhibited a broad diffuse halo in a selected-area diffraction pattern (SAD), which is typically a fully amorphous phase. The SAD pattern in the lower left inset taken from the interface between the bright and dark phases clearly demonstrates the presence of a nanocrystalline phase together with the amorphous phase. Highresolution TEM images (Fig. 2b and c), taken from the marked areas in Fig. 2a further confirm the respective amorphous and nanocrystalline character of the phases. Quantitative analysis of the TEM images showed that the fraction of the amorphous phase is 78%. The size of the crystallites was less than 20 nm and the crystallites congregate into islands surrounded by the amorphous matrix. The formation of amorphous phase by the cryorolling and subsequent SMAT process appears to bear some similarity with that observed in other SPD processes. Presumably, it is caused by the crystal structure collapse, which was induced by the high internal lattice strain and the dissolution of small atoms into the large lattice [13–15]. Fig. 3 shows the typical engineering stress–strain curves for the Ti-based A/NC composite under micro-compressive loading at a strain rate of 2.5 10 4 s 1. In present study, the microcompression specimen axially was compressed to a strain of 28%. The composite exhibited an elastic strain limit up to 2.5%, followed by plastic deformation. A strain hardening behavior was evident when deformed beyond yielding to a total compressive ductility up to 18% strain. When the strain is 420%, the sharp increase of the stress is most possibly caused by the friction occurring at the contacts surfaces of the specimen and the indenter [8]. In contrast to most amorphous alloys where the
b
c 100nm
5nm
2nm
Fig. 2. (a) TEM bright-field (BF) image of the microstructure 20 mm below the top surface, revealing that the top layer is composed of an amorphous phase (bright region) and nanocrystallites (dark region). The upper right inset is the SAD pattern of the amorphous phase and the lower left inset is the SAD pattern of the nanocrystallite; (b) and (c) high-resolution (HR) TEM images obtained from the regions marked in (a).
D.K. Yang et al. / Materials Science & Engineering A 560 (2013) 339–342
elastic–plastic transition is abrupt and the deformation is characterized by strain ‘pop-in’ events under micro-compression [16–20], the current A/NC composite displays a continuing elastic–plastic transition and gradual strain. SEM images representing the same pillar before and after deformation are shown in the inset of Fig. 3. The increase in lateral dimensions and reduction in height of the pillar suggest a high deformability of this A/NC composite. This behavior is very similar to some highly ductile Zr- or Al-based amorphous materials, e.g. [8,20]. Due to the continuous elastic–plastic deformation transition, we take the 0.2% offset yield strength as the onset of plasticity. The 0.2% offset
Bulk MHS Ti Loading stop
Fig. 3. Engineering stress–strain curves for microscale compression specimen of Ti-based A/NC composite. Inset shows SEM images of a micropillar before and after compression to strain of 28%.
a
341
yield strength of A/NC composite micropillar is 1.8 GPa, an increase of 5.9% over the Ti-based bulk MGs ( 1.7 GPa) [6]. The slight increase in the strength may be caused by the geometrical taper of the pillar, a strengthening artefact caused by the FIB damaged layer or a decrease in the defect population in the microscale specimen [17–20]. The elastic modulus was determined using the method developed in [16] and then corrected to eliminate the effect of the base compliance [21]. The measured elastic modulus is 97.3 GPa, which is comparable to the quoted modulus of bulk Ti based MG specimens ( 90 GPa) [6]. This demonstrates that the misalignment effect in the current study is negligible. Compression loading was stopped at the yielding, strainhardening and stress-drop deformation stages to understand how plastic strains were accommodated and the origin of the strain-hardening during deformation (Fig. 3). Typical surface morphologies of these deformed micropillars are shown in Fig. 4. Uniform localized flows with 10–50 nm wide formed approximately vertical to the compression direction on the surface of the pillar immediately after yielding ( 2% strain) (Fig. 4a). Localized flow increased with further deformation into the strain hardening region (Fig. 4b). These localized flows combined to form a rough surface morphology to support the increasing deformation. This indicates that a homogenous nucleation and distribution of a high density of localized flows accommodate the applied strain rather than accumulation and cause catastrophic failure along one dominant shear band. When the strain increased to 18%, a large number of shear bands with space of 30–80 nm distributed on the surface of the specimen (Fig. 4c). The multiplication, interaction and intersection of the shear bands create a
b
c
200nm
200nm
200nm
Fig. 4. Surface morphology change of micropillar at different deformation stages obtained by high resolution SEM: (a) right after yielding ( 2% strain), localized flows originated on the surface; (b) ‘strain-hardening’ stage ( 8% strain), localized flows linked to form a rough surface morphology and (c) further strain to 18%, the multiplication, interaction and intersection of the shear bands on the surface.
1nm
10nm
10nm
Fig. 5. (a) TEM BF images of the microstructures of the A/NC layer after compression deformation to 5% strain, showing abundant nanocrystallites distributed homogeneously in the amorphous matrix. The upper right inset is the SAD pattern corresponding to the area and the low left inset is the HRTEM image recorded from the black speck, confirming the presence of nanocrystallites and (b) BF TEM images showing the microstructure of the A/NC layer after strained to 18%. Small nanograins are clearly observed with a typical size of 5–20 nm. The inset shows the corresponding SAD pattern.
342
D.K. Yang et al. / Materials Science & Engineering A 560 (2013) 339–342
microcrack extension was effectively arrested when it propagated through the amorphous matrix containing the precipitated nanocrystallites and the tip of the suppressed microcrack propagated in a jagged manner inside the amorphous matrix. The precipitated nanocrystallites thus act as obstacles to the direct crack propagation during loading.
2nm 4. Conclusions
10nm Fig. 6. TEM observations showing that the microcrack extension becomes arrested when it propagates across the amorphous matrix containing the precipitated nanocrystallites and the tip of the suppressed microcrack propagates in a jagged pattern. The upper right inset is the SAD pattern corresponding to the area and the lower left inset is the HRTEM image recorded from a black speck.
local restriction effect, which arrests the further propagation of these bands, thus restricting the potential catastrophic failure, similar to that shown in [5,8] for bulk MGs. To further explore the deformation micro-mechanisms, bulk MHS Ti rectangular specimens with dimensions of 5 5 10 mm3 were compressed at a strain rate of 2.5 10 4 s 1. The compression testing result for bulk MHS Ti was given in Fig. 3. It can be seen that the yield strength of A/NC composite is much higher than bulk MHT Ti ( 1.1 GPa) which is composed by A/NC layer, NG layer and UFG core. This reflects that A/NC structure has higher strength than NG or UFG structure. The A/NC layer of the tested specimen was investigated using TEM. Bright field images of the A/NC layer microstructure after 5% compressive strain show abundant dark regions with an average size of 5 nm distributed homogeneously in the amorphous matrix (Fig. 5a). The HR TEM image and the SAD pattern (inset in Fig. 5a) of the corresponding area verify the crystalline nature of the dark regions. The homogenous distribution of the precipitated crystallites were absent from the undeformed specimen (Fig. 2a), indicating that this is not an artefact from TEM sample preparation. Small nanograins are clearly observed with a typical size of 5–20 nm when further strained to 18% (Fig. 5b). The corresponding SAD pattern shows the typical character of nanograined polycrystalline materials. It appears that compressive deformation induced the formation of nanocrystallites, similar to deformation-induced nanocrystallite formation in amorphous matrices in other studies [4,7,22–26], with [4,8,26] showing that nanocrystal precipitation not only promotes the nucleation of shear bands but also alters shear band propagating directions and assists their branching and multiplication. The nanocrystal precipitations shown in Fig. 5 should play a similar role in hindering, branching, and deflecting the propagation of shear bands, resulting in intersection, interaction and multiplication of shear bands, as shown in Fig. 4. TEM observations have provided evidence that the A/NC layer (after strain to 35%) is able to disperse the strain localization upon compression by the interaction of the precipitated nanocrystallites with the propagation of the shear bands and microcracks. As shown in Fig. 6, the
In summary, a Ti-based A/NC composite was produced by sequentially combining cryorolling and SMAT. This composite showed high ductility and strain-hardening, which is attributed to the homogeneous nucleation of shear localized flows and continuous multiplication and interaction of shear bands during deformation. Microscopic observation of the deformed composite shows the compressive loading facilities nanocrystallite precipitation in the amorphous matrix. The precipitates appear to arrest and deflect crack propagation through the amorphous matrix, contributing to the high plasticity.
Acknowledgments DKY acknowledges the financial support through Alfred Deakin Postdoctoral Fellowship. PH acknowledges the support from Australian Research Council through Australia Laureate Fellowship. References [1] M.F. Ashby, A.L. Greer, Scripta Mater. 54 (2006) 321. [2] D.C. Hofmann, J.Y. Suh, A. Wiest, G. Duan, M.L. Lind, M.D. Demetriou, W.L. Johnson, Nature 451 (2008) 1085. [3] L. Tian, Y.Q. Cheng, Z.W. Shan, C.C. Wang, X.D. Han, J. Sun, E. Ma, Nature Commun.. 3 (2012) 609. [4] M.W. Chen, A. Inoue, W. Zhang, T. Sakurai, Phys. Rev. Lett. 96 (2006) 245502. [5] H.K. Kim, K.B. Lee, J.C. Lee, Mater. Sci. Eng. A 552 (2012) 399. [6] X.J. Gu, S.J. Poon, G.J. Shiflet, J.J. Lewandowski, Acta Mater. 58 (2010) 1708. [7] J.C. Lee, Y.C. Kim, J.P. Ahn, H.S. Kim, Acta Mater. 53 (2005) 129. [8] Y.H. Liu, G. Wang, R.J. Wang, D.Q. Zhao, M.X. Pan, W.H. Wang, Science 315 (2007) 1385. [9] H.B. Yu, J. Hu, X.X. Xia, B.A. Sun, X.X. Li, W.H. Wang, H.Y. Bai, Scripta Mater. 61 (2009) 640. [10] K. Lu, J. Lu, Mater. Sci. Eng. A 375–377 (2004) 38. [11] D.K. Yang, P. Cizek, D. Fabijanic, J.T. Wang, B.S. Li, P.D. Hodgson, Work Hardening in Ultrafine-Grained Titanium: Multilayering and Grading, Submitted for publication. [12] M.D. Uchic, D.M. Dimiduk, Mater. Sci. Eng. A 400–401 (2005) 268. [13] F. Sun, P. Rojas, A. Zuniga, E.J. Lavernia, Mater. Sci. Eng. A 430 (2006) 90. [14] A. Sagel, H. Sieber, H.J. Fecht, J.H. Ferepezko, Acta Mater. 46 (1998) 4233. [15] S. Ohsaki, S. Kato, N. Tsuji, T. Ohkubo, K. Hono, Acta Mater. 55 (2007) 2885. [16] C.J. Lee, J.C. Huang, T.G. Nieh, Appl. Phys. Lett. 91 (2007) 161913. [17] Y.H. Lee, Y.T. Cheng, H.S. Chou, H.M. Chen, X.H. Du, C.I. Chang, J.C. Huang, S.R. Jian, J.S.C. Jang, T.G. Nieh, Scripta Mater. 58 (2008) 890. [18] B.E. Schuster, Q. Wei, M.H. Ervin, S.O. Hruszkewycz, M.K. Miller, T.C. Hufnagel, K.T. Ramesh, Scripta Mater. 57 (2007) 517. [19] L.C. Zhang, F. Jiang, Y.L. Zhao, J.F. Zhang, L. He, J. Sun, Mater. Sci. Eng. A 527 (2010) 4122. [20] A. Bharathula, S.W. Lee, W.J. Wright, K.M. Flores, Acta Mater. 58 (2010) 5789. [21] H. Zhang, B.E. Schuster, Q. Wei, K.T. Ramesh, Scripta Mater. 54 (2006) 181. [22] H. Chen, Y. He, G.J. Shiflet, S.J. Poon, Nature 367 (1994) 541. [23] W.H. Jiang, M. Atzmon, Acta Mater. 51 (2003) 4095. [24] J.J. Kim, Y. Choi, S. Suresh, A.S. Argon, Science 295 (2002) 654. [25] J.C. Lee, Y.C. Kim, J.P. Ahn, H.S. Kim, S.H. Lee, B.J. Lee, Acta Mater. 52 (2003) 1525. [26] S.W. Lee, M.Y. Huh, E. Fleury, J.C. Lee, Acta Mater. 54 (2006) 349.