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Bimodal microstructure and deformation of cryomilled bulk nanocrystalline Al–7.5Mg alloy
Materials Science and Engineering A 410–411 (2005) 462–467
Bimodal microstructure and deformation of cryomilled bulk nanocrystalline Al–7.5Mg alloy Z. Lee a,c,∗ , D.B. Witkin b , V. Radmilovic c , E.J. Lavernia d , S.R. Nutt a a
b
Department of Materials Science, University of Southern California, Los Angeles, CA 90089-0241, USA Department of Chemical Engineering and Materials Science, University of California, Irvine, CA 92697-2575, USA c National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, USA d Department of Chemical Engineering and Materials Science, University of California, Davis, CA 95616, USA Received in revised form 25 January 2005
Abstract The microstructure, mechanical properties and deformation response of bimodal structured nanocrystalline Al–7.5Mg alloy were investigated. Grain refinement was achieved by cryomilling of atomized Al–7.5Mg powders, and then cryomilled nanocrystalline powders blended with 15 and 30% unmilled coarse-grained powders were consolidated by hot isostatic pressing followed by extrusion to produce bulk nanocrystalline alloys. Bimodal bulk nanocrystalline Al–7.5Mg alloys, which were comprised of nanocrystalline grains separated by coarse-grain regions, show balanced mechanical properties of enhanced yield and ultimate strength and reasonable ductility and toughness compared to comparable conventional alloys and nanocrystalline metals. The investigation of tensile and hardness test suggests unusual deformation mechanisms and interactions between ductile coarse-grain bands and nanocrystalline regions. © 2005 Elsevier B.V. All rights reserved. Keywords: Nanocrystalline; Bimodal; Cryomilling; Aluminum; Deformation
1. Introduction In recent years, nanocrystalline (nanostructured or ultrafinegrained) metals have become the topic of numerous scientific and technological studies because these materials exhibit remarkable improvements in strength and create the possibility of weight savings. However, nanocrystalline materials generally possess insufficient ductility and a reduced toughness compared to their coarse-grained conventional counterparts. Attempts to address the loss of ductility of nanocrystalline metals include: inducing the formation of a bimodal grain microstructure with grains sufficiently large to facilitate dislocation activity [1] and annealing to promote grain growth [2]. The annealing approach results in strength degradation due to normal and abnormal grain growth of the nanocrystalline grains. In a recent report, a nonuniform bimodal grain size distribution with micrometer-sized grains embedded inside a matrix of nanocrystalline and ultrafine (<300 nm) grains was formed in pure Cu by cold rolling followed
∗
Corresponding author. Tel.: +1 5104864590; fax: +1 5104865888. E-mail address: [email protected] (Z. Lee).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.08.104
by annealing to achieve a high tensile ductility with relatively high strength [2]. In other reports, Legros et al. [3] and Tellkamp et al. [1] also observed that incorporation of coarser grains improved the ductility of nanocrystalline Cu and Al 5083 alloy, respectively. The coarser grains were formed by recrystallization during warm compaction [3] or abnormal grain growth during consolidation by hot isostatic pressing (HIPing) and extrusion [1]. In both cases, the uniaxial tensile fracture strains were higher than typical nanocrystalline metals, while considerable strengthening was achieved relative to the properties of conventional materials. Previous findings suggest that the presence of coarser grains within the nanocrystalline matrix may enhance the ductility of nanocrystalline materials [4–6]. However, in practice, selective grain growth and achievement of target grain size and phase of embedded coarse-grains may be difficult or impossible to control by thermal treatment. Furthermore, this method has inherent limitations in the design and control of the specific microstructure, which is closely dependent on balanced properties of ductility and strength. In contrast, a bimodal (or multi-scale) grain structure can be achieved by consolidation of blended powders. In this process, cryomilled nanocrystalline and coarse-grained
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transmission electron microscopy (SEM and TEM) and optical microscopy. The stress–strain behavior was measured using uniaxial tensile tests, and microhardness measurements were performed on extruded samples to reveal the interactions between coarse-grain and nanocrystalline regions. 2. Experimental procedures
Fig. 1. Incorporating coarse-grains to improve ductility of nanocrystalline metals (a) by annealing of bulk nanocrystalline materials, and (b) by consolidation of blended powders followed by extrusion.
powders (and sometimes, intermediate grains and dispersoids) can be combined in specific proportions, enabling the design and manufacture of materials that achieve the desired balance of enhanced strength with acceptable ductility and toughness. A schematic comparison of these two methods is illustrated in Fig. 1. Moreover, for the issue of structural applications, mechanical alloying allows the production of nanocrystalline metal samples of a sufficient size with the enhanced mechanical properties, while several other techniques have been limited to the fabrication of small samples to explore the intrinsic properties of the materials [7–9]. In the present work, ductile-phase toughening in bimodal structured Al–7.5Mg was achieved by deliberate blending of unmilled coarse-grained powders with cryomilled nanocrystalline powders in select proportions. The microstructures and deformation mechanisms were investigated using scanning and
Prealloyed Al–7.5%Mg (in weight percent) powders were used to produce materials for this work. Nanocrystalline powders were prepared using low-energy mechanical attrition at a cryogenic temperature (cryomilling) with a stainless steel vessel and milling balls with a diameter of 6.4 mm. The ballto-charge ratio was 36:1, with stearic acid added at 0.25% of the powder weight to moderate the cold welding process. The cryomilling apparatus was operated at 180 rpm for 8 h, and maintained at a temperature of −190 ◦ C using flowing liquid nitrogen in the vessel, submerging the powder and balls. Cryomilled nanocrystalline Al–7.5Mg powders (average grain size of ∼20 nm) were combined with 0, 15 and 30 wt.% unmilled coarse-grained Al–7.5Mg powders (particle size of 4–40 m) and blended. These samples were loaded in aluminum cans for vacuum degassing at 400 ◦ C. The cans were sealed and consolidated in HIPing at 325 ◦ C and 7 MPa to produce bimodal samples. The consolidated billets were extruded using an extrusion ratio of 6.5:1 and proprietary rate and temperature parameters. The resulting extrusions were machined to produce cylindrical tensile specimens (tensile axis parallel to the extrusion direction) with a gauge length of 13.5 mm and a gauge diameter of 3.5 mm. Tensile tests were performed on a universal testing machine at room temperature, using a nominal strain rate of 10−3 s−1 . The microstructure of extrusions was examined using metallography and TEM. Microhardness measurements were performed with a Vickers indenter on the polished and chemically etched surfaces of longitudinal and transverse sections. Microindentations were made using a 10 g mass for loading for coarse-grain and nanocrystalline regions, and macroindentations were made using 1 kg indenter which sampled both coarse-grain and nanocrystalline regions. Brinell indentations using 60 kg normal to the extrusion direction also were performed to investigate apparent subsurface plastic deformation. The polished and chemically etched surfaces and indentations were observed by optical microscopy and SEM. The polished cross-sections of extrusions along extrusion direction were analyzed by X-ray diffraction (XRD). 3. Results and discussion The chemically etched sections of as-extruded bimodal samples are shown in Fig. 2. The alloy containing 0% unmilled (i.e. 100% cryomilled) powder exhibited uniform nanocrystalline grains with a few percent of residual coarse-grains. In contrast, the samples containing 15 and 30% unmilled powders revealed light coarse-grain (CG) regions and darker nanocrystalline (NC) regions in the images. Coarse-grain regions formed elongated bands that extended along the extrusion direction, while the nanocrystalline regions formed a continuous matrix. The CG
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Fig. 2. Bimodal microstructures of as-extruded Al–7.5Mg alloy with 0, 15 and 30% coarse-grain content according to extrusion and transverse direction. The bright and dark parts indicate coarse-grain bands and nanocrystalline region, respectively.
bands exhibited elongated grains and sub-grains with the grain size of approximately 1000 nm, compared with 100–300 nm grains in equiaxed nanocrystalline regions according to TEM observations. The formation of sub-grains within the coarsegrain bands may be facilitated by the rotational forces introduced by radial compression of the billet in the extrusion die. In the sample produced with 30% unmilled powder, the CG bands were about 240 m long and 20 m wide. The average spacing between CG bands was about 25 m, and in general they were dispersed uniformly. In the sample produced from 15% unmilled powder, the CG bands were shorter and narrower than those in the 30% CG sample. Also, the spacing between the CG bands was greater in the 15% CG sample than in the 30% CG sample. The grain size distributions of the as-extruded alloys are shown in histograms in Fig. 3. The mean grain size estimated from approximately 200 grains was 203 nm for 0% CG, 221 nm for 15% CG and 313 nm for 30% CG, respectively. The number of coarse-grains increased with increasing volume fraction of unmilled powder from 0 to 30%. The interfaces between
nanocrystalline region and CG bands are remarkably discrete and abrupt. The composition within the coarse-grain bands and the nanocrystalline regions was consistent with the alloy composition of Al–7.5Mg. The XRD patterns of 0, 15 and 30% CG samples are shown in Fig. 4. The XRD scan along the normal to the extrusion direction revealed a notable increase in the (2 2 0) peak intensity with increasing CG content, while the XRD scan along the extrusion direction showed a moderate increase in the (3 1 1) peak. The increase in the (2 2 0) peak intensity may result from the contribution of increasing volume fraction of fibrous coarse-grains. The (2 2 0) peak may correspond to the plane normal to the radial direction of the extrusion rod [10]. The stress–strain curves obtained from uniaxial tensile tests of the different bimodal samples are displayed as a function of CG content in Fig. 5 [11]. The sample produced from cryomilled powder alone (i.e. 0% CG) exhibited a yield strength more than four times that of conventional Al 5083, a commercial alloy with a similar composition but having a grain size of
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Fig. 3. Histogram of grain size distribution of as-extruded Al–7.5Mg alloy with 0, 15 and 30% coarse-grain content.
about 5 m. This sample, however, showed very little ductility and negligible work-hardening. In contrast, the 30% CG alloy exhibited elastic-nearly perfectly plastic stress–strain behavior. This phenomenon is consistent with the uniaxial tension stress–strain behavior reported for nanocrystalline Al Ti Cu [5,6] and Al alloy [12]. The yielding was followed by extended periods of mild work softening. The tensile stress–strain behavior of Al–7.5Mg with 15% CG exhibited intermediate behavior, with only a slight improvement in ductility and brief work softening. The ductility and toughness of bimodal Al–7.5Mg alloy with 30% CG increased over 300% compared with the 0% CG sample, while the yield stress decreased just 14%. The peak flow stresses of the Al–7.5Mg with 0, 15 and 30% CG were about 847, 778 and 734 MPa at room temperature, respectively. These values are notable for an aluminum alloy, given that the flow stress for conventional 5083 Al is about 281 MPa. Especially noteworthy is the elongation-to-failure of Al–7.5Mg with 30% CG fraction, which was about 7% at room temperature. These results indicate a correlation between toughness and CG content. However, unlike alloys with a single grain size, the strength and ductility of the present material cannot be attributed solely to grain size refinement. These results also may indicate a change in deformation mechanism stemming from the unusual structure. The Vickers hardness of individual CG and NC regions was measured using a 10 gf load, thereby permitting separate sampling of CG and NC regions, as shown in Fig. 6. The size difference of indentations from individual CG and NC regions are evident in the image (a). The hardness of NC regions was essentially constant in alloys with 0, 15 and 30% CG, while the hardness of CG regions decreased with increasing CG content. The insensitivity of NC hardness to CG content stems from the higher hardness, which restricted plastic flow to a relatively small zone beneath the indenter and limited the interaction with adjoining coarse-grain regions. However, the hardness of the ductile coarse-grain regions showed a clear dependence on CG
Fig. 4. The XRD patterns of as-extruded Al–7.5Mg alloys with 0, 15 and 30% coarse-grain content. (a) The XRD scan along the normal to extrusion direction, and (b) the XRD scan along the extrusion direction.
Fig. 5. Uniaxial tensile stress–strain behaviors of conventional 5083 Al alloy and as-extruded Al–7.5Mg alloy with 0, 15 and 30% coarse-grain content.
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was dependent on coarse-grain content, while that of nanocrystalline regions was not. The Vickers indentations using a 1 kgf load were also performed on longitudinal transverse sections of extrusions, which was sufficient to sample several NC regions and CG bands. SEM micrographs of a 1 kgf Vickers indentation on 30% CG are shown in Fig. 7. The flow patterns are unusual for a single-phase material, and are unlikely to occur except in duplex structures. In particular, the images show barrel-shaped indentation faces, which resulted from extensive flow of coarse-grain material around the faces of the indenter. In contrast, where the indenter contacted nanocrystalline regions, little flow occurred, and the impression was much clearer. These observations illustrate the substantial difference in flow stress between NC and CG regions, and the tendency for the harder matrix to constrain the flow within softer, CG material.
Fig. 6. (a) Vickers indentation of individual coarse-grain and nanocrystalline regions using a 10 gf load. (b) Vickers hardness measured from individual coarsegrain and nanocrystalline regions of as-extruded bimodal Al–7.5Mg alloys.
content, reflecting a larger plastic zone beneath the indenter. Still, the hardness of individual CG regions in the 30% CG sample was larger than conventional 5083 Al, reflecting the flow constraint of the nanocrystalline matrix. The hardness of Al–7.5Mg with 0% coarse-grain was slightly lower than that of nanocrystalline regions in Al–7.5Mg 15% and 30% CG, possibly because 0% CG contained a few percent of residual coarsegrains. Consequently, the hardness of local coarse-grain regions
Fig. 7. The scanning electron micrograph of Vickers indentation boundary using 1 kgf load on Al–7.5Mg alloy with 30% coarse-grain content along the transverse direction.
Fig. 8. The cross-sectional micrographs of Brinell indentations using 60 kg load on Al–7.5Mg alloy with 0% (a), 15% (b) and 30% (c) coarse-grain content along the transverse direction, respectively.
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Patterns of plastic flow were further revealed by examination of Brinell macrohardness indentations. Subsurface plastic deformation was evident in polished and etched cross-sections of the indented surface. The extent of deformation was revealed by the bending of CG bands, as shown in Fig. 8. The depth of the plastic zone beneath the indentation increased with increasing of CG content. In fact, the apparent deformation depth in samples with 30% CG was nearly three times that of samples with 0% CG. Interestingly, this increase is similar to the increase in tensile elongation observed when the CG content was increased from 0 to 30%, as shown in Fig. 5.
The bimodal grain structures exhibit unusual deformation mechanisms similar to the ductile-phase toughening of brittle materials. Coarse-grain bands tend to deform locally at stress concentrations, arresting cracks by local blunting and resisting crack growth by bridging of crack wakes. The deformation mechanisms can be altered by controlling the morphology and dispersion of the CG phase, as well as the interface properties. Future experiments will be directed at understanding these deformation mechanisms and the patterns of strain distribution within the microstructure.
4. Conclusions
Support from the Office of Naval Research (contract ONR00014-03-1-0149 and ONR00014-03-C-0163) is gratefully acknowledged.
Bimodal structures of Al–7.5Mg composed of nanocrystalline grains and elongated coarse-grains were produced by consolidation of cryomilled powders, resulting in ductilephase toughening. Bimodal bulk nanocrystalline Al Mg alloys showed balanced mechanical properties, including enhanced yield and ultimate strength, and reasonable ductility and toughness compared to comparable conventional alloys and with materials comprised of nanocrystalline grains only. The present work presents what may be a new paradigm in metals processing. Powders of different grain sizes can be blended prior to consolidation to produce specific distributions of grain sizes that were previously impossible. Using this process, one can design single-phase microstructures with local variations in grain size (and strength) to achieve unique combinations of properties. The distribution of grain sizes can be controlled through the proportions of the powder blend to produce alloys with prescribed combinations of strength, ductility, and toughness. The approach is expected to yield similar performance enhancements with other aluminum alloys, as well as with titanium and ferrous alloys.
Acknowledgement
References [1] V.L. Tellkamp, A. Melmed, E.J. Lavernia, Metall. Mater. Trans. A 32 (2001) 2335–2343. [2] Y. Wang, M. Chen, F. Zhou, E. Ma, Nature 419 (2002) 912–915. [3] M. Legros, B.R. Elliot, M.N. Rittner, J.R. Weertman, K.J. Hemker, Phil. Mag. A 80 (4) (2000) 1017–1026. [4] C.C. Koch, D.G. Morris, K. Lu, A. Inoue, MRS Bull. 24 (1999) 54–58. [5] R.W. Hayes, R. Rodriguez, E.J. Lavernia, Acta Mater. 49 (2001) 4055–4068. [6] Z. Lee, R. Rodriguez, R.W. Hayes, E.J. Lavernia, S.R. Nutt, Metall. Mater. Trans. A 34 (2003) 1473–1481. [7] R.W. Siegel, MRS Bull. 15 (1990) 60–67. [8] C.C. Koch, NanoStruct. Mater. 9 (1997) 13–22. [9] C. Suryanarayana, C.C. Koch, in: C. Suryanarayana (Ed.), NonEquilibrium Processing of Materials, Pergamon, New York, 1999, pp. 313–344. [10] D.N. Lee, Y.H. Chung, M.C. Shin, Scripta Metall. 17 (1983) 339–342. [11] D. Witkin, Z. Lee, R. Rodriguez, S.R. Nutt, E.J. Lavernia, Scripta Mater. 49 (2003) 297–302. [12] X.K. Sun, H.T. Cong, M. Sun, M.C. Yang, Metall. Mater. Trans. A 31 (2000) 1017–1024.
Bimodal microstructure and deformation of cryomilled bulk nanocrystalline Al–7.5Mg alloy Z. Lee a,c,∗ , D.B. Witkin b , V. Radmilovic c , E.J. Lavernia d , S.R. Nutt a a
b
Department of Materials Science, University of Southern California, Los Angeles, CA 90089-0241, USA Department of Chemical Engineering and Materials Science, University of California, Irvine, CA 92697-2575, USA c National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, USA d Department of Chemical Engineering and Materials Science, University of California, Davis, CA 95616, USA Received in revised form 25 January 2005
Abstract The microstructure, mechanical properties and deformation response of bimodal structured nanocrystalline Al–7.5Mg alloy were investigated. Grain refinement was achieved by cryomilling of atomized Al–7.5Mg powders, and then cryomilled nanocrystalline powders blended with 15 and 30% unmilled coarse-grained powders were consolidated by hot isostatic pressing followed by extrusion to produce bulk nanocrystalline alloys. Bimodal bulk nanocrystalline Al–7.5Mg alloys, which were comprised of nanocrystalline grains separated by coarse-grain regions, show balanced mechanical properties of enhanced yield and ultimate strength and reasonable ductility and toughness compared to comparable conventional alloys and nanocrystalline metals. The investigation of tensile and hardness test suggests unusual deformation mechanisms and interactions between ductile coarse-grain bands and nanocrystalline regions. © 2005 Elsevier B.V. All rights reserved. Keywords: Nanocrystalline; Bimodal; Cryomilling; Aluminum; Deformation
1. Introduction In recent years, nanocrystalline (nanostructured or ultrafinegrained) metals have become the topic of numerous scientific and technological studies because these materials exhibit remarkable improvements in strength and create the possibility of weight savings. However, nanocrystalline materials generally possess insufficient ductility and a reduced toughness compared to their coarse-grained conventional counterparts. Attempts to address the loss of ductility of nanocrystalline metals include: inducing the formation of a bimodal grain microstructure with grains sufficiently large to facilitate dislocation activity [1] and annealing to promote grain growth [2]. The annealing approach results in strength degradation due to normal and abnormal grain growth of the nanocrystalline grains. In a recent report, a nonuniform bimodal grain size distribution with micrometer-sized grains embedded inside a matrix of nanocrystalline and ultrafine (<300 nm) grains was formed in pure Cu by cold rolling followed
∗
Corresponding author. Tel.: +1 5104864590; fax: +1 5104865888. E-mail address: [email protected] (Z. Lee).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.08.104
by annealing to achieve a high tensile ductility with relatively high strength [2]. In other reports, Legros et al. [3] and Tellkamp et al. [1] also observed that incorporation of coarser grains improved the ductility of nanocrystalline Cu and Al 5083 alloy, respectively. The coarser grains were formed by recrystallization during warm compaction [3] or abnormal grain growth during consolidation by hot isostatic pressing (HIPing) and extrusion [1]. In both cases, the uniaxial tensile fracture strains were higher than typical nanocrystalline metals, while considerable strengthening was achieved relative to the properties of conventional materials. Previous findings suggest that the presence of coarser grains within the nanocrystalline matrix may enhance the ductility of nanocrystalline materials [4–6]. However, in practice, selective grain growth and achievement of target grain size and phase of embedded coarse-grains may be difficult or impossible to control by thermal treatment. Furthermore, this method has inherent limitations in the design and control of the specific microstructure, which is closely dependent on balanced properties of ductility and strength. In contrast, a bimodal (or multi-scale) grain structure can be achieved by consolidation of blended powders. In this process, cryomilled nanocrystalline and coarse-grained
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transmission electron microscopy (SEM and TEM) and optical microscopy. The stress–strain behavior was measured using uniaxial tensile tests, and microhardness measurements were performed on extruded samples to reveal the interactions between coarse-grain and nanocrystalline regions. 2. Experimental procedures
Fig. 1. Incorporating coarse-grains to improve ductility of nanocrystalline metals (a) by annealing of bulk nanocrystalline materials, and (b) by consolidation of blended powders followed by extrusion.
powders (and sometimes, intermediate grains and dispersoids) can be combined in specific proportions, enabling the design and manufacture of materials that achieve the desired balance of enhanced strength with acceptable ductility and toughness. A schematic comparison of these two methods is illustrated in Fig. 1. Moreover, for the issue of structural applications, mechanical alloying allows the production of nanocrystalline metal samples of a sufficient size with the enhanced mechanical properties, while several other techniques have been limited to the fabrication of small samples to explore the intrinsic properties of the materials [7–9]. In the present work, ductile-phase toughening in bimodal structured Al–7.5Mg was achieved by deliberate blending of unmilled coarse-grained powders with cryomilled nanocrystalline powders in select proportions. The microstructures and deformation mechanisms were investigated using scanning and
Prealloyed Al–7.5%Mg (in weight percent) powders were used to produce materials for this work. Nanocrystalline powders were prepared using low-energy mechanical attrition at a cryogenic temperature (cryomilling) with a stainless steel vessel and milling balls with a diameter of 6.4 mm. The ballto-charge ratio was 36:1, with stearic acid added at 0.25% of the powder weight to moderate the cold welding process. The cryomilling apparatus was operated at 180 rpm for 8 h, and maintained at a temperature of −190 ◦ C using flowing liquid nitrogen in the vessel, submerging the powder and balls. Cryomilled nanocrystalline Al–7.5Mg powders (average grain size of ∼20 nm) were combined with 0, 15 and 30 wt.% unmilled coarse-grained Al–7.5Mg powders (particle size of 4–40 m) and blended. These samples were loaded in aluminum cans for vacuum degassing at 400 ◦ C. The cans were sealed and consolidated in HIPing at 325 ◦ C and 7 MPa to produce bimodal samples. The consolidated billets were extruded using an extrusion ratio of 6.5:1 and proprietary rate and temperature parameters. The resulting extrusions were machined to produce cylindrical tensile specimens (tensile axis parallel to the extrusion direction) with a gauge length of 13.5 mm and a gauge diameter of 3.5 mm. Tensile tests were performed on a universal testing machine at room temperature, using a nominal strain rate of 10−3 s−1 . The microstructure of extrusions was examined using metallography and TEM. Microhardness measurements were performed with a Vickers indenter on the polished and chemically etched surfaces of longitudinal and transverse sections. Microindentations were made using a 10 g mass for loading for coarse-grain and nanocrystalline regions, and macroindentations were made using 1 kg indenter which sampled both coarse-grain and nanocrystalline regions. Brinell indentations using 60 kg normal to the extrusion direction also were performed to investigate apparent subsurface plastic deformation. The polished and chemically etched surfaces and indentations were observed by optical microscopy and SEM. The polished cross-sections of extrusions along extrusion direction were analyzed by X-ray diffraction (XRD). 3. Results and discussion The chemically etched sections of as-extruded bimodal samples are shown in Fig. 2. The alloy containing 0% unmilled (i.e. 100% cryomilled) powder exhibited uniform nanocrystalline grains with a few percent of residual coarse-grains. In contrast, the samples containing 15 and 30% unmilled powders revealed light coarse-grain (CG) regions and darker nanocrystalline (NC) regions in the images. Coarse-grain regions formed elongated bands that extended along the extrusion direction, while the nanocrystalline regions formed a continuous matrix. The CG
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Fig. 2. Bimodal microstructures of as-extruded Al–7.5Mg alloy with 0, 15 and 30% coarse-grain content according to extrusion and transverse direction. The bright and dark parts indicate coarse-grain bands and nanocrystalline region, respectively.
bands exhibited elongated grains and sub-grains with the grain size of approximately 1000 nm, compared with 100–300 nm grains in equiaxed nanocrystalline regions according to TEM observations. The formation of sub-grains within the coarsegrain bands may be facilitated by the rotational forces introduced by radial compression of the billet in the extrusion die. In the sample produced with 30% unmilled powder, the CG bands were about 240 m long and 20 m wide. The average spacing between CG bands was about 25 m, and in general they were dispersed uniformly. In the sample produced from 15% unmilled powder, the CG bands were shorter and narrower than those in the 30% CG sample. Also, the spacing between the CG bands was greater in the 15% CG sample than in the 30% CG sample. The grain size distributions of the as-extruded alloys are shown in histograms in Fig. 3. The mean grain size estimated from approximately 200 grains was 203 nm for 0% CG, 221 nm for 15% CG and 313 nm for 30% CG, respectively. The number of coarse-grains increased with increasing volume fraction of unmilled powder from 0 to 30%. The interfaces between
nanocrystalline region and CG bands are remarkably discrete and abrupt. The composition within the coarse-grain bands and the nanocrystalline regions was consistent with the alloy composition of Al–7.5Mg. The XRD patterns of 0, 15 and 30% CG samples are shown in Fig. 4. The XRD scan along the normal to the extrusion direction revealed a notable increase in the (2 2 0) peak intensity with increasing CG content, while the XRD scan along the extrusion direction showed a moderate increase in the (3 1 1) peak. The increase in the (2 2 0) peak intensity may result from the contribution of increasing volume fraction of fibrous coarse-grains. The (2 2 0) peak may correspond to the plane normal to the radial direction of the extrusion rod [10]. The stress–strain curves obtained from uniaxial tensile tests of the different bimodal samples are displayed as a function of CG content in Fig. 5 [11]. The sample produced from cryomilled powder alone (i.e. 0% CG) exhibited a yield strength more than four times that of conventional Al 5083, a commercial alloy with a similar composition but having a grain size of
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Fig. 3. Histogram of grain size distribution of as-extruded Al–7.5Mg alloy with 0, 15 and 30% coarse-grain content.
about 5 m. This sample, however, showed very little ductility and negligible work-hardening. In contrast, the 30% CG alloy exhibited elastic-nearly perfectly plastic stress–strain behavior. This phenomenon is consistent with the uniaxial tension stress–strain behavior reported for nanocrystalline Al Ti Cu [5,6] and Al alloy [12]. The yielding was followed by extended periods of mild work softening. The tensile stress–strain behavior of Al–7.5Mg with 15% CG exhibited intermediate behavior, with only a slight improvement in ductility and brief work softening. The ductility and toughness of bimodal Al–7.5Mg alloy with 30% CG increased over 300% compared with the 0% CG sample, while the yield stress decreased just 14%. The peak flow stresses of the Al–7.5Mg with 0, 15 and 30% CG were about 847, 778 and 734 MPa at room temperature, respectively. These values are notable for an aluminum alloy, given that the flow stress for conventional 5083 Al is about 281 MPa. Especially noteworthy is the elongation-to-failure of Al–7.5Mg with 30% CG fraction, which was about 7% at room temperature. These results indicate a correlation between toughness and CG content. However, unlike alloys with a single grain size, the strength and ductility of the present material cannot be attributed solely to grain size refinement. These results also may indicate a change in deformation mechanism stemming from the unusual structure. The Vickers hardness of individual CG and NC regions was measured using a 10 gf load, thereby permitting separate sampling of CG and NC regions, as shown in Fig. 6. The size difference of indentations from individual CG and NC regions are evident in the image (a). The hardness of NC regions was essentially constant in alloys with 0, 15 and 30% CG, while the hardness of CG regions decreased with increasing CG content. The insensitivity of NC hardness to CG content stems from the higher hardness, which restricted plastic flow to a relatively small zone beneath the indenter and limited the interaction with adjoining coarse-grain regions. However, the hardness of the ductile coarse-grain regions showed a clear dependence on CG
Fig. 4. The XRD patterns of as-extruded Al–7.5Mg alloys with 0, 15 and 30% coarse-grain content. (a) The XRD scan along the normal to extrusion direction, and (b) the XRD scan along the extrusion direction.
Fig. 5. Uniaxial tensile stress–strain behaviors of conventional 5083 Al alloy and as-extruded Al–7.5Mg alloy with 0, 15 and 30% coarse-grain content.
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was dependent on coarse-grain content, while that of nanocrystalline regions was not. The Vickers indentations using a 1 kgf load were also performed on longitudinal transverse sections of extrusions, which was sufficient to sample several NC regions and CG bands. SEM micrographs of a 1 kgf Vickers indentation on 30% CG are shown in Fig. 7. The flow patterns are unusual for a single-phase material, and are unlikely to occur except in duplex structures. In particular, the images show barrel-shaped indentation faces, which resulted from extensive flow of coarse-grain material around the faces of the indenter. In contrast, where the indenter contacted nanocrystalline regions, little flow occurred, and the impression was much clearer. These observations illustrate the substantial difference in flow stress between NC and CG regions, and the tendency for the harder matrix to constrain the flow within softer, CG material.
Fig. 6. (a) Vickers indentation of individual coarse-grain and nanocrystalline regions using a 10 gf load. (b) Vickers hardness measured from individual coarsegrain and nanocrystalline regions of as-extruded bimodal Al–7.5Mg alloys.
content, reflecting a larger plastic zone beneath the indenter. Still, the hardness of individual CG regions in the 30% CG sample was larger than conventional 5083 Al, reflecting the flow constraint of the nanocrystalline matrix. The hardness of Al–7.5Mg with 0% coarse-grain was slightly lower than that of nanocrystalline regions in Al–7.5Mg 15% and 30% CG, possibly because 0% CG contained a few percent of residual coarsegrains. Consequently, the hardness of local coarse-grain regions
Fig. 7. The scanning electron micrograph of Vickers indentation boundary using 1 kgf load on Al–7.5Mg alloy with 30% coarse-grain content along the transverse direction.
Fig. 8. The cross-sectional micrographs of Brinell indentations using 60 kg load on Al–7.5Mg alloy with 0% (a), 15% (b) and 30% (c) coarse-grain content along the transverse direction, respectively.
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Patterns of plastic flow were further revealed by examination of Brinell macrohardness indentations. Subsurface plastic deformation was evident in polished and etched cross-sections of the indented surface. The extent of deformation was revealed by the bending of CG bands, as shown in Fig. 8. The depth of the plastic zone beneath the indentation increased with increasing of CG content. In fact, the apparent deformation depth in samples with 30% CG was nearly three times that of samples with 0% CG. Interestingly, this increase is similar to the increase in tensile elongation observed when the CG content was increased from 0 to 30%, as shown in Fig. 5.
The bimodal grain structures exhibit unusual deformation mechanisms similar to the ductile-phase toughening of brittle materials. Coarse-grain bands tend to deform locally at stress concentrations, arresting cracks by local blunting and resisting crack growth by bridging of crack wakes. The deformation mechanisms can be altered by controlling the morphology and dispersion of the CG phase, as well as the interface properties. Future experiments will be directed at understanding these deformation mechanisms and the patterns of strain distribution within the microstructure.
4. Conclusions
Support from the Office of Naval Research (contract ONR00014-03-1-0149 and ONR00014-03-C-0163) is gratefully acknowledged.
Bimodal structures of Al–7.5Mg composed of nanocrystalline grains and elongated coarse-grains were produced by consolidation of cryomilled powders, resulting in ductilephase toughening. Bimodal bulk nanocrystalline Al Mg alloys showed balanced mechanical properties, including enhanced yield and ultimate strength, and reasonable ductility and toughness compared to comparable conventional alloys and with materials comprised of nanocrystalline grains only. The present work presents what may be a new paradigm in metals processing. Powders of different grain sizes can be blended prior to consolidation to produce specific distributions of grain sizes that were previously impossible. Using this process, one can design single-phase microstructures with local variations in grain size (and strength) to achieve unique combinations of properties. The distribution of grain sizes can be controlled through the proportions of the powder blend to produce alloys with prescribed combinations of strength, ductility, and toughness. The approach is expected to yield similar performance enhancements with other aluminum alloys, as well as with titanium and ferrous alloys.
Acknowledgement
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