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Thermal stability of an ultrafine grain β-Ti alloy
Materials Science & Engineering A 556 (2012) 582–587
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Thermal stability of an ultrafine grain b-Ti alloy D. Kent a,b,n, W.L. Xiao b, G. Wang a,b, Z. Yu c, M.S. Dargusch a,b a b c
Queensland Centre for Advanced Materials Processing and Manufacturing (AMPAM), School of Mechanical and Mining Engineering, The University of Queensland, Australia Defence Materials Technology Centre, Australia Biomaterial Research Centre, Northwest Institute for Nonferrous Metal Research, Xian 710016, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 30 March 2012 Received in revised form 7 July 2012 Accepted 7 July 2012 Available online 16 July 2012
The aim of this research was to investigate the thermal stability of an ultrafine grained metastable b Ti alloy, Ti–25Nb–3Zr–3Mo–2Sn, which was formed using a modified accumulative roll bonding technique. The thermal stability was assessed during annealing heat treatments through property and microstructure observations. Only small changes were detected in the hardness during annealing at temperatures of 400 1C–500 1C. There was no change in the b grain size after prolonged annealing treatments at 400 1C and limited grain growth at 500 1C. At these temperatures coarsening of the a phase and some recovery is evident. Significant recrystallisation of the b phase in conjunction with grain growth occurred at an annealing temperature of 600 1C and was associated with substantial reduction in the hardness. An ultrafine grain structure was retained after prolonged annealing at this temperature. & 2012 Elsevier B.V. All rights reserved.
Keywords: Annealing Electron microscopy Titanium alloys Grain growth
1. Introduction Significant efforts have been devoted to the development of new metastable b titanium alloys for biomedical applications because of their favourable properties in conjunction with compositions which do not include Ni and other toxic elements [1,2]. In addition to biocompatibility, Ti alloys for surgical implants need to be mechanically compatible [3,4]. For orthopaedic implants high strength in conjunction with a low elastic modulus is critical to reducing detrimental ‘‘stress shielding’’ effects. For other implant applications factors such as the surface properties and fatigue resistance are of primary importance. With their potential for enhanced strengths and greater fatigue resistance there is considerable interest in the development of ultrafine grain (UFG) and nanocrystalline (NC) versions of these new metastable b titanium alloys [5–12]. A new metastable b Ti alloy, containing Nb, Zr, Mo and Sn, was recently reported by the authors for use in bio-implant applications [11–14]. The Ti–25Nb–3Zr–3Mo–2Sn alloy has a low modulus of elasticity, high strength and exhibits considerable plasticity and pseudoelastic character. Through the application of grain refinement via severe plastic deformation (SPD), using a modified accumulative roll bonding (ARB) technique, the strength of the alloy was significantly enhanced [11]. A 70% improvement
n Corresponding author at: Queensland Centre for Advanced Materials Processing and Manufacturing (AMPAM), School of Mechanical and Mining Engineering, The University of Queensland, Australia. Tel.: þ61 7 3365 3801; fax: þ15 71 7 3355 3888. E-mail address: [email protected] (D. Kent).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.07.030
in the ultimate tensile strength and a two-fold increase in the 0.5% proof strength of the alloy were achieved while maintaining levels of ductility of around 4.5%. TEM investigations have shown that the ARB processed b Ti–25Nb–3Zr–3Mo–2Sn alloy microstructure is primarily comprised of ultrafine b grains heavily elongated in the rolling direction with a fine dispersion of nanocrystalline a phase precipitates located on the b grain boundaries. The b grains typically have a flattened morphology with a grain size of approximately 130 nm when viewed normal to the rolling direction and a mean grain boundary interval of around 70 nm viewed from the transverse direction. Occasional large b grains in the order of 1–2 mm were also observed amongst the heavily refined structure and nanocrystalline grains with grain sizes in the order of 20 nm were present within shear bands. In order to further the development and application of the UFG metastable b Ti–25Nb–3Zr–3Mo–2Sn alloy it is important to investigate its thermal stability. Thermal treatments are often used in order to relieve residual stresses or to improve the balance between strength and ductility in the as-processed UFG b Ti alloys. However, fine, non-equilibrium microstructures can be prone to rapid grain growth at elevated temperatures, which may significantly impact on their properties and performance. Also for UFG materials produced by SPD rapid grain growth can occur at relatively low temperatures [15,16]. However, in some cases UFG materials achieve good thermal stability at higher temperatures due to the influence of impurities, pores or second phase particles due to grain boundary pinning effects [7,17–19]. Hence the objective of this research was to investigate the thermal stability of the UFG metastable b Ti–25Nb–3Zr–3Mo– 2Sn alloy microstructure.
D. Kent et al. / Materials Science & Engineering A 556 (2012) 582–587
583
2. Method
3. Results and discussion
The Ti–25Nb–3Zr–3Mo–2Sn (wt%) alloy was produced from commercially pure Ti sponge (99.5 wt%), pure Zr bars (99.7 wt%), pure Sn bars (99.9 wt%), pure Mo powder (99.8 wt%) and an intermediate Nb–47wt% Ti alloy. It was melted twice by nonconsumable arc melting to ensure chemical homogeneity and low levels of impurities (Fer0.10 wt%, Sir0.04 wt%, Cr0.03 wt%, Nr0.02 wt%, H r0.003 wt%). The ingot was forged and then hot rolled to a thickness of 1.0 mm at approximately 850 1C (above the b transus). The hot-rolled sheets were cleaned with ethanol, clad in Ti foil, solution treated under a protective Argon atmosphere at a temperature of 710 1C for 1 h and air cooled. Rolling was conducted on a 550 mm (diam.) 650 mm roll set using a roll speed of 10 ms 1. The 1.0 mm solution treated sheet was initially cold rolled at ambient temperature to a thickness of 0.2 mm. The sheet was then sectioned in half, stacked and rolled together. Stacking and rolling comprises a single ARB cycle and the process was repeated 3 times to give a total of 4 cycles, including the initial 80% reduction from 1 mm to 0.2 mm. A 50% total reduction was used for each of the ARB cycles, involving 3 individual roll passes with reductions of 30%, then 15% and finally 5%, such that each time the sheet was cut, stacked and rolled a constant 0.2 mm final sheet thickness was maintained. The sheet dimensions were 0.2 mm 150 mm 200 mm. The surface of the sheets was chemically cleaned using a HF:HNO3:H2O¼1:5:4 acid solution between each cycle. The surface of the Ti ARB sheet was protected during rolling by cladding with stainless steel sheet. The stainless steel clad ARB assembly was preheated to 400 1C for 20 min before rolling. Annealing heat treatments were undertaken on the UFG Ti– 25Nb–3Zr–3Mo–2Sn (wt%) alloy produced after 4 cycles of rolling. The annealing treatments were conducted in a MIHM-VOGT P6/B box furnace. Samples were encased in vacuum sealed quartz tubes and placed in the preheated furnace at the annealing temperature for times up to 180 min. Microhardness tests were performed on a Struers Duramin 20 Micro Vickers hardness tester using a 200 g load. Samples for transmission electron microscopy (TEM) were prepared using twin-jet electropolishing (Tenupol-3, Struers) with a solution containing HClO4 (5%), 2-butoxyethanol (35%) and methanol at 30 1C and 20 V. Some samples were further thinned using a Gatan precision ion polishing system (PIPS, model 691) with Arbeam energy of 4 KeV at an incident angle of 741. TEM was performed using a Philips Tecnai 20 FEG instrument.
A bright field (BF) TEM image of the alloy showing the heavily refined grain structure in the as-rolled condition is presented in Fig. 1(a). The grain boundaries are wavy and not well defined and there is extensive diffraction contrast both within the grains and at grain boundaries. These are features typical of UFG materials formed via SPD and are associated with the non-equilibrium state of the material associated with the high levels of strain and associated dislocation activity imparted to the material during processing [15,19,20]. A selected area diffraction pattern (SADP) from this region shown inset, shows a ring type pattern characteristic of very fine and randomly oriented grains with large grain boundary misorientations, which can be indexed to the b and a phases. Due to the indistinct grain structure of the UFG material under BF imaging conditions, a dark field imaging technique was developed to show contrast between the grains for grain size measurements. The technique is a derivative of hollow-cone dark field microscopy, a method in which an image is formed by summing the electron intensity at the image plane as a chosen annular region of the diffraction pattern circumferentially scans over the objective aperture [21]. A similar technique was employed in this study to obtain dark field (DF) images. However, rather than taking a single long exposure while continuously scanning the diffraction pattern over the objective aperture, a series of 3 or more DF images were taken from the same region of the sample using different portions of the diffraction pattern, selected by rotating the diffraction pattern over the objective aperture between each exposure. Coloured composite dark field (CDF) images were then created by combining a series of three of these DF images using the Analysis—colour mix function within the Digital Micrograph software. Fig. 1(b) is a CDF image of the same area shown in Fig. 1(a) formed using the DF images obtained using representative aperture positions indicated on the SADP (inset) by the coloured circles. The CDF image clearly reveals the grain structure which primarily consists of irregularly shaped b grains with an average grain size determined by the line intercept method of approximately 130 nm [11]. To assess the thermal stability of the UFG microstructure, changes in the properties were evaluated using microhardness measurements over the course of annealing treatments at temperatures ranging from 400 1C to 600 1C for times of up to 3 h. The results of these measurements are presented in Fig. 2. The microhardness of the as-rolled UFG b-Ti alloy was approximately
Fig. 1. TEM of the as-rolled microstructure (4 cycles). (a) BF image reveals a heavily refined grain structure. Inset: SADP exhibiting a ring type pattern typical of very fine and randomly oriented grains indexed to the b and a phases. (b) CDF image showing irregularly shaped b grains with an average grain size of approximately 130 nm. Inset: SADP with aperture positions represented by coloured circles used to form the CDF image.
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310 Hv. Annealing at 400 1C resulted in an increase in the hardness of around 50 Hv, while for the annealing treatment at 450 1C the hardness remained effectively constant. At 500 1C there was an initial rapid drop in the microhardness of around 30 Hv after 15 min followed by a more gradual decrease in the hardness with further annealing time. At 600 1C there was a considerable rapid reduction in the hardness of around 100 Hv within the first
Fig. 2. Change in microhardness over time for different annealing temperatures.
10 min of annealing, after which the hardness remains effectively constant at around 210 Hv. TEM analysis was undertaken in order to assess changes in the microstructure associated with the variations in microhardness detected during annealing. A BF TEM image of a sample annealed for 60 min at 400 1C is shown in Fig. 3(a). At this temperature an increase in the microhardness of around 50 Hv was observed. In comparison to the as-rolled microstructure which was shown in Fig. 1(a), there is reduced diffraction contrast associated with dislocation networks within grains and at the grain boundaries and the boundaries are more distinct. These features may be associated with the process of recovery involving the annihilation of dislocations at the grain boundary walls and via dipole interactions [22,23]. However, it should be noted that the appearance of these features can also be influenced by specific TEM imaging conditions such as the sample thickness [20]. Careful examination of the SADP, shown in Fig. 3(b), also reveals an increased intensity for reflections within the a phase diffraction rings, which is particularly apparent for the inner a (100) reflections. This may be attributed to coarsening of the a phase precipitates during annealing [14,24]. The proportion of a phase did not change considerably during annealing according to the results from the X-ray diffraction analysis, presumably due to its being near the equilibrium as a result of the multiple heating cycles undergone during processing as it was heated to 400 1C before each ARB pass [11]. A DF image formed using a portion of the inner a (100) diffraction ring is shown in Fig. 3(c). It reveals nanoscale and some coarser a precipitates
Fig. 3. TEM from a sample annealed at 400 1C for 60 min. (a) BF image showing reduced diffraction contrast associated with dislocation networks and boundaries which are more distinct in comparison to the as rolled condition. (b) SADP exhibiting a ring type pattern typical of very fine and randomly oriented grains which can be indexed to the b and a phases. (c) A DF formed using a portion of the inner a (100) diffraction ring reveals both nanoscale and some coarser a phase precipitates. (d) A CDF image of the region shown in (a).
D. Kent et al. / Materials Science & Engineering A 556 (2012) 582–587
located on the grain boundaries of the predominantly b phase microstructure. Fig. 3(d) shows a CDF image of the region shown in the BF image (Fig. 3(a)). It reveals that the b grain size is of a similar order of magnitude to that of the as-rolled material and the wavy nature of the grain boundaries is preserved. A thorough analysis of the b grain size during annealing is presented in Fig. 7.
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A BF TEM image of a sample annealed for 10 min at 500 1C is shown in Fig. 4(a) and a CDF image of the same region is shown in Fig. 4(b). After this annealing treatment, there was a small decrease in the microhardness of around 20–30 Hv. Some grains exhibit well defined grain boundaries with a curved appearance as shown in the higher magnification image, Fig. 4(c), contrasting with the wavy appearance of grain boundaries in the as rolled
Fig. 4. TEM from a sample annealed at 500 1C for 10 min. (a) BF image reveals a heavily refined grain structure. (b) CDF image of the region shown in (a) formed using portions of the low order diffraction rings. Inset: SADP exhibiting exhibiting a ring type pattern which can be indexed to the b and a phases. (c) Higher magnification BF image showing well defined grain boundaries with a curved appearance. (d) BF image showing an example of an isolated pocket of lath type a phase precipitates which were observed in occassional relatively large b grains within the UFG microstructure.
Fig. 5. TEM from a sample annealed at 600 1C for 10 min. (a) BF image reveals comparatively large grains with straight sided grain boundaries characteristic of a recrystallised grain structure with SADP, inset, exhibiting diffraction spots and streaks which can be indexed to the b phase. (b) CDF image of the recrystallised grains formed using the low order b phase reflections.
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condition and after annealing at 400 1C. This suggests that the early stages of recrystallisation commence at the higher annealing temperature of 500 1C. Recovery leading to reductions in lattice microstrains and the early stages of recrystallisation may account for the decrease in hardness observed during annealing at 500 1C [23,25]. Examination of the SADP, inset Fig. 4(b), again reveals some increased intensity for reflections within the inner a (100) diffraction ring in comparison to the as-rolled material due to coarsening of the a phase during annealing. There were also isolated pockets of lath type a phase precipitates, shown in Fig. 4(d), which are of a similar morphology to those observed in the coarse grained form of the alloy after aging treatments, and which in this case were observed in the occasional relatively large b grains present within the predominantly UFG microstructure [14]. A BF TEM image of a sample annealed for 10 min at 600 1C is shown in Fig. 5(a). The b grains show a markedly changed equiaxed structure with distinct and predominantly straight grain
boundaries. This annealing treatment resulted in a substantial decrease in the microhardness of around 100 Hv and it is clear that this is associated with recrystallisation of the b phase and subsequent grain growth. The SADP, shown inset in Fig. 5(a), is no longer in the form of continuous diffraction rings and the reflections, which can be indexed to the b phase, are in the form of spots or streaks due to the larger b grain size. The CDF image of the region depicted in Fig. 5(a) is shown in Fig. 5(b), showing the b grains with an average grain size in the order of 0.5–1 mm. An analysis of the b grain size was undertaken from multiple CDF images for each annealing condition using the line intercept method. Examples of the CDF images were presented in Figs. 1, 3–5 and a series of BF–CDF image pairs showing the evolution of the grain size over the course annealing at 500 1C is presented in Fig. 6. They show the gradual increase in the b grain size during annealing at this temperature. The results of b grain size measurements over the course of annealing at temperatures of 400 1C,
Fig. 6. Series of TEM BF–CDF image pairs showing the evolution of the grain size over the course annealing at 500 1C. (a) 10 min, (b) 30 min and (c) 60 min.
D. Kent et al. / Materials Science & Engineering A 556 (2012) 582–587
Fig. 7. An analysis of the b grain size versus annealing time for annealing temperatures of 400 1C, 500 1C and 600 1C undertaken from CDF images using the line intercept method.
500 1C and 600 1C are presented in Fig. 7. It shows that the grain size did not change significantly during annealing at 400 1C and increased gradually during annealing at 500 1C. However, annealing at 600 1C led to a rapid increase in the average grain size within the first 10 min to approximately 0.72 mm, after which it remains effectively constant. The TEM analysis has shown that coarsening of the a phase and the process of recovery, associated with the annihilation of dislocations and reduction of lattice strains, occur simultaneously in the temperature range of 400–500 1C. Significant recrystallisation of the b phase and associated grain growth does not occur until annealing temperatures of 600 1C and above. Despite the recrystallisation of the b phase occurring during annealing at 600 1C, an UFG structure was still retained after 60 min at this temperature. The UFG b Ti alloys show considerably greater thermal stabilty than other UFG Ti alloys produced from commercially pure titanium, for which significant grain growth is observed after annealing at temperatures of 450 1C and above [16]. These results are consistent with the observations of Li et al [7] for a similar nanostructured b alloy, Ti– 24Nb–4Zr–7.9Sn, which showed that significant growth of the b grains to conventional grain sizes larger than 10 mm does not occur until annealing temperatures of 700 1C or above. Nanocrystalline a phase precipitates, which in this case are present after ARB processing, coarsen during annealing and are believed to enhance the thermal stability of the microstructures in comparison to that of other nanostructured or UFG Ti alloys with single phase compositions [7]. Furthermore, the effects of solute drag may also play a more prominent role in retarding the rate of grain growth in the heavily alloyed metastable b Ti alloys in comparison to that of the commercially pure Ti alloys [26,27]. Further work to clarify the role of solute drag and the evolution of the a phase during annealing and their influence on thermal stability will be published in the future. An important feature of metastable UFG b Ti alloys, such as the Ti–25Nb–3Zr–3Mo–2Sn alloy, is their capacity to retain an UFG grain structure at relatively high annealing temperatures around 600 1C, while undergoing recrystallisation of the b phase. It permits the use of thermal annealing treatments to relieve residual stresses and anisotropy of properties imparted to the material during processing, while preserving the potential benefits of UFG Ti alloys such as an enhanced in vitro fibroblast response or improved mechanical properties [5,6].
treatments. Coarsening of the a phase occurred during annealing at temperatures of 400–500 1C in conjunction with recovery involving the annihilation of dislocations in the b phase. Minimal changes were observed in the hardness at these annealing temperatures, with a small increase in the hardness at 400 1C and a small decrease at 500 1C. There was no change in the b grain size after prolonged annealing treatments at 400 1C and limited grain growth after prolonged periods at 500 1C. Recrystallisation of the b phase and substantial grain growth occurred at an annealing temperature of 600 1C, which was associated with significant reductions in the hardness. However, a UFG grain structure with average grain size around 0.72 mm was still retained after prolonged annealing at this temperature. The UFG metastable b Ti alloys have significantly greater thermal stability than commercial purity UFG Ti alloys.
Acknowledgements The authors acknowledge the facilities, and the scientific and technical assistance, of the Australian Microscopy and Microanalysis Research Facility at the Centre for Microscopy and Microanalysis, The University of Queensland.
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4. Conclusions [26]
The thermal stability of an UFG metastable b Ti alloy, Ti– 25Nb–3Zr–3Mo–2Sn, was assessed during annealing heat
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Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Thermal stability of an ultrafine grain b-Ti alloy D. Kent a,b,n, W.L. Xiao b, G. Wang a,b, Z. Yu c, M.S. Dargusch a,b a b c
Queensland Centre for Advanced Materials Processing and Manufacturing (AMPAM), School of Mechanical and Mining Engineering, The University of Queensland, Australia Defence Materials Technology Centre, Australia Biomaterial Research Centre, Northwest Institute for Nonferrous Metal Research, Xian 710016, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 30 March 2012 Received in revised form 7 July 2012 Accepted 7 July 2012 Available online 16 July 2012
The aim of this research was to investigate the thermal stability of an ultrafine grained metastable b Ti alloy, Ti–25Nb–3Zr–3Mo–2Sn, which was formed using a modified accumulative roll bonding technique. The thermal stability was assessed during annealing heat treatments through property and microstructure observations. Only small changes were detected in the hardness during annealing at temperatures of 400 1C–500 1C. There was no change in the b grain size after prolonged annealing treatments at 400 1C and limited grain growth at 500 1C. At these temperatures coarsening of the a phase and some recovery is evident. Significant recrystallisation of the b phase in conjunction with grain growth occurred at an annealing temperature of 600 1C and was associated with substantial reduction in the hardness. An ultrafine grain structure was retained after prolonged annealing at this temperature. & 2012 Elsevier B.V. All rights reserved.
Keywords: Annealing Electron microscopy Titanium alloys Grain growth
1. Introduction Significant efforts have been devoted to the development of new metastable b titanium alloys for biomedical applications because of their favourable properties in conjunction with compositions which do not include Ni and other toxic elements [1,2]. In addition to biocompatibility, Ti alloys for surgical implants need to be mechanically compatible [3,4]. For orthopaedic implants high strength in conjunction with a low elastic modulus is critical to reducing detrimental ‘‘stress shielding’’ effects. For other implant applications factors such as the surface properties and fatigue resistance are of primary importance. With their potential for enhanced strengths and greater fatigue resistance there is considerable interest in the development of ultrafine grain (UFG) and nanocrystalline (NC) versions of these new metastable b titanium alloys [5–12]. A new metastable b Ti alloy, containing Nb, Zr, Mo and Sn, was recently reported by the authors for use in bio-implant applications [11–14]. The Ti–25Nb–3Zr–3Mo–2Sn alloy has a low modulus of elasticity, high strength and exhibits considerable plasticity and pseudoelastic character. Through the application of grain refinement via severe plastic deformation (SPD), using a modified accumulative roll bonding (ARB) technique, the strength of the alloy was significantly enhanced [11]. A 70% improvement
n Corresponding author at: Queensland Centre for Advanced Materials Processing and Manufacturing (AMPAM), School of Mechanical and Mining Engineering, The University of Queensland, Australia. Tel.: þ61 7 3365 3801; fax: þ15 71 7 3355 3888. E-mail address: [email protected] (D. Kent).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.07.030
in the ultimate tensile strength and a two-fold increase in the 0.5% proof strength of the alloy were achieved while maintaining levels of ductility of around 4.5%. TEM investigations have shown that the ARB processed b Ti–25Nb–3Zr–3Mo–2Sn alloy microstructure is primarily comprised of ultrafine b grains heavily elongated in the rolling direction with a fine dispersion of nanocrystalline a phase precipitates located on the b grain boundaries. The b grains typically have a flattened morphology with a grain size of approximately 130 nm when viewed normal to the rolling direction and a mean grain boundary interval of around 70 nm viewed from the transverse direction. Occasional large b grains in the order of 1–2 mm were also observed amongst the heavily refined structure and nanocrystalline grains with grain sizes in the order of 20 nm were present within shear bands. In order to further the development and application of the UFG metastable b Ti–25Nb–3Zr–3Mo–2Sn alloy it is important to investigate its thermal stability. Thermal treatments are often used in order to relieve residual stresses or to improve the balance between strength and ductility in the as-processed UFG b Ti alloys. However, fine, non-equilibrium microstructures can be prone to rapid grain growth at elevated temperatures, which may significantly impact on their properties and performance. Also for UFG materials produced by SPD rapid grain growth can occur at relatively low temperatures [15,16]. However, in some cases UFG materials achieve good thermal stability at higher temperatures due to the influence of impurities, pores or second phase particles due to grain boundary pinning effects [7,17–19]. Hence the objective of this research was to investigate the thermal stability of the UFG metastable b Ti–25Nb–3Zr–3Mo– 2Sn alloy microstructure.
D. Kent et al. / Materials Science & Engineering A 556 (2012) 582–587
583
2. Method
3. Results and discussion
The Ti–25Nb–3Zr–3Mo–2Sn (wt%) alloy was produced from commercially pure Ti sponge (99.5 wt%), pure Zr bars (99.7 wt%), pure Sn bars (99.9 wt%), pure Mo powder (99.8 wt%) and an intermediate Nb–47wt% Ti alloy. It was melted twice by nonconsumable arc melting to ensure chemical homogeneity and low levels of impurities (Fer0.10 wt%, Sir0.04 wt%, Cr0.03 wt%, Nr0.02 wt%, H r0.003 wt%). The ingot was forged and then hot rolled to a thickness of 1.0 mm at approximately 850 1C (above the b transus). The hot-rolled sheets were cleaned with ethanol, clad in Ti foil, solution treated under a protective Argon atmosphere at a temperature of 710 1C for 1 h and air cooled. Rolling was conducted on a 550 mm (diam.) 650 mm roll set using a roll speed of 10 ms 1. The 1.0 mm solution treated sheet was initially cold rolled at ambient temperature to a thickness of 0.2 mm. The sheet was then sectioned in half, stacked and rolled together. Stacking and rolling comprises a single ARB cycle and the process was repeated 3 times to give a total of 4 cycles, including the initial 80% reduction from 1 mm to 0.2 mm. A 50% total reduction was used for each of the ARB cycles, involving 3 individual roll passes with reductions of 30%, then 15% and finally 5%, such that each time the sheet was cut, stacked and rolled a constant 0.2 mm final sheet thickness was maintained. The sheet dimensions were 0.2 mm 150 mm 200 mm. The surface of the sheets was chemically cleaned using a HF:HNO3:H2O¼1:5:4 acid solution between each cycle. The surface of the Ti ARB sheet was protected during rolling by cladding with stainless steel sheet. The stainless steel clad ARB assembly was preheated to 400 1C for 20 min before rolling. Annealing heat treatments were undertaken on the UFG Ti– 25Nb–3Zr–3Mo–2Sn (wt%) alloy produced after 4 cycles of rolling. The annealing treatments were conducted in a MIHM-VOGT P6/B box furnace. Samples were encased in vacuum sealed quartz tubes and placed in the preheated furnace at the annealing temperature for times up to 180 min. Microhardness tests were performed on a Struers Duramin 20 Micro Vickers hardness tester using a 200 g load. Samples for transmission electron microscopy (TEM) were prepared using twin-jet electropolishing (Tenupol-3, Struers) with a solution containing HClO4 (5%), 2-butoxyethanol (35%) and methanol at 30 1C and 20 V. Some samples were further thinned using a Gatan precision ion polishing system (PIPS, model 691) with Arbeam energy of 4 KeV at an incident angle of 741. TEM was performed using a Philips Tecnai 20 FEG instrument.
A bright field (BF) TEM image of the alloy showing the heavily refined grain structure in the as-rolled condition is presented in Fig. 1(a). The grain boundaries are wavy and not well defined and there is extensive diffraction contrast both within the grains and at grain boundaries. These are features typical of UFG materials formed via SPD and are associated with the non-equilibrium state of the material associated with the high levels of strain and associated dislocation activity imparted to the material during processing [15,19,20]. A selected area diffraction pattern (SADP) from this region shown inset, shows a ring type pattern characteristic of very fine and randomly oriented grains with large grain boundary misorientations, which can be indexed to the b and a phases. Due to the indistinct grain structure of the UFG material under BF imaging conditions, a dark field imaging technique was developed to show contrast between the grains for grain size measurements. The technique is a derivative of hollow-cone dark field microscopy, a method in which an image is formed by summing the electron intensity at the image plane as a chosen annular region of the diffraction pattern circumferentially scans over the objective aperture [21]. A similar technique was employed in this study to obtain dark field (DF) images. However, rather than taking a single long exposure while continuously scanning the diffraction pattern over the objective aperture, a series of 3 or more DF images were taken from the same region of the sample using different portions of the diffraction pattern, selected by rotating the diffraction pattern over the objective aperture between each exposure. Coloured composite dark field (CDF) images were then created by combining a series of three of these DF images using the Analysis—colour mix function within the Digital Micrograph software. Fig. 1(b) is a CDF image of the same area shown in Fig. 1(a) formed using the DF images obtained using representative aperture positions indicated on the SADP (inset) by the coloured circles. The CDF image clearly reveals the grain structure which primarily consists of irregularly shaped b grains with an average grain size determined by the line intercept method of approximately 130 nm [11]. To assess the thermal stability of the UFG microstructure, changes in the properties were evaluated using microhardness measurements over the course of annealing treatments at temperatures ranging from 400 1C to 600 1C for times of up to 3 h. The results of these measurements are presented in Fig. 2. The microhardness of the as-rolled UFG b-Ti alloy was approximately
Fig. 1. TEM of the as-rolled microstructure (4 cycles). (a) BF image reveals a heavily refined grain structure. Inset: SADP exhibiting a ring type pattern typical of very fine and randomly oriented grains indexed to the b and a phases. (b) CDF image showing irregularly shaped b grains with an average grain size of approximately 130 nm. Inset: SADP with aperture positions represented by coloured circles used to form the CDF image.
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310 Hv. Annealing at 400 1C resulted in an increase in the hardness of around 50 Hv, while for the annealing treatment at 450 1C the hardness remained effectively constant. At 500 1C there was an initial rapid drop in the microhardness of around 30 Hv after 15 min followed by a more gradual decrease in the hardness with further annealing time. At 600 1C there was a considerable rapid reduction in the hardness of around 100 Hv within the first
Fig. 2. Change in microhardness over time for different annealing temperatures.
10 min of annealing, after which the hardness remains effectively constant at around 210 Hv. TEM analysis was undertaken in order to assess changes in the microstructure associated with the variations in microhardness detected during annealing. A BF TEM image of a sample annealed for 60 min at 400 1C is shown in Fig. 3(a). At this temperature an increase in the microhardness of around 50 Hv was observed. In comparison to the as-rolled microstructure which was shown in Fig. 1(a), there is reduced diffraction contrast associated with dislocation networks within grains and at the grain boundaries and the boundaries are more distinct. These features may be associated with the process of recovery involving the annihilation of dislocations at the grain boundary walls and via dipole interactions [22,23]. However, it should be noted that the appearance of these features can also be influenced by specific TEM imaging conditions such as the sample thickness [20]. Careful examination of the SADP, shown in Fig. 3(b), also reveals an increased intensity for reflections within the a phase diffraction rings, which is particularly apparent for the inner a (100) reflections. This may be attributed to coarsening of the a phase precipitates during annealing [14,24]. The proportion of a phase did not change considerably during annealing according to the results from the X-ray diffraction analysis, presumably due to its being near the equilibrium as a result of the multiple heating cycles undergone during processing as it was heated to 400 1C before each ARB pass [11]. A DF image formed using a portion of the inner a (100) diffraction ring is shown in Fig. 3(c). It reveals nanoscale and some coarser a precipitates
Fig. 3. TEM from a sample annealed at 400 1C for 60 min. (a) BF image showing reduced diffraction contrast associated with dislocation networks and boundaries which are more distinct in comparison to the as rolled condition. (b) SADP exhibiting a ring type pattern typical of very fine and randomly oriented grains which can be indexed to the b and a phases. (c) A DF formed using a portion of the inner a (100) diffraction ring reveals both nanoscale and some coarser a phase precipitates. (d) A CDF image of the region shown in (a).
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located on the grain boundaries of the predominantly b phase microstructure. Fig. 3(d) shows a CDF image of the region shown in the BF image (Fig. 3(a)). It reveals that the b grain size is of a similar order of magnitude to that of the as-rolled material and the wavy nature of the grain boundaries is preserved. A thorough analysis of the b grain size during annealing is presented in Fig. 7.
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A BF TEM image of a sample annealed for 10 min at 500 1C is shown in Fig. 4(a) and a CDF image of the same region is shown in Fig. 4(b). After this annealing treatment, there was a small decrease in the microhardness of around 20–30 Hv. Some grains exhibit well defined grain boundaries with a curved appearance as shown in the higher magnification image, Fig. 4(c), contrasting with the wavy appearance of grain boundaries in the as rolled
Fig. 4. TEM from a sample annealed at 500 1C for 10 min. (a) BF image reveals a heavily refined grain structure. (b) CDF image of the region shown in (a) formed using portions of the low order diffraction rings. Inset: SADP exhibiting exhibiting a ring type pattern which can be indexed to the b and a phases. (c) Higher magnification BF image showing well defined grain boundaries with a curved appearance. (d) BF image showing an example of an isolated pocket of lath type a phase precipitates which were observed in occassional relatively large b grains within the UFG microstructure.
Fig. 5. TEM from a sample annealed at 600 1C for 10 min. (a) BF image reveals comparatively large grains with straight sided grain boundaries characteristic of a recrystallised grain structure with SADP, inset, exhibiting diffraction spots and streaks which can be indexed to the b phase. (b) CDF image of the recrystallised grains formed using the low order b phase reflections.
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condition and after annealing at 400 1C. This suggests that the early stages of recrystallisation commence at the higher annealing temperature of 500 1C. Recovery leading to reductions in lattice microstrains and the early stages of recrystallisation may account for the decrease in hardness observed during annealing at 500 1C [23,25]. Examination of the SADP, inset Fig. 4(b), again reveals some increased intensity for reflections within the inner a (100) diffraction ring in comparison to the as-rolled material due to coarsening of the a phase during annealing. There were also isolated pockets of lath type a phase precipitates, shown in Fig. 4(d), which are of a similar morphology to those observed in the coarse grained form of the alloy after aging treatments, and which in this case were observed in the occasional relatively large b grains present within the predominantly UFG microstructure [14]. A BF TEM image of a sample annealed for 10 min at 600 1C is shown in Fig. 5(a). The b grains show a markedly changed equiaxed structure with distinct and predominantly straight grain
boundaries. This annealing treatment resulted in a substantial decrease in the microhardness of around 100 Hv and it is clear that this is associated with recrystallisation of the b phase and subsequent grain growth. The SADP, shown inset in Fig. 5(a), is no longer in the form of continuous diffraction rings and the reflections, which can be indexed to the b phase, are in the form of spots or streaks due to the larger b grain size. The CDF image of the region depicted in Fig. 5(a) is shown in Fig. 5(b), showing the b grains with an average grain size in the order of 0.5–1 mm. An analysis of the b grain size was undertaken from multiple CDF images for each annealing condition using the line intercept method. Examples of the CDF images were presented in Figs. 1, 3–5 and a series of BF–CDF image pairs showing the evolution of the grain size over the course annealing at 500 1C is presented in Fig. 6. They show the gradual increase in the b grain size during annealing at this temperature. The results of b grain size measurements over the course of annealing at temperatures of 400 1C,
Fig. 6. Series of TEM BF–CDF image pairs showing the evolution of the grain size over the course annealing at 500 1C. (a) 10 min, (b) 30 min and (c) 60 min.
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Fig. 7. An analysis of the b grain size versus annealing time for annealing temperatures of 400 1C, 500 1C and 600 1C undertaken from CDF images using the line intercept method.
500 1C and 600 1C are presented in Fig. 7. It shows that the grain size did not change significantly during annealing at 400 1C and increased gradually during annealing at 500 1C. However, annealing at 600 1C led to a rapid increase in the average grain size within the first 10 min to approximately 0.72 mm, after which it remains effectively constant. The TEM analysis has shown that coarsening of the a phase and the process of recovery, associated with the annihilation of dislocations and reduction of lattice strains, occur simultaneously in the temperature range of 400–500 1C. Significant recrystallisation of the b phase and associated grain growth does not occur until annealing temperatures of 600 1C and above. Despite the recrystallisation of the b phase occurring during annealing at 600 1C, an UFG structure was still retained after 60 min at this temperature. The UFG b Ti alloys show considerably greater thermal stabilty than other UFG Ti alloys produced from commercially pure titanium, for which significant grain growth is observed after annealing at temperatures of 450 1C and above [16]. These results are consistent with the observations of Li et al [7] for a similar nanostructured b alloy, Ti– 24Nb–4Zr–7.9Sn, which showed that significant growth of the b grains to conventional grain sizes larger than 10 mm does not occur until annealing temperatures of 700 1C or above. Nanocrystalline a phase precipitates, which in this case are present after ARB processing, coarsen during annealing and are believed to enhance the thermal stability of the microstructures in comparison to that of other nanostructured or UFG Ti alloys with single phase compositions [7]. Furthermore, the effects of solute drag may also play a more prominent role in retarding the rate of grain growth in the heavily alloyed metastable b Ti alloys in comparison to that of the commercially pure Ti alloys [26,27]. Further work to clarify the role of solute drag and the evolution of the a phase during annealing and their influence on thermal stability will be published in the future. An important feature of metastable UFG b Ti alloys, such as the Ti–25Nb–3Zr–3Mo–2Sn alloy, is their capacity to retain an UFG grain structure at relatively high annealing temperatures around 600 1C, while undergoing recrystallisation of the b phase. It permits the use of thermal annealing treatments to relieve residual stresses and anisotropy of properties imparted to the material during processing, while preserving the potential benefits of UFG Ti alloys such as an enhanced in vitro fibroblast response or improved mechanical properties [5,6].
treatments. Coarsening of the a phase occurred during annealing at temperatures of 400–500 1C in conjunction with recovery involving the annihilation of dislocations in the b phase. Minimal changes were observed in the hardness at these annealing temperatures, with a small increase in the hardness at 400 1C and a small decrease at 500 1C. There was no change in the b grain size after prolonged annealing treatments at 400 1C and limited grain growth after prolonged periods at 500 1C. Recrystallisation of the b phase and substantial grain growth occurred at an annealing temperature of 600 1C, which was associated with significant reductions in the hardness. However, a UFG grain structure with average grain size around 0.72 mm was still retained after prolonged annealing at this temperature. The UFG metastable b Ti alloys have significantly greater thermal stability than commercial purity UFG Ti alloys.
Acknowledgements The authors acknowledge the facilities, and the scientific and technical assistance, of the Australian Microscopy and Microanalysis Research Facility at the Centre for Microscopy and Microanalysis, The University of Queensland.
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4. Conclusions [26]
The thermal stability of an UFG metastable b Ti alloy, Ti– 25Nb–3Zr–3Mo–2Sn, was assessed during annealing heat
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