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Improvement of room temperature ductility for Mo and Fe modified Ti2AlNb alloy
Materials Science and Engineering A 528 (2010) 355–362
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Improvement of room temperature ductility for Mo and Fe modified Ti2 AlNb alloy Satoshi Emura a,∗ , Kaneaki Tsuzaki a,b , Koichi Tsuchiya a,b a b
National Institute for Materials Science, Tsukuba 305-0047, Japan Graduate School of Pure and Applied Sciences, University of Tsukuba, Ibaraki 305-0047, Japan
a r t i c l e
i n f o
Article history: Received 24 May 2010 Received in revised form 1 September 2010 Accepted 3 September 2010
Keywords: Orthorhombic phase Thermomechanical treatment VGS structure Tensile elongation Segregation
a b s t r a c t The effect of hot bar rolling and a subsequent annealing in the (B2 + ␣2 ) two-phase region on the mechanical properties was investigated for the Ti–25Al–14Nb–2Mo–1Fe (mol%) orthorhombic phase base alloy. After this thermomechanical treatment in the (B2 + ␣2 ) two-phase region, a ‘Van Gogh’s Sky (VGS)’ structure, with wavy and curled bands of spherical ␣2 precipitates, was obtained. Electron probe micro-analysis (EPMA) examinations indicated that the VGS bands correspond to the Nb and Mo lean regions. The width and spacing of VGS bands varied with thermomechanical treatment conditions (forging temperature before hot bar rolling as well as annealing temperature after the rolling). The samples with the VGS structure exhibited the higher tensile elongation-to-failure at room temperature. The creep resistance at 923 K decreased with the existence of the VGS structure. These mechanical properties were affected by the width and spacing of the VGS bands. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Ti–Al–Nb alloy based on Ti2 AlNb intermetallic with the ordered orthorhombic phase (O phase) [1,2] offers a higher strengthto-weight ratio, better creep resistance, and better workability than conventional titanium aluminides such as TiAl-based and Ti3 Al-based alloys [3–7]. For these qualities, Ti2 AlNb-based alloys have attracted much attention as potential materials for advanced aerospace applications such as aircraft engine parts [6,7]. One prototypical alloy is a Ti–22Al–27Nb (mol%), which consists mainly of a (O + B2 (CsCl)) two-phase microstructure [3–5]. Ti–22Al–27Nb alloy shows well-balanced mechanical properties. For example, Rowe reported that the Ti–22Al–27Nb alloy solution treated at 1273 K followed by aging at 1033 K exhibits a high tensile strength (TS) of around 1080 MPa and tensile elongation of about 5% at room temperature, and TS of above 650 MPa and elongation of above 16% when tested at 1033 K [4]. One major drawback of Ti–22Al–27Nb alloy is its extremely high Nb content, which corresponds to around 45 mass% exceeding the Ti content. Nb as additional element has some problems such as high cost and high density. Efforts have been made to reduce the Nb content of Ti–22Al–27Nb alloy by substituting Nb with other alloying elements, such as W and Mo [8,9]. The present authors have found that multiple additions of Mo and Fe are effective to
∗ Corresponding author at: Hybrid Materials Center, National Institute for Materials Science, 1-2-1 Sengen Tsukuba, Ibaraki 305-0047, Japan. Tel.: +81 29 859 2529; fax: +81 29 859 2101. E-mail address: [email protected] (S. Emura). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.09.003
reduce Nb content and also to improve tensile and creep strength at elevated temperatures up to 923 K [10]. This Mo and Fe modified alloy, however, exhibits poor room temperature tensile elongation of less than 1%. Reduction of the grain size of the prior B2 phase aids improvement of the room temperature elongation-to-failure of Ti2 AlNb alloys. We have succeeded in controlling and refining the prior B2 grain size of Ti–22Al–27Nb alloy processed by powder metallurgy using the pinning effect of ␣2 (D019 ) particles homogeneously dispersed in the B2 matrix during hot bar rolling and subsequent annealing in the (B2 + ␣2 ) two-phase region [11]. This fine-grained material with the average prior B2 grain size of 8 m exhibited a room temperature tensile elongation of more than 15%. In the present study, we have applied this thermomechanical treatment to the Mo and Fe modified alloy produced by levitation melting, and obtained a unique microstructure as shown in Fig. 1. Instead of a homogeneous distribution of the ␣2 phase particles in the above-mentioned powder metallurgy Ti–22Al–27Nb alloy, the distribution was heterogeneous forming wavy and curly bands when viewed from the rolling direction in these samples. Similar microstructures have been observed by other researchers [12–15], especially in Nb-containing B2 phase-based alloys. They referred to this microstructure as ‘Van Gogh’s Sky (VGS)’ structure because it resembles Van Gogh’s sky paintings. It was also reported that Ti–Nb–Al and Ti–Zr-Nb–Al alloys with the VGS structure exhibited larger tensile elongation-to-failures at room temperature [13]. So it is expected that the room temperature tensile elongation-tofailure of Mo and Fe modified Ti2 AlNb alloys can be improved when processed into this VGS structure. Meanwhile, because Ti2 AlNb alloys are expected to be high-temperature materials, the high-
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Fig. 1. Typical OM image of Ti–25Al–14Nb–2Mo–1Fe with Van Gogh’s Sky microstructure. (Homogenized, hot forged, hot rolled and annealed at 1273 K followed by furnace cooling at 0.03 K/s.)
temperature mechanical properties such as creep resistance are also important. In the present study, we investigated the effect of the VGS structure on the room temperature tensile properties and the high-temperature creep properties of the Mo and Fe modified Ti2 AlNb alloy. 2. Experimental The Ti–25Al–14Nb–2Mo–1Fe (mol%) alloy was prepared by the cold crucible levitation melting. Each ingot with around 70 mm diameter and around 60 mm length had a weight of around 1.2 kg. The ingots were hot forged and then hot rolled into 11.8 mm square bars followed by annealing for 3.6 ks. We have prepared samples with three different thermomechanical treatments. The thermomechanical treatments are schematically shown in Fig. 2. For the sample A, the ingot was hot forged, hot bar rolled and annealed at 1373 K in the B2 single phase region. For the samples B and C, the forged bar was hot bar rolled and annealed in the (B2 + ␣2 ) two-phase region. The difference between samples B and C was
Fig. 3. X-ray diffraction profiles of Ti–25Al–14Nb–2Mo–1Fe with different processing conditions.
the forging temperature before the hot bar rolling. For sample B, all steps of the thermomechanical treatment were performed at 1273 K in the (B2 + ␣2 ) two-phase region. For sample C, the forging was performed at 1523 K in the B2 single phase region, followed by bar rolling at 1273 K. The annealing temperature for sample C was 1293 K, which is 20 K higher than that for sample B but both are in the (B2 + ␣2 ) two-phase region. In all the treatments, the hot bar rolling was performed for 16 passes, and the samples were reheated after every pass. For evaluating the microstructure and mechanical properties, the annealed materials were all furnace-cooled (cooling rate: 0.03 K/s) to room temperature and then aged at 1073 K in the (O + B2) two-phase region for 360 ks followed by air-cooling to stabilize the microstructure. The phase identification was made by an X-ray diffraction (XRD) analysis using a Cu-K␣ radiation operated at 40 kV–300 mA. X-ray profiles were taken on the planes normal to the rolling direction (RD plane, as shown in Fig. 2). Microstructural observations were performed on an optical microscope (OM) and a scanning electron
Fig. 2. A schematic drawing of the thermomechanical treatment for Ti–25Al–14Nb–2Mo–1Fe. CCLM, AC, TD and RD refer to cold crucible levitation melting, air-cooling, transverse direction and rolling direction, respectively.
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Fig. 4. OM images of Ti–25Al–14Nb–2Mo–1Fe observed from the RD planes ((a), (c), and (e)) and the TD planes ((b), (d), and (f)) for different samples, A ((a) and (b)), B ((c) and (d)) and C ((e) and (f)).
microscope (SEM). Both the RD planes and the planes parallel to the rolling direction (TD plane, as shown in Fig. 2) were observed. Mechanically polished specimens for OM and SEM observations were etched using a solution of distilled water, nitric acid and hydrofluoric acid (90:8:2 in volume). To measure the volume fraction of the ␣2 phase particles, the SEM images were processed using image processing software (Image J). The elemental distributions were mapped out by the electron probe micro-analysis (EPMA) with a spacial resolution of 1 m. Room-temperature tensile tests were carried out at a cross-head speed of 0.3 mm/min (strain rate: 3 × 10−4 s−1 ) in air. The stress was measured by a load
cell (capacity: 30 kN). The strain was calculated from the crosshead displacement, and the total elongation value was obtained after eliminating the elastic deformation amount. Specimens for the tensile test had a 3.5 mm diameter and a 16 mm gauge length. Deformed and fractured samples were examined by SEM. Creep tests were conducted in air under a constant tensile load using a dead-weight creep rupture machine. Creep specimens had a 4 mm diameter and a 25 mm effective gauge length. The temperature and initial applied stress were 923 K and 310 MPa, respectively. Two thermocouples to measure the specimen temperature were attached at the upper and lower part of the gauge section, and the
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average of the two measured values was adopted as the specimen temperature.
3. Results and discussion Fig. 3 shows the X-ray diffraction profiles of the Ti–25Al–14Nb–2Mo–1Fe alloy. These profiles were taken on the RD plane. All the samples consist of O, B2 and ␣2 phase. In sample C, the amount of the ␣2 phase was a little smaller compared with that in sample B. This could be due to the higher annealing temperature after the rolling. Fig. 4 illustrates the optical micrographs of the Ti–25Al–14Nb–2Mo–1Fe alloy after the different working conditions. In sample A, processed in the B2 single phase region, large equiaxed grains with grain boundary ␣2 phase particles were observed. And there is no significant difference in the microstructures between those observed on the RD plane (Fig. 4(a)) and the TD plane (Fig. 4(b)). Meanwhile, in samples B and C, processed in the (B2 + ␣2 ) two-phase region, the so-called VGS structure was seen. The RD plane exhibited wavy and curled bands of ␣2 phase particles as seen in Fig. 4(c). These ␣2 phase bands roughly align along the rolling direction, as observed in the TD plane (Fig. 4(d)). There were some differences in the distributions of ␣2 phase particles between samples B and C. In sample C, with a higher forging temperature of 1523 K and a higher annealing temperature of 1293 K, the widths of the ␣2 phase bands were somewhat larger and the spacing of the bands was also larger. Fig. 5 illustrates magnified SEM images of the Ti–25Al–14Nb–2Mo–1Fe alloy. All images were taken on the RD planes. In sample A (Fig. 5(a)), the ␣2 phase can be seen with the dark contrast along the grain boundaries. Inside the grains, fine acicular O phase (the gray contrast region) precipitated in the B2 phase matrix (the light contrast region). In sample B (Fig. 5(b)), curled VGS bands, which consisted of spherical ␣2 precipitates with a diameter of between 2 and 5 m, were observed. Between the bands, there was a fine (O + B2) two-phase microstructure similar to the one in sample A. Although the microstructural features were similar in sample C (Fig. 5(c)), the dispersed ␣2 precipitates were finer and more angular in shape. By processing the low-magnification SEM images using the image analysis software, the volume fraction of dispersed ␣2 precipitates was determined to be 20.6 ± 1.0% and 18.5 ± 0.7% in samples B and C, respectively. The ␣2 phase fraction was smaller in sample C, which is in agreement with the X-ray diffraction results. Fig. 6 shows the backscattered electron images (Fig. 6(a) and (b)) and elemental distribution maps for Nb (Fig. 6(c) and (d)) and Mo (Fig. 6(e) and (f)) obtained by EPMA for samples B and C. The VGS bands of the spherical ␣2 phase particles seen in Fig. 6(a) and (b) coincided with Nb and Mo lean regions in Fig. 6(c)–(f). Naka et al. reported [13] that plastic instabilities arise from a region-toregion variation in hardness during the high-temperature extrusion of Nb-containing B2 phase-based alloys. He insisted that this variation originated from the dendritic segregation in the starting ingot, and the metallographic contrasts of the VGS structure were related to the local variation in chemical composition. We also have confirmed the segregation of alloying elements (especially Nb and Mo) with the spacing of around 100 m in as consolidated ingot using energy dispersive X-ray spectroscopy. In the present study, the Nb-containing alloy was processed by hot bar rolling. Since the deformation state of hot bar rolling is similar to that of the hot extrusion, VGS structure related to the local segregation can be formed during the hot bar rolling. So it is supposed that the dendritic segregation in the ingot, which is coincident with the region-to-region variation in hardness, has evolved into wavy and curled bands of Nb and Mo lean regions due to the inhomoge-
Fig. 5. SEM images of Ti–25Al–14Nb–2Mo–1Fe. (a) Sample A, (b) sample B, and (c) sample C. The observed plane is normal to the rolling direction.
neous deformation during the hot bar rolling. The bar rolling was performed at 1273 K in the (B2 + ␣2 ) two-phase region, and thus precipitation of the ␣2 phase particles must be occurred simultaneously. Nb and Mo are both the B2 phase stabilizing elements, so the ␣2 phase preferably precipitated at the Nb and Mo lean regions. Finally, during the annealing in the (B2 + ␣2 ) two-phase region (at 1273 K for sample B and at 1293 K for sample C, respectively), the precipitated ␣2 phase particles grew spherically and VGS-like ␣2 phase bands were constructed. We have applied the hot bar rolling in the (B2 + ␣2 ) two-phase region to the Ti–22Al–27Nb alloy processed by powder metallurgy [11]. In this case, we obtained a homogeneous distribution of the ␣2 phase instead of the VGSlike structure. We have used the gas-atomized (i.e. rapidly cooled) Ti–22Al–27Nb alloy powder with a diameter of less than 150 m as the starting powder. Even if the elemental segregation existed in the gas-atomized powder, the spacing of the segregation must be very fine. Furthermore, these powders were consolidated using hot isostatic pressing at high temperature (1373 K) for long period (7.2 ks). Therefore, elemental segregation in this powder metal-
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Fig. 6. EPMA mappings for Nb and Mo distribution in sample B ((a), (c), and (e)) and sample C ((b), (d), and (f)) of Ti–25Al–14Nb–2Mo–1Fe. (a) and (b) backscattered electron images. (c, d) and (e, f) are the Nb and Mo maps, respectively. Concentrations are expressed by mass%. The observed plane is normal to the rolling direction.
lurgy processed material was significantly small, and VGS-like structure seems to be unfavorable. The reason why sample C exhibited wider VGS bands and wider spacing could also be explained by the segregation of Nb and Mo. Sample C was forged at a much higher temperature (1523 K) than sample B before hot bar rolling. At this temperature, the diffusion
rates of Nb and Mo are much higher. Using the diffusion coefficients of these atoms in -Ti [16], the diffusion rates of Nb and Mo are estimated to be around four times higher than those at 1273 K. So, after the forging, the degree of segregation in sample C, which was forged at 1523 K, was supposed to become lower compared with that in sample B, which was forged at 1273 K. This means that the
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Fig. 7. Room temperature tensile stress–strain curves of Ti–25Al–14Nb–2Mo–1Fe.
local variation of hardness in sample C before bar rolling became smaller than that in sample B. Therefore, it is supposed that the VGS-like segregation in sample C became less significant and the width and spacing of the VGS bands became wider after the hot bar rolling. Room temperature tensile stress–strain curves of the Ti–25Al–14Nb–2Mo–1Fe alloy are shown in Fig. 7. In Fig. 7, tensile strain was calculated from the cross-head displacement. The values of 0.2% proof stress (YS), tensile strength (TS) and total elongation (EL) are summarized in Table 1. All the samples were broken without necking. Sample A (without VGS structure) exhibited the lowest EL value of all the microstructures we have evaluated [10]. On the contrary, samples B and C (with VGS structure) exhibited higher EL values. Sample B exhibited the largest EL value of 6.3% along with the highest YS (705 MPa) and TS (992 MPa). This EL value is comparable to that of the Ti–22Al–27Nb alloy solution treated at 1273 K followed by aging at 1033 K [4]. Meanwhile, the EL value of sample C was 1.7%, which is higher than that of sample A. Fig. 8(a)–(c) shows the SEM images of the tensile fracture surfaces of samples A–C, respectively. The SEM images of the cross section areas beneath the tensile fracture surfaces are shown in Fig. 9. Sample A (without VGS structure) exhibited a coarse brittle fracture surface. Large cracks occurred mainly inside the grains. These large cracks resulted in the poor room temperature El value of sample A. On the other hand, sample B exhibited a much finer and dimple-like fracture surface. Beneath the fracture surface, a large number of small subcracks were observed. These subcracks were often arrested and/or deflected at the ␣2 phase bands, as shown in Fig. 9(d). These crack arrest and crack deflection seem to have prevented premature fracture and enhance the EL values. Sample C exhibited another type of fracture features. The fracture surface was a little coarser compared to that of sample B, and deflected and/or arrested subcracks were not observed as often as in sample B. Coarser fracture surface and smaller number of the small sub-
Table 1 Room temperature tensile properties of the Ti–25Al–14Nb–2Mo–1Fe alloy with different processing conditions. YS, TS and EL refer to 0.2% proof stress, tensile strength and total elongation, respectively.
Sample A Sample B Sample C
YS (MPa)
TS (MPa)
EL (%)
– 705 603
781 992 765
0.1 6.3 1.7
Fig. 8. SEM images of the fracture surfaces of Ti–25Al–14Nb–2Mo–1Fe after tensile tests. (a) Sample A, (b) sample B, and (c) sample C.
cracks suggest that the crack propagation resistance of sample C was lower than that of sample B. Therefore the EL value of sample C seems to be lower than that of sample B. Even though some long cracks parallel to the tensile direction were also observed in sample C, these cracks do not seem to be directly related to the lower EL values of sample C. Rather, these long parallel cracks indicate the mechanical anisotropy of sample C. Fig. 10 shows the creep curves of samples A–C of the Ti–25Al–14Nb–2Mo–1Fe alloy tested at 923 K under the initial applied stress of 310 MPa. As a reference, the result of Ti–22Al–27Nb alloy with similar microstructure to the sample A [10] is also shown in Fig. 10. Contrary to the room temperature EL values, the samples with the VGS structure exhibited lower creep resistance. This reduction in the creep resistance seems to be due to the fine VGS microstructure. However all the specimens of Ti–25Al–14Nb–2Mo–1Fe alloy exhibited a higher creep resistance compared with that of the Ti–22Al–27Nb alloy.
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Fig. 9. SEM images of transverse sections of the tensile tested specimens of Ti–25Al–14Nb–2Mo–1Fe. Tensile axis is parallel to the horizontal direction. (a) and (b) sample A, (c) and (d) sample B, (e) and (f) sample C.
4. Conclusions To improve the room temperature elongation-to-failure of the Ti–25Al–14Nb–2Mo–1Fe (mol%) orthorhombic phase base alloys, hot bar rolling and a subsequent annealing were performed in the (B2 + ␣2 ) two-phase region.
Fig. 10. Creep curves of Ti–25Al–14Nb–2Mo–1Fe with and without VGS structure. Samples were tested at 923 K under the initial applied stress of 310 MPa. The result of Ti–22Al–27Nb [10] is also shown for comparison.
(1) After the thermomechanical treatment in the (B2 + ␣2 ) twophase region, a ‘Van Gogh’s Sky (VGS)’ structure, with wavy and curled bands of spherical ␣2 precipitates, was obtained. (2) The VGS bands corresponded to the Nb and Mo lean regions. The formation of this microstructure may be related to the dendritic segregation of Nb and Mo during the solidification. This segregation has evolved into wavy and curled bands of Nb and Mo lean regions due to the inhomogeneous deformation during the hot bar rolling. ␣2 precipitation in the Nb and Mo lean regions also occurred during forging and hot bar rolling.
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(3) The sample forged at the higher temperature (1523 K) before bar rolling and annealed at the higher temperature (1293 K) after bar rolling exhibited wider VGS bands and wider spacing. (4) The samples with the VGS structure exhibited higher tensile elongation-to-failure at room temperature. The creep resistance at 923 K decreased with the existence of the VGS structure. These mechanical properties were affected by the width and spacing of the VGS bands. References [1] D. Banerjee, A.K. Gogia, T.K. Nandy, V.A. Joshi, Acta Metall. 36 (1988) 871– 882. [2] K. Muraleedharan, T.K. Nandy, D. Banerjee, S. Lele, Intermetallics 3 (1995) 187–199. [3] R.G. Rowe, in: Y.-W. Kim, R.R. Boyer (Eds.), Microstructure/Property Relationships in Titanium Aluminides and Alloys, TMS, Warrendale, 1991, pp. 387– 398.
[4] R.G. Rowe, D.G. Konitzer, A.P. Woodfield, J.C. Chesnutt, in: L.A. Johnson, D.P. Pope, J.O. Stiegler (Eds.), High-Temperature Ordered Intermetallic Alloys IV, Materials Research Society, Pittsburgh, 1991, pp. 703–708. [5] C.J. Boehlert, B.S. Majumdar, V. Seetharaman, D.B. Miracle, Mater. Trans. A 30 (1999) 2305–2323. [6] J. Kumpfert, C. Leyens, in: M.V. Nathal, R. Darolia, C.T. Liu, P.L. Martin, D.B. Miracle, R. Wagner, M. Yamaguchi (Eds.), Structural Intermetallics 1997, TMS, Warrendale, 1996, pp. 895–904. [7] J. Kumpfert, W.A. Kaysser, Z. Metallkd. 92 (2001) 128–134. [8] F. Tang, S. Nakazawa, M. Hagiwara, Mater. Sci. Eng. A 329–331 (2002) 492–498. [9] J. Cao, F. Bai, Z. Li, Mater. Sci. Eng. A 424 (2006) 47–52. [10] Y. Mao, M. Hagiwara, S. Emura, Scripta Mater. 57 (2007) 261–264. [11] S. Emura, A. Araoka, M. Hagiwara, Scripta Mater. 48 (2003) 629–634. [12] M. Thomas, S. Naka, T. Khan, in: P.A. Blenkinsop, W.J. Evans, H.M. Flower (Eds.), Titanium’95: Science and Technology, The Institute of Materials, London, 1996, pp. 388–395. [13] S. Naka, M. Marty, M. Thomas, T. Khan, Mater. Sci. Eng. A 192/193 (1995) 69–76. [14] E.A. Loria, Mater. Sci. Eng. A 271 (1999) 430–438. [15] S.M. Allameh, R.W. Hayes, M. Li, E.A. Loria, D.J. Srolovitz, W.O. Soboyejo, Mater. Sci. Eng. A 328 (2002) 122–132. [16] G.B. Gibbs, D. Graham, D.H. Tomlin, Philos. Mag. 8 (1963) 1269–1282.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Improvement of room temperature ductility for Mo and Fe modified Ti2 AlNb alloy Satoshi Emura a,∗ , Kaneaki Tsuzaki a,b , Koichi Tsuchiya a,b a b
National Institute for Materials Science, Tsukuba 305-0047, Japan Graduate School of Pure and Applied Sciences, University of Tsukuba, Ibaraki 305-0047, Japan
a r t i c l e
i n f o
Article history: Received 24 May 2010 Received in revised form 1 September 2010 Accepted 3 September 2010
Keywords: Orthorhombic phase Thermomechanical treatment VGS structure Tensile elongation Segregation
a b s t r a c t The effect of hot bar rolling and a subsequent annealing in the (B2 + ␣2 ) two-phase region on the mechanical properties was investigated for the Ti–25Al–14Nb–2Mo–1Fe (mol%) orthorhombic phase base alloy. After this thermomechanical treatment in the (B2 + ␣2 ) two-phase region, a ‘Van Gogh’s Sky (VGS)’ structure, with wavy and curled bands of spherical ␣2 precipitates, was obtained. Electron probe micro-analysis (EPMA) examinations indicated that the VGS bands correspond to the Nb and Mo lean regions. The width and spacing of VGS bands varied with thermomechanical treatment conditions (forging temperature before hot bar rolling as well as annealing temperature after the rolling). The samples with the VGS structure exhibited the higher tensile elongation-to-failure at room temperature. The creep resistance at 923 K decreased with the existence of the VGS structure. These mechanical properties were affected by the width and spacing of the VGS bands. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Ti–Al–Nb alloy based on Ti2 AlNb intermetallic with the ordered orthorhombic phase (O phase) [1,2] offers a higher strengthto-weight ratio, better creep resistance, and better workability than conventional titanium aluminides such as TiAl-based and Ti3 Al-based alloys [3–7]. For these qualities, Ti2 AlNb-based alloys have attracted much attention as potential materials for advanced aerospace applications such as aircraft engine parts [6,7]. One prototypical alloy is a Ti–22Al–27Nb (mol%), which consists mainly of a (O + B2 (CsCl)) two-phase microstructure [3–5]. Ti–22Al–27Nb alloy shows well-balanced mechanical properties. For example, Rowe reported that the Ti–22Al–27Nb alloy solution treated at 1273 K followed by aging at 1033 K exhibits a high tensile strength (TS) of around 1080 MPa and tensile elongation of about 5% at room temperature, and TS of above 650 MPa and elongation of above 16% when tested at 1033 K [4]. One major drawback of Ti–22Al–27Nb alloy is its extremely high Nb content, which corresponds to around 45 mass% exceeding the Ti content. Nb as additional element has some problems such as high cost and high density. Efforts have been made to reduce the Nb content of Ti–22Al–27Nb alloy by substituting Nb with other alloying elements, such as W and Mo [8,9]. The present authors have found that multiple additions of Mo and Fe are effective to
∗ Corresponding author at: Hybrid Materials Center, National Institute for Materials Science, 1-2-1 Sengen Tsukuba, Ibaraki 305-0047, Japan. Tel.: +81 29 859 2529; fax: +81 29 859 2101. E-mail address: [email protected] (S. Emura). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.09.003
reduce Nb content and also to improve tensile and creep strength at elevated temperatures up to 923 K [10]. This Mo and Fe modified alloy, however, exhibits poor room temperature tensile elongation of less than 1%. Reduction of the grain size of the prior B2 phase aids improvement of the room temperature elongation-to-failure of Ti2 AlNb alloys. We have succeeded in controlling and refining the prior B2 grain size of Ti–22Al–27Nb alloy processed by powder metallurgy using the pinning effect of ␣2 (D019 ) particles homogeneously dispersed in the B2 matrix during hot bar rolling and subsequent annealing in the (B2 + ␣2 ) two-phase region [11]. This fine-grained material with the average prior B2 grain size of 8 m exhibited a room temperature tensile elongation of more than 15%. In the present study, we have applied this thermomechanical treatment to the Mo and Fe modified alloy produced by levitation melting, and obtained a unique microstructure as shown in Fig. 1. Instead of a homogeneous distribution of the ␣2 phase particles in the above-mentioned powder metallurgy Ti–22Al–27Nb alloy, the distribution was heterogeneous forming wavy and curly bands when viewed from the rolling direction in these samples. Similar microstructures have been observed by other researchers [12–15], especially in Nb-containing B2 phase-based alloys. They referred to this microstructure as ‘Van Gogh’s Sky (VGS)’ structure because it resembles Van Gogh’s sky paintings. It was also reported that Ti–Nb–Al and Ti–Zr-Nb–Al alloys with the VGS structure exhibited larger tensile elongation-to-failures at room temperature [13]. So it is expected that the room temperature tensile elongation-tofailure of Mo and Fe modified Ti2 AlNb alloys can be improved when processed into this VGS structure. Meanwhile, because Ti2 AlNb alloys are expected to be high-temperature materials, the high-
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Fig. 1. Typical OM image of Ti–25Al–14Nb–2Mo–1Fe with Van Gogh’s Sky microstructure. (Homogenized, hot forged, hot rolled and annealed at 1273 K followed by furnace cooling at 0.03 K/s.)
temperature mechanical properties such as creep resistance are also important. In the present study, we investigated the effect of the VGS structure on the room temperature tensile properties and the high-temperature creep properties of the Mo and Fe modified Ti2 AlNb alloy. 2. Experimental The Ti–25Al–14Nb–2Mo–1Fe (mol%) alloy was prepared by the cold crucible levitation melting. Each ingot with around 70 mm diameter and around 60 mm length had a weight of around 1.2 kg. The ingots were hot forged and then hot rolled into 11.8 mm square bars followed by annealing for 3.6 ks. We have prepared samples with three different thermomechanical treatments. The thermomechanical treatments are schematically shown in Fig. 2. For the sample A, the ingot was hot forged, hot bar rolled and annealed at 1373 K in the B2 single phase region. For the samples B and C, the forged bar was hot bar rolled and annealed in the (B2 + ␣2 ) two-phase region. The difference between samples B and C was
Fig. 3. X-ray diffraction profiles of Ti–25Al–14Nb–2Mo–1Fe with different processing conditions.
the forging temperature before the hot bar rolling. For sample B, all steps of the thermomechanical treatment were performed at 1273 K in the (B2 + ␣2 ) two-phase region. For sample C, the forging was performed at 1523 K in the B2 single phase region, followed by bar rolling at 1273 K. The annealing temperature for sample C was 1293 K, which is 20 K higher than that for sample B but both are in the (B2 + ␣2 ) two-phase region. In all the treatments, the hot bar rolling was performed for 16 passes, and the samples were reheated after every pass. For evaluating the microstructure and mechanical properties, the annealed materials were all furnace-cooled (cooling rate: 0.03 K/s) to room temperature and then aged at 1073 K in the (O + B2) two-phase region for 360 ks followed by air-cooling to stabilize the microstructure. The phase identification was made by an X-ray diffraction (XRD) analysis using a Cu-K␣ radiation operated at 40 kV–300 mA. X-ray profiles were taken on the planes normal to the rolling direction (RD plane, as shown in Fig. 2). Microstructural observations were performed on an optical microscope (OM) and a scanning electron
Fig. 2. A schematic drawing of the thermomechanical treatment for Ti–25Al–14Nb–2Mo–1Fe. CCLM, AC, TD and RD refer to cold crucible levitation melting, air-cooling, transverse direction and rolling direction, respectively.
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Fig. 4. OM images of Ti–25Al–14Nb–2Mo–1Fe observed from the RD planes ((a), (c), and (e)) and the TD planes ((b), (d), and (f)) for different samples, A ((a) and (b)), B ((c) and (d)) and C ((e) and (f)).
microscope (SEM). Both the RD planes and the planes parallel to the rolling direction (TD plane, as shown in Fig. 2) were observed. Mechanically polished specimens for OM and SEM observations were etched using a solution of distilled water, nitric acid and hydrofluoric acid (90:8:2 in volume). To measure the volume fraction of the ␣2 phase particles, the SEM images were processed using image processing software (Image J). The elemental distributions were mapped out by the electron probe micro-analysis (EPMA) with a spacial resolution of 1 m. Room-temperature tensile tests were carried out at a cross-head speed of 0.3 mm/min (strain rate: 3 × 10−4 s−1 ) in air. The stress was measured by a load
cell (capacity: 30 kN). The strain was calculated from the crosshead displacement, and the total elongation value was obtained after eliminating the elastic deformation amount. Specimens for the tensile test had a 3.5 mm diameter and a 16 mm gauge length. Deformed and fractured samples were examined by SEM. Creep tests were conducted in air under a constant tensile load using a dead-weight creep rupture machine. Creep specimens had a 4 mm diameter and a 25 mm effective gauge length. The temperature and initial applied stress were 923 K and 310 MPa, respectively. Two thermocouples to measure the specimen temperature were attached at the upper and lower part of the gauge section, and the
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average of the two measured values was adopted as the specimen temperature.
3. Results and discussion Fig. 3 shows the X-ray diffraction profiles of the Ti–25Al–14Nb–2Mo–1Fe alloy. These profiles were taken on the RD plane. All the samples consist of O, B2 and ␣2 phase. In sample C, the amount of the ␣2 phase was a little smaller compared with that in sample B. This could be due to the higher annealing temperature after the rolling. Fig. 4 illustrates the optical micrographs of the Ti–25Al–14Nb–2Mo–1Fe alloy after the different working conditions. In sample A, processed in the B2 single phase region, large equiaxed grains with grain boundary ␣2 phase particles were observed. And there is no significant difference in the microstructures between those observed on the RD plane (Fig. 4(a)) and the TD plane (Fig. 4(b)). Meanwhile, in samples B and C, processed in the (B2 + ␣2 ) two-phase region, the so-called VGS structure was seen. The RD plane exhibited wavy and curled bands of ␣2 phase particles as seen in Fig. 4(c). These ␣2 phase bands roughly align along the rolling direction, as observed in the TD plane (Fig. 4(d)). There were some differences in the distributions of ␣2 phase particles between samples B and C. In sample C, with a higher forging temperature of 1523 K and a higher annealing temperature of 1293 K, the widths of the ␣2 phase bands were somewhat larger and the spacing of the bands was also larger. Fig. 5 illustrates magnified SEM images of the Ti–25Al–14Nb–2Mo–1Fe alloy. All images were taken on the RD planes. In sample A (Fig. 5(a)), the ␣2 phase can be seen with the dark contrast along the grain boundaries. Inside the grains, fine acicular O phase (the gray contrast region) precipitated in the B2 phase matrix (the light contrast region). In sample B (Fig. 5(b)), curled VGS bands, which consisted of spherical ␣2 precipitates with a diameter of between 2 and 5 m, were observed. Between the bands, there was a fine (O + B2) two-phase microstructure similar to the one in sample A. Although the microstructural features were similar in sample C (Fig. 5(c)), the dispersed ␣2 precipitates were finer and more angular in shape. By processing the low-magnification SEM images using the image analysis software, the volume fraction of dispersed ␣2 precipitates was determined to be 20.6 ± 1.0% and 18.5 ± 0.7% in samples B and C, respectively. The ␣2 phase fraction was smaller in sample C, which is in agreement with the X-ray diffraction results. Fig. 6 shows the backscattered electron images (Fig. 6(a) and (b)) and elemental distribution maps for Nb (Fig. 6(c) and (d)) and Mo (Fig. 6(e) and (f)) obtained by EPMA for samples B and C. The VGS bands of the spherical ␣2 phase particles seen in Fig. 6(a) and (b) coincided with Nb and Mo lean regions in Fig. 6(c)–(f). Naka et al. reported [13] that plastic instabilities arise from a region-toregion variation in hardness during the high-temperature extrusion of Nb-containing B2 phase-based alloys. He insisted that this variation originated from the dendritic segregation in the starting ingot, and the metallographic contrasts of the VGS structure were related to the local variation in chemical composition. We also have confirmed the segregation of alloying elements (especially Nb and Mo) with the spacing of around 100 m in as consolidated ingot using energy dispersive X-ray spectroscopy. In the present study, the Nb-containing alloy was processed by hot bar rolling. Since the deformation state of hot bar rolling is similar to that of the hot extrusion, VGS structure related to the local segregation can be formed during the hot bar rolling. So it is supposed that the dendritic segregation in the ingot, which is coincident with the region-to-region variation in hardness, has evolved into wavy and curled bands of Nb and Mo lean regions due to the inhomoge-
Fig. 5. SEM images of Ti–25Al–14Nb–2Mo–1Fe. (a) Sample A, (b) sample B, and (c) sample C. The observed plane is normal to the rolling direction.
neous deformation during the hot bar rolling. The bar rolling was performed at 1273 K in the (B2 + ␣2 ) two-phase region, and thus precipitation of the ␣2 phase particles must be occurred simultaneously. Nb and Mo are both the B2 phase stabilizing elements, so the ␣2 phase preferably precipitated at the Nb and Mo lean regions. Finally, during the annealing in the (B2 + ␣2 ) two-phase region (at 1273 K for sample B and at 1293 K for sample C, respectively), the precipitated ␣2 phase particles grew spherically and VGS-like ␣2 phase bands were constructed. We have applied the hot bar rolling in the (B2 + ␣2 ) two-phase region to the Ti–22Al–27Nb alloy processed by powder metallurgy [11]. In this case, we obtained a homogeneous distribution of the ␣2 phase instead of the VGSlike structure. We have used the gas-atomized (i.e. rapidly cooled) Ti–22Al–27Nb alloy powder with a diameter of less than 150 m as the starting powder. Even if the elemental segregation existed in the gas-atomized powder, the spacing of the segregation must be very fine. Furthermore, these powders were consolidated using hot isostatic pressing at high temperature (1373 K) for long period (7.2 ks). Therefore, elemental segregation in this powder metal-
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Fig. 6. EPMA mappings for Nb and Mo distribution in sample B ((a), (c), and (e)) and sample C ((b), (d), and (f)) of Ti–25Al–14Nb–2Mo–1Fe. (a) and (b) backscattered electron images. (c, d) and (e, f) are the Nb and Mo maps, respectively. Concentrations are expressed by mass%. The observed plane is normal to the rolling direction.
lurgy processed material was significantly small, and VGS-like structure seems to be unfavorable. The reason why sample C exhibited wider VGS bands and wider spacing could also be explained by the segregation of Nb and Mo. Sample C was forged at a much higher temperature (1523 K) than sample B before hot bar rolling. At this temperature, the diffusion
rates of Nb and Mo are much higher. Using the diffusion coefficients of these atoms in -Ti [16], the diffusion rates of Nb and Mo are estimated to be around four times higher than those at 1273 K. So, after the forging, the degree of segregation in sample C, which was forged at 1523 K, was supposed to become lower compared with that in sample B, which was forged at 1273 K. This means that the
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Fig. 7. Room temperature tensile stress–strain curves of Ti–25Al–14Nb–2Mo–1Fe.
local variation of hardness in sample C before bar rolling became smaller than that in sample B. Therefore, it is supposed that the VGS-like segregation in sample C became less significant and the width and spacing of the VGS bands became wider after the hot bar rolling. Room temperature tensile stress–strain curves of the Ti–25Al–14Nb–2Mo–1Fe alloy are shown in Fig. 7. In Fig. 7, tensile strain was calculated from the cross-head displacement. The values of 0.2% proof stress (YS), tensile strength (TS) and total elongation (EL) are summarized in Table 1. All the samples were broken without necking. Sample A (without VGS structure) exhibited the lowest EL value of all the microstructures we have evaluated [10]. On the contrary, samples B and C (with VGS structure) exhibited higher EL values. Sample B exhibited the largest EL value of 6.3% along with the highest YS (705 MPa) and TS (992 MPa). This EL value is comparable to that of the Ti–22Al–27Nb alloy solution treated at 1273 K followed by aging at 1033 K [4]. Meanwhile, the EL value of sample C was 1.7%, which is higher than that of sample A. Fig. 8(a)–(c) shows the SEM images of the tensile fracture surfaces of samples A–C, respectively. The SEM images of the cross section areas beneath the tensile fracture surfaces are shown in Fig. 9. Sample A (without VGS structure) exhibited a coarse brittle fracture surface. Large cracks occurred mainly inside the grains. These large cracks resulted in the poor room temperature El value of sample A. On the other hand, sample B exhibited a much finer and dimple-like fracture surface. Beneath the fracture surface, a large number of small subcracks were observed. These subcracks were often arrested and/or deflected at the ␣2 phase bands, as shown in Fig. 9(d). These crack arrest and crack deflection seem to have prevented premature fracture and enhance the EL values. Sample C exhibited another type of fracture features. The fracture surface was a little coarser compared to that of sample B, and deflected and/or arrested subcracks were not observed as often as in sample B. Coarser fracture surface and smaller number of the small sub-
Table 1 Room temperature tensile properties of the Ti–25Al–14Nb–2Mo–1Fe alloy with different processing conditions. YS, TS and EL refer to 0.2% proof stress, tensile strength and total elongation, respectively.
Sample A Sample B Sample C
YS (MPa)
TS (MPa)
EL (%)
– 705 603
781 992 765
0.1 6.3 1.7
Fig. 8. SEM images of the fracture surfaces of Ti–25Al–14Nb–2Mo–1Fe after tensile tests. (a) Sample A, (b) sample B, and (c) sample C.
cracks suggest that the crack propagation resistance of sample C was lower than that of sample B. Therefore the EL value of sample C seems to be lower than that of sample B. Even though some long cracks parallel to the tensile direction were also observed in sample C, these cracks do not seem to be directly related to the lower EL values of sample C. Rather, these long parallel cracks indicate the mechanical anisotropy of sample C. Fig. 10 shows the creep curves of samples A–C of the Ti–25Al–14Nb–2Mo–1Fe alloy tested at 923 K under the initial applied stress of 310 MPa. As a reference, the result of Ti–22Al–27Nb alloy with similar microstructure to the sample A [10] is also shown in Fig. 10. Contrary to the room temperature EL values, the samples with the VGS structure exhibited lower creep resistance. This reduction in the creep resistance seems to be due to the fine VGS microstructure. However all the specimens of Ti–25Al–14Nb–2Mo–1Fe alloy exhibited a higher creep resistance compared with that of the Ti–22Al–27Nb alloy.
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Fig. 9. SEM images of transverse sections of the tensile tested specimens of Ti–25Al–14Nb–2Mo–1Fe. Tensile axis is parallel to the horizontal direction. (a) and (b) sample A, (c) and (d) sample B, (e) and (f) sample C.
4. Conclusions To improve the room temperature elongation-to-failure of the Ti–25Al–14Nb–2Mo–1Fe (mol%) orthorhombic phase base alloys, hot bar rolling and a subsequent annealing were performed in the (B2 + ␣2 ) two-phase region.
Fig. 10. Creep curves of Ti–25Al–14Nb–2Mo–1Fe with and without VGS structure. Samples were tested at 923 K under the initial applied stress of 310 MPa. The result of Ti–22Al–27Nb [10] is also shown for comparison.
(1) After the thermomechanical treatment in the (B2 + ␣2 ) twophase region, a ‘Van Gogh’s Sky (VGS)’ structure, with wavy and curled bands of spherical ␣2 precipitates, was obtained. (2) The VGS bands corresponded to the Nb and Mo lean regions. The formation of this microstructure may be related to the dendritic segregation of Nb and Mo during the solidification. This segregation has evolved into wavy and curled bands of Nb and Mo lean regions due to the inhomogeneous deformation during the hot bar rolling. ␣2 precipitation in the Nb and Mo lean regions also occurred during forging and hot bar rolling.
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(3) The sample forged at the higher temperature (1523 K) before bar rolling and annealed at the higher temperature (1293 K) after bar rolling exhibited wider VGS bands and wider spacing. (4) The samples with the VGS structure exhibited higher tensile elongation-to-failure at room temperature. The creep resistance at 923 K decreased with the existence of the VGS structure. These mechanical properties were affected by the width and spacing of the VGS bands. References [1] D. Banerjee, A.K. Gogia, T.K. Nandy, V.A. Joshi, Acta Metall. 36 (1988) 871– 882. [2] K. Muraleedharan, T.K. Nandy, D. Banerjee, S. Lele, Intermetallics 3 (1995) 187–199. [3] R.G. Rowe, in: Y.-W. Kim, R.R. Boyer (Eds.), Microstructure/Property Relationships in Titanium Aluminides and Alloys, TMS, Warrendale, 1991, pp. 387– 398.
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