Heterogeneous twin formation and its effect on tensile properties in Ti–Mo based β titanium alloys

Materials Science and Engineering A 554 (2012) 53–60

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Heterogeneous twin formation and its effect on tensile properties in Ti–Mo based ␤ titanium alloys X.H. Min a,∗ , K. Tsuzaki a,b , S. Emura a , K. 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 21 March 2012 Received in revised form 31 May 2012 Accepted 5 June 2012 Available online 17 June 2012 Keywords: ␤ titanium alloys Phase stability Twinning Ductility Heterogeneous elemental distribution

a b s t r a c t The deformation modes and the tensile properties were investigated in a Ti–15Mo–5Zr alloy with different heat treatments. Both the {3 3 2}1 1 3 twinning and the dislocation slip occurred in the samples with the heterogeneous distribution of Mo and Zr atoms. On the other hand, the twinning disappeared when the elemental distribution became homogeneous. The high yield strength and significant uniform elongation resulted from the combination of the twinning and the slip, and the more significant work hardening rate was obtained by the enhancement of the twinning. The heterogeneous twin formation was mainly discussed based on the effects of the interval of heterogeneous elemental distribution and the grain orientation. © 2012 Elsevier B.V. All rights reserved.

1. Introduction ␤ titanium alloys are known for their use as structural materials due to their high specific strength and good corrosion resistance. However, their relatively poor formability [1,2] is still a problem that affects their applications from an engineering viewpoint. An increase in the strength–ductility relationship is desirable for widening the range of their applications especially under complex loading conditions. For example, the recently developed transformation induced plasticity (TRIP) and twinning induced plasticity (TWIP) steels possess an excellent combination of strength and ductility [3], which is favorable for making complex geometrical components. However, a simultaneous improvement in both the strength, especially in yield strength, and the ductility seems to be difficult in ␤ titanium alloys. A series of special thermomechanical processing control is known to enhance the yield strength of ␤ titanium alloys by the ␣ phase precipitation [4–6]; however, the ductility, especially the uniform elongation, is limited. So far an increase in the strength–ductility relationship has not received much attention on the basis of the review of literatures for ␤ titanium alloys. Only a few

∗ Corresponding author at: National Institute for Materials Science, Research Center for Strategic Materials, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan. Tel.: +81 29 859 2529; fax: +81 29 859 2101. E-mail addresses: [email protected], [email protected] (X.H. Min). 0921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.06.009

studies [7,8] have reported the possibility of a ductility improvement through a plasticity induced transformation in titanium alloys (PiTTi), which is similar to the mechanism of the above-mentioned TRIP steels. On the other hand, a very useful method for obtaining good ductility in ␤ titanium alloys is by {3 3 2}1 1 3 twinning [9,10]; however, there remains a key issue on how to improve their yield strength. In the previous study, the present authors [11] found that a combination of different deformation modes, namely dislocation slip and {3 3 2}1 1 3 twinning was effective for achieving high yield strength and good ductility in two ␤ titanium alloys, i.e. Ti–15Mo–5Zr and Ti–10Mo–2Fe alloys. The high yield strength was mainly caused by the slip, while the large uniform elongation was generated by the twinning through significant work hardening. Moreover, the present authors suggested that the presence of various deformation modes as mentioned-above is caused not only by the grain orientation but also by the segregation of alloying elements, i.e. the heterogeneous elemental distribution such as Mo, Zr and Fe. The generally used transition metals such as Mo, Nb, V, Zr, Ta, Cr and Fe are known to segregate easily in ␤ titanium alloys, and much effort has been made to reduce or eliminate the segregation of the alloys [12–15]. However, due to the existence of the heterogeneous elemental distribution, for example, in the Ti–15Mo–5Zr alloy, the twinning occurs easily in the low ␤ phase stability region containing low Mo and high Zr contents, while the slip occurs in the high ␤ phase stability region containing high Mo and low Zr contents. This result leads to a heterogeneous deformation in the grains where the

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number fraction of the grains with twins is saturated to about 56% after being deformed to a tensile strain of 4.0%. Here, at least two questions still remain. One is whether or not the twinning can be enhanced further in this alloy by adjusting the heterogeneous elemental distribution; the other is how to change in the deformation mode if the elemental distribution becomes homogeneous. The aim of this study is to investigate the effects of the elemental distribution on the deformation modes further, namely the dislocation slip and the {3 3 2}1 1 3 twinning, and on the tensile properties in a Ti–15Mo–5Zr alloy, and to discuss the possibility of improvement in the tensile properties by controlling the deformation microstructure. 2. Experimental 2.1. Samples preparation A Ti–15Mo–5Zr alloy was prepared by cold crucible levitation melting. Fig. 1 shows a schematic drawing of the heat treatment for this alloy [11]. An ingot of 80 mm in diameter and 30 mm in length was treated at 1273 K for 3.6 ks, hot forged at 1273 K into a block of 60 mm (l) × 60 mm (w) × 40 mm (t), and then hot rolled into a plate of 240 mm (l) × 60 mm (w) × 10 mm (t) at 1173 K followed by air cooling. The plate was cut into 40 mm (l) × 60 mm (w) × 10 mm (t) pieces, and then the pieces were solution treated (ST) at 1173 K for 3.6 ks followed by water quenching. The heat treatments were carried out in the air. Note that the ␤ transus temperature for this alloy is around 1000 K based on the relationship between the ␤ transus temperature and the chemical compositions in the titanium alloys reported by Ouchi [16]. The principal axes of the rolled plate corresponding to the normal direction (ND), rolling direction (RD) and transverse direction (TD) are defined as shown in Fig. 1. The analyzed chemical composition of the alloy after the solution treatment was 15.7Mo–5.4Zr–0.103O–0.004N–0.008C (mass%) with a balance of Ti. Three types of samples were used in this study. The ST sample refers to the above-mentioned solution treated sample. The CR–ST sample was produced by subjecting the ST sample to further cold rolling (CR) with a thickness reduction ratio of 70% and then carrying out a solution treatment for changing the heterogeneous elemental distribution. On the other hand, the H–ST sample was made by homogenizing (H) the ST sample at a higher temperature of 1473 K for 14.4 ks in an argon atmosphere, hot rolling it to a thickness of 3 mm at 1173 K and then conducting a solution treatment. This heat treated sample is used to achieve a homogeneous elemental distribution.

CCLM

1273K Forging Rolling 3.6ks 1173K

Temperature

1173K 3.6ks Tβ ≈ 1000K

AC

Time

60

80

ND

60

40 30

WQ

10 60

240

60 10 40

Fig. 2. Optical micrographs of the samples. (a) ST sample, (b) CR–ST sample and (c) H–ST sample. The observed plane is normal to the TD and the horizontal direction is parallel to the RD.

2.2. Microstructural characterization The specimens were mechanically polished and etched using a solution of distilled water, nitric acid and hydrofluoric acid (100:3:2 in volume) and then observed with an optical microscope (OM). The microstructures of the mechanically polished specimens without being etched were also observed with a field emission scanning electron microscope (FE-SEM), and the elemental distributions were investigated quantitatively by the energy dispersive spectroscopy (EDS) line–spot analysis with a step size of 1 ␮m performed on a JSM-7001F system. The same system with orientation analyzing accessories was employed to conduct the electron backscattered diffraction (EBSD) analysis with a step size of 1 m. The phase identification was made by the X-ray diffraction (XRD) analysis with a RINT-2500 diffractometer using a Cu–K␣ radiation operated at 40 kV and 300 mA. The Vickers hardness measurements were carried out with a load of 49 N at 10 positions for each specimen.

2.3. Tensile test TD RD

Fig. 1. A schematic drawing of the heat treatment to obtain the solution treated (ST) sample from the Ti–15Mo–5Zr ingot. CCLM, AC and WQ refer to cold crucible levitation melting, air cooling and water quenching, respectively.

Tensile specimens with a width of 4 mm and a thickness of 2 mm in the gage section with a length of 18 mm were cut from the ST, CR–ST and H–ST samples by electric discharge machining. The tensile tests were carried out at ambient temperature with an initial strain rate of 2.78 × 10−4 s−1 .

X.H. Min et al. / Materials Science and Engineering A 554 (2012) 53–60

3. Results 3.1. Microstructure and elemental distribution Fig. 2 shows the optical micrographs of the ST, CR–ST and H–ST samples. The OM observations reveal a series of ‘bands’, parallel to the rolling direction (the horizontal direction in the figure), in the ST and CR–ST samples as shown in Fig. 2(a) and (b), respectively, while no bands appear in the H–ST sample in Fig. 2(c). The density of the bands seems to be slightly higher in the CR–ST sample than in the ST sample. In addition, the observations show the presence of the ␤ phase in all the samples with no significant difference in the grain size among the samples. XRD profiles indicate that the diffraction peaks for the athermal ␻ phase are very weak in three samples, and significant difference was not observed between them. The Vickers hardness, which is about 278, 275 and 284 Hv in the ST, CR–ST and H–ST samples, respectively, also implies the similar presence of the athermal ␻ phase in three samples even with the different heat treatments in this study. Fig. 3 shows the backscattered electron images and the distributions of Mo and Zr in the three samples. The images in Fig. 3(a) and

55

(c) show a series of ‘bands’ in the ST and CR–ST samples, respectively, while there are no bands in the H–ST sample in Fig. 3(e). The band in the CR–ST sample has a higher density and is thinner than that of the ST sample. Note that the effect of roughness of the specimens on the bands is very small since the specimens are not subjected to etching for the SEM observations. In addition to the effect of the grain orientation on the contrast of the images, the contrast of the bands reflects the different content of elements. Fig. 3(b), (d) and (f) shows the distributions of Mo and Zr in the samples along the analyzed lines. The bright bands correspond to a high content of Mo and a low content of Zr, and the dark ones to a low content of Mo and a high content of Zr in the ST and CR–ST samples. On the other hand, the distributions of Mo and Zr are almost the same along the analyzed line in the H–ST sample. As shown in Table 1, the content of Mo and Zr varies from 12.06 to 17.00 and 3.84 to 6.19 mass%, respectively, in the ST sample. The content of Mo and Zr in the CR–ST sample varies from 12.84 to 16.53 and 4.30 to 5.84 mass%, respectively, and has a narrower range than that in the ST sample. The H–ST sample has the narrowest range of content of Mo and Zr, respectively, from 13.24 to 15.24 and 4.50 to 5.75 mass%. Table 1 shows that the average contents of Mo and Zr in the analyzed local areas are 14.13 and

Fig. 3. Backscattered electron images (a, c and e), and elemental distribution of Mo and Zr along the white lines (b, d, and f) in the samples. (a and b) ST sample, (c and d) CT–ST sample and (e and f) H–ST sample. The analyzed plane is normal to the TD and the vertical direction is parallel to the RD.

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Table 1 Analyzed contents of Mo and Zr by the energy dispersive spectroscopy line analysis, and calculated Mo equivalency based on the contents of Mo and Zr in the ST, CR–ST and H–ST samples (mass%). Samples

ST CR–ST H–ST

Mo content

Zr content

Mo equivalency

Min

Max

Ave

Min

Max

Ave

Min

Max

Ave

12.06 12.84 13.24

17.00 16.53 15.24

14.13 14.59 14.31

3.84 4.30 4.50

6.19 5.84 5.75

5.25 5.04 5.14

14.97 15.38 15.56

19.08 18.78 17.51

16.60 16.96 16.72

5.25 mass% in the ST sample, 14.59 and 5.04 mass% in the CR–ST sample, and 14.31 and 5.14 mass% in the H–ST sample, respectively. Thus, the microstructure observations and the EDS analyses indicate that the distributions of Mo and Zr in the ST and CR–ST samples are clearly heterogeneous, while they are almost homogeneous in the H–ST sample. In addition to the varied amplitude, the varied interval (width of the bands) of Mo and Zr in the CR–ST sample is also smaller than that in the ST sample as shown in Fig. 3(b) and (d). 3.2. Tensile property Fig. 4(a) shows the nominal stress–strain curves of the ST, CR–ST and H–ST samples along with the Ti–15Mo alloy. Note that the Ti–15Mo alloy was subjected to the same heat treatment conditions as that of the ST sample. Their true stress–strain curves and the work hardening rates are shown in Fig. 4(b). Table 2 provides the tensile properties of 0.2% proof stress, i.e. yield strength (YS), tensile strength (TS), uniform elongation (uEL), total elongation (tEL) and the work hardening rate (WHR) at a 4.0% tensile strain. As

1000

(a)

H-ST

Stress (MPa)

800

CR-ST

Ti-15Mo

600

ST 400 200 0 0

10

20

30

40

50

60

True stress, work hardening rate (MPa)

Strain (%) 2500

(b) 2000

Ti-15Mo

1500

ST

CR-ST

1000 500

H-ST

0 0

0.05

0.1

0.15

0.2

0.25

0.3

0.35

True strain Fig. 4. Nominal stress–strain curves (a), and true stress–strain curves and corresponding work hardening rates (d/dε), where  and ε are true stress and true strain respectively (b), in the ST, CR–ST and H–ST samples. The results of the Ti–15Mo alloy [11] are also shown for comparison.

mentioned previously [11], the Ti–15Mo alloy shows the low yield strength of 439 MPa and large uniform elongation of 27% through significant work hardening, and the ST sample has a combination of the high yield strength of 760 MPa and significant uniform elongation of 10%. In this study, the CR–ST sample shows the high yield strength of 734 MPa and significant uniform elongation of 13%. Note that the work hardening rate of 1825 MPa at a 4.0% tensile strain in the CR–ST sample (Table 2) is more significant than that of 1120 MPa in the ST sample in Fig. 4(b). On the other hand, the yield strength of 956 MPa in the H–ST sample is the highest among the samples, while the uniform elongation is almost 0% even with a total elongation of 12%. The calculated work hardening rate shown in Fig. 4(b) indicates the initiation of necking soon after the yielding in the H–ST sample.

3.3. Deformation microstructure Fig. 5 shows the optical micrographs of the ST, CR–ST and H–ST samples after being deformed to fracture, and of the CR–ST sample after being deformed to a tensile strain of 4.0%. In Fig. 5(a), the plate-like features observed in some of the grains in the fractured ST sample are identified as the {3 3 2}1 1 3 twins as described in the previous study [11]. The fractured CR–ST sample basically has the similar deformation microstructure as shown in Fig. 5(b). In addition, the twins are heavily deformed, and some of the twins become lenticular in shape. The density of the twins is higher and the number fraction of the grains with the twins is larger than in the fractured ST sample. On the other hand, the twins are not observed in the grains in the fractured H–ST sample as shown in Fig. 5(c); however, the slip traces are clearly noticeable at a higher magnification in Fig. 5(d). Fig. 5(e) shows the deformation microstructure of the 4.0% strained CR–ST sample. The twins are thinner and its density is lower than that of the fractured sample (Fig. 5(b)), while the number fraction of the grains with twins does not seem to change much. In addition, the slip traces are clearly observed in the grains without the twins which are indicated by the white arrows at a higher magnification as shown in Fig. 5(f). A similar slip trace was confirmed by the transmission electron microscope observation in a Ti–15Mo–1Fe alloy in the previous study [17]. The EBSD measurement for the 4.0% strained CR–ST sample was used to further confirm the {3 3 2}1 1 3 twins through the inverse pole figure map in Fig. 6(a), and their boundaries delineated by the red lines in Fig. 6(b). Based on the optical observation and the EBSD measurement, the twin formation of the CR–ST sample is heterogeneous among the grains and is similar to that of the ST sample. To discuss the difference in the twin formation between the two samples quantitatively, the number fraction of the grains containing the twins was obtained from the optical observations of 2000 grains in the 4.0% strained and fractured CR–ST sample. Fig. 7 shows the change in the number fraction of the grains with the twins against the strain in the three samples along with the Ti–15Mo alloy. The number fraction is about 73% (4.0% strain) and 75% (fracture) in the CR–ST sample, which is larger than 55% (4.0% strain) and 57% (fracture) in the ST sample [11]; however, the number fraction is 0% in the H–ST sample, and 100% in the Ti–15Mo alloy. Here, it can be concluded that the CR–ST sample is deformed by the dislocation slip and the {3 3 2}1 1 3 twinning results in high strength and significant uniform elongation. The higher work hardening rate compared to that of the ST sample is due to the enhancement of the twinning. On the other hand, the H–ST sample is deformed only by the slip, and the balance of its tensile properties shows deterioration.

X.H. Min et al. / Materials Science and Engineering A 554 (2012) 53–60

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Table 2 Tensile properties for 0.2% proof stress (YS), tensile strength (TS), uniform elongation (uEL), total elongation (tEL), work hardening rate (WHR) at 4.0% strain and the number fraction of grains with twins (twin fraction) in the ST, CR–ST and H–ST samples along with the Ti–15Mo alloy. Samples

ST CR–ST H–ST Ti–15Mo

YS (MPa)

760 734 956 439

TS (MPa)

801 820 956 714

uEL (%)

10 13 0.2 27

4. Discussion This study further confirmed that a combination of both the dislocation slip and the {3 3 2}1 1 3 twinning is effective for achieving the high yield strength and large uniform elongation in the CR–ST sample. Moreover, a better strength–ductility relationship is obtained through the enhancement of the twinning in the CR–ST sample. This study also showed an interesting result of the absence of the twin formation in the H–ST sample. The next section discusses the changes in the twin formation among the samples. The present authors [11] reported that the various deformation modes in the ST sample cannot be explained solely by the

tEL (%)

19 23 12 49

WHR (MPa)

1120 1825 360 1839

Twin fraction (%)

Ref.

4% strain

Fracture

55.2 73.1 0 100

56.8 74.5 0 100

[11]

[11]

effect of the grain orientation, and emphasized the effect of the heterogeneous elemental distribution of Mo and Zr. The SEM–EDS mappings for the 4.0% strained ST sample indicated that the twinning occurred in the region with low Mo and high Zr content corresponding to the low ␤ phase stability, while the slip occurred in the region with high Mo and low Zr content corresponding to the high ␤ phase stability [11]. Here, it is needed to consider why the twin formation was enhanced in the CR–ST sample and absent in the H–ST sample as shown in Figs. 5 and 7. The factors affecting the twinning such as the testing condition [18], grain size [19], grain orientation [20], phase stability [11,21] should all be considered. All the tensile

Fig. 5. Optical micrographs of the samples after being deformed to fracture in (a) the ST sample, (b) CR–ST sample, (c and d) H–ST sample, and to a tensile strain of 4.0% in (e and f) CR–ST sample. The observed plane is normal to the ND and the tensile axis parallel to the RD is the horizontal direction.

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Number fraction of grains with twins (%)

Fig. 6. EBSD maps of the CR–ST sample after being deformed to a tensile strain of 4.0%. (a) RD inverse pole figure map with the color code presented in the upper right corner and (b) boundaries of the {3 3 2}1 1 3 twins delineated by the red lines and of the grains by the gray lines. The observed plane is normal to the ND and the tensile axis parallel to the RD is the horizontal direction. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of the article.)

Fig. 8. Orientations of the grains in the standard stereographic triangle of RD. (a) ST sample, (b) CR–ST sample and (c) H–ST sample. The dotted lines indicate the contours with equal Schmid factor for the (23 3)[3¯ 1 1] twinning system.

100

Ti-15Mo 80

CR-ST

60

ST

40 20

H-ST

0 0

10

20

30

40

50

Strain (%) Fig. 7. Change in the number fraction of grains with {3 3 2}1 1 3 twins against tensile strain in the ST, CR–ST and H–ST samples. The result of the Ti–15Mo alloy [11] is also shown for comparison.

tests were carried out at ambient temperature and the grain size with an average value of 100 ␮m was almost the same between the samples based on the OM observations (Fig. 2). Hence, the effects of the first two factors are considered to be negligible. Fig. 8 shows the orientations of the grains in the samples within an area of 800 ␮m × 800 ␮m in the standard stereographic triangle of the RD. The dotted lines in the triangle indicate the contours with equal Schmid factor for the (2 3 3)[3¯ 1 1] twinning system, which shows an easily operative twinning system among 12 possible {3 3 2}1 1 3 twinning systems during the tensile deformation [22]. Fig. 8(a) and (c) indicates that the ST and H–ST samples show almost uniform distribution of the grain orientations. However, the CR–ST sample shows that few orientations appear near the 0 0 1 as shown in Fig. 8(b).

Mo equivalency (mass%)

20

Mo equivalency (mass%)

The deformation twinning is known to depend strongly on the grain orientation [23]. For example, many investigators [24,25] studied the orientation dependence of the {1 1 1}1 1 2 twinning and the dislocation slip on the TWIP steels, and they reported that the orientation close to the 1¯ 1 1 is favorable for the twinning, while it near the 0 0 1 is easily deformed by the multiple slip. This selection for the deformation modes can be basically interpreted by the difference in the Schmid factor between partial Shockley dislocation and perfect dislocation. In addition, the crystallographic texture affecting the deformation twining based on the Schmid factor criterion in the extruded magnesium alloy tubes was reported by Godet et al. [26]. The similar orientation dependence is also present for the {3 3 2}1 1 3 twinning and the dislocation slip in the ␤ titanium alloys. Hanada et al. [22] reported that the orientation around the 1¯ 1 1 is favorable for the twinning, while it in the neighborhood of the 0 0 1 to the 0 1 1 is favorable for the slip, divided into two regions in the standard stereographic triangle based on the different Schmid factor between two deformation modes. Therefore, the presence of initial texture in the CR–ST sample before tensile deformation is thought to be one of the reasons for the enhancement in the twin formation. Note that the varied amplitude and interval of Mo and Zr in the CR–ST sample is different from that of the ST as shown in Fig. 3, indicating that the varied region of the ␤ phase stability differs between the samples. In this study, the Mo equivalency is used to evaluate the ␤ phase stability, which is the ratio of the level of a given stabilizer to the one of Mo required for the same degree of ␤ phase stability [27–29]. One mass% Zr is equivalent to 0.47 mass% Mo. In addition, the Mo equivalency is also useful for evaluating the deformation modes [30–32]. Fig. 9 shows the change in the Mo equivalency of the samples based on the contents of Mo and Zr (Fig. 3). The Mo equivalency varies from 14.97 to 19.08 mass%, 15.38 to 18.78 mass%, and 15.56 to 17.51 mass% in the ST, CR–ST and H–ST samples, respectively (Table 1). This result indicates that the varied amplitude of Mo equivalency is slightly smaller in the CR–ST sample than in the ST sample, while the average values are almost the same as that in the H–ST sample (Table 1). In addition, the interval of the Mo equivalency is about 100 ␮m in the ST sample, and about 50 ␮m in the CR–ST sample. Here, an interval is defined as the distance where the Mo equivalency is higher or lower than the average value. These results suggest that the higher number fraction of the grains with the twins in the CR–ST sample can also be caused by the narrower interval of the Mo equivalency. Fig. 10 shows the correlation between the change in the Mo equivalency and the twin formation in the grains of the ST and CR–ST samples. The grain size is the same between the samples, and the Mo equivalency

20

Mo equivalency (mass%)

X.H. Min et al. / Materials Science and Engineering A 554 (2012) 53–60

20

59

(a) ST

19 18 17 16 15 14

(b) CR-ST

19 18 17 16 15 14

(c) H-ST

19 18 17 16 15 14 0

100

200

300

400

500

Distance ( μm) Fig. 9. Change in the Mo equivalency with distance in (a) the ST sample, (b) the CR–ST sample and (c) the H–ST sample. The Mo equivalency is calculated by using the analyzed contents of Mo and Zr as shown in Fig. 3.

is assumed to show a sinusoidal variation. The Mo equivalency of the H–ST sample is used as the average amplitude, which is indicated by the straight line in the figure. The existence of only the slip deformation in the H–ST sample indicates that the critical Mo equivalency between the twinning and the slip is lower than this value. As shown in Fig. 10(a), when the amplitude of the Mo equivalency is lower than the average value, the grain is deformed easily by the twinning, and when it is equal to or higher, the grain is deformed by the slip. On the other hand, when the interval of the Mo equivalency is decreased from a sinusoidal dotted line to a sinusoidal solid line as shown in Fig. 10(b), the grain contains

Fig. 10. Correlation between the Mo equivalency and the twin formation in the grains for (a) the ST sample, and change in the twin formation with reducing the interval of Mo equivalency from the dotted line to the solid line in the grains for (b) the CR–ST sample.

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both the low and high Mo equivalency to result in the presence of the twins. Here, a critical size of the region with low Mo equivalency is thought to exist for the nucleation of the twins, and a reduced interval of the Mo equivalency is intended to increase the nucleation site for the twinning. Therefore, mostly grains containing the region of low ␤ stability are deformed by the twinning in the CR–ST sample since the ‘bands’ become denser and thinner (Figs. 2 and 3). Hanada and Izumi [10,20] reported the different deformation modes in various ␤ titanium alloys, and described that the occurrence of the {3 3 2}1 1 3 twinning was closely related to the stability of the ␤ phase with respect to the athermal ␻ phase. In the macroscopic scale, the XRD profiles and the Vickers hardness showed that the presence of the athermal ␻ phase has no significant difference in three samples as mentioned in Section 3.1. However, the formation of the athermal ␻ phase is thought to be affected by the heterogeneous elemental distribution, especially in the CR–ST sample. The previous studies [17,33] reported that the spots of the athermal ␻ phase from the selected area electron diffraction patters became weak and diffuse with increasing the Mo equivalency in either the Ti–15Mo based alloys or the Ti–10Mo based alloys, and the deformation mode changed from the {3 3 2}1 1 3 twinning to the dislocation slip. In the present study, it is inferred that the regions with the low Mo equivalency in the CR–ST sample are favorable for the presence of the athermal ␻ phase accompanying the twin formation. In addition, impurities, especially oxygen (O), should be considered in the H–ST sample since the homogenization was carried out at a high temperature of 1473 K. The content of O was analyzed to be 0.139 mass% in the H–ST sample, which is slightly higher than 0.103 mass% in the ST sample. Note that the effect of O on the ␤ transus temperature can be evaluated as reported by Ouchi [16]. For example, the ␤ transus temperature increases by 147.7 K with 1 mass% O, and it decreases by 10.3 K with 1 mass% Mo. The increase in the ␤ transus temperature is 5.3 K in the H–ST sample due to the change in the content of O. This result indicates that the effect of O on the slip deformation in the H–ST sample is negligible. It can be concluded that the heterogeneous twin formation, which is mainly affected by the grain orientation and the segregation of alloying elements, is thought to be effective for achieving a good strength–ductility relationship in the Ti–15Mo–5Zr alloy. To enhance the deformation twinning by the grain orientation as shown in Fig. 8, the orientations are desirable to be as close as possible to the 1¯ 1 1, while it is difficult as mentioned previously [11]. Here, a further optimization of the heterogeneous elemental distribution can be desirable by controlling the heat treatment conditions. 5. Conclusions The microstructures, the deformation modes and the tensile properties were investigated in the Ti–15Mo–5Zr alloy subjected to various heat treatment conditions, classified as ST, CR–ST and H–ST samples. The main results are summarized as follows:

(1) The heterogeneous elemental distribution of Mo and Zr atoms are present in the ST and CR–ST samples and absent in the H–ST sample. The interval of the distribution is narrower in the CR–ST sample than in the ST sample. (2) The ST and CR–ST samples are deformed by both the {3 3 2}1 1 3 twinning and the dislocation slip to result in high strength and significant uniform elongation. The CR–ST sample has a more significant work hardening rate than that of the ST sample due to the enhancement of the twinning that is mainly caused by the narrower interval of the distribution and the grain orientation. On the other hand, the H–ST sample is deformed by the slip, which leads to high strength and negligible uniform elongation. (3) The heterogeneous twin formation is effective for achieving a good strength–ductility relationship. References [1] [2] [3] [4] [5] [6] [7]

[8]

[9] [10] [11] [12] [13] [14]

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