**4. Results and discussion**

### **4.1. Change in β-phase stability with alloys' composition**

The β-phase stability increases with increasing content of the β-stabilizing elements. Shown in figure 4.a are the X-ray diffraction patterns taken from the A-alloys in the ST condition. The α" martensite phase is the predominant phase in A00 alloy beside small amount of β and ω phas‐ es. The α" martensite phase is suppressed by O addition to the A00 β-phase alloys. As shown in figure 4.a, the addition of 1 mol.% O to A00 alloy was very effective in suppressing the α" mar‐ tensite phase as observed in A01 alloy. Also, it is clear from this figure that the addition of Fe to the A01 alloy resulted in stabilizing the β-phase as observed in the A11 alloy and the α" marten‐ site phase couldn't be observed by XRD at room temperature.

It can be deducted from the XRD results that the A11 is a single β-phase alloy. So, the least stable single β-phase alloys in the A series is A11 alloy. It is concluded that the co-addition of Fe and O to these β-type alloys is very effective in suppressing the α" martensite phase and instantaneously stabilize β-phase of the alloy. An important observation from figure 4.a is that, the co-addition of only 1% O and 1% Fe could modify greatly the phase stability of alloy A00, with the α" predominant phase, to alloy A11, with the single β-phase. Therefore, the β-phase stability in these alloys increases in the order, A11> A01> A00.

**Figure 4.** XRD profiles of the alloys Z1-4 (a), and A-alloys (b) in the solution treated at 1223K, ST, after 90% cold rolling.


**Table 2.** X-ray peak intensity ratios of Z-alloys and A-alloy in the CR and ST conditions.

By the same way for Z-alloys, shown in figure 4.b are the X-ray diffraction patterns taken from the Z-alloys in the ST condition. A single β-phase was predominant in the alloys Z1, Z2, and Z3. Only in the alloy Z4, the martensite α''-phase coexisted with the β-phase, as shown in figure 4.b. Therefore, theβ/β+α'' boundary shown in figure 3 is located between the alloys Z3 and Z4 as indicated by a dotted curve. So, the alloy Z3 is the least stable single β-phase alloy which is defined as the alloy containing a least amount of the β-stabilizing ele‐ ments to get a β single phase [20,25]. According to the *<sup>B</sup>*¯ *o*-*M*¯ *d* diagram shown in figure 3, the alloy closer to the β/β+α'' boundary has the lower β-phase stability, so the β-phase stability in these alloys decreased in the order, Z1> Z2> Z3> Z4. The stability was highest in the alloy Z1 and lowest in alloy Z4.

#### **4.2. Textures developed by cold rolling**

The phases existing in the specimen and its pole figures were identified by the conventional X-ray diffraction (XRD) using a Ni-filtered Cu-Kα radiation. Electron back scattered diffrac‐ tion (EBSD) analysis was also made using a HITACHI S-3000H scanning electron micro‐ scope (SEM) equipped with a OXFORD INCA Crystal EBSD detector, operated at an

performed using the optical microscope (OM), the scanning electron microscope (SEM) and

The β-phase stability increases with increasing content of the β-stabilizing elements. Shown in figure 4.a are the X-ray diffraction patterns taken from the A-alloys in the ST condition. The α" martensite phase is the predominant phase in A00 alloy beside small amount of β and ω phas‐ es. The α" martensite phase is suppressed by O addition to the A00 β-phase alloys. As shown in figure 4.a, the addition of 1 mol.% O to A00 alloy was very effective in suppressing the α" mar‐ tensite phase as observed in A01 alloy. Also, it is clear from this figure that the addition of Fe to the A01 alloy resulted in stabilizing the β-phase as observed in the A11 alloy and the α" marten‐

It can be deducted from the XRD results that the A11 is a single β-phase alloy. So, the least stable single β-phase alloys in the A series is A11 alloy. It is concluded that the co-addition of Fe and O to these β-type alloys is very effective in suppressing the α" martensite phase and instantaneously stabilize β-phase of the alloy. An important observation from figure 4.a is that, the co-addition of only 1% O and 1% Fe could modify greatly the phase stability of alloy A00, with the α" predominant phase, to alloy A11, with the single β-phase. Therefore,

**Alloy O Fe V Cr Mo Nb Ta Zr Z1** 4 2 9 7 30 **Z2** 3 15 3 25 **Z3** 1 8 14 15 **Z4** 4 20 5

**A00** 17 6 **A01** 1 17 6 **A11** 1 1 17 6

. The microstructural characterization was

acceleration voltage of 20 kV and a tilt angle of 71o

**Table 1.** Chemical compositions of the designed A and Z-alloys, at.%.

**4.1. Change in β-phase stability with alloys' composition**

site phase couldn't be observed by XRD at room temperature.

the β-phase stability in these alloys increases in the order, A11> A01> A00.

**4. Results and discussion**

the transmission electron microscopy (TEM).

124 Recent Developments in the Study of Recrystallization

Most of the annealed β-type Ti-alloys have random equiaxed grains and their X-ray peak in‐ tensity of the {110}β plane is the highest among the different atomic planes reflections. How‐ ever, the {100}β<110> rolling texture is formed normally after high reduction ratio of the βtype Ti-alloys due to dislocation slipping and grains rotation. In this cold rolling texture, the {200}β planes aligned parallel to the rolling plane preferentially. Figure 5.a shows the X-ray diffraction patterns taken from the Z-alloys in the 60% cold rolled (60CR) condition. It is known that the cold rolling texture is formed in conventional β- type Ti alloys. As a result, the measured X-ray peak intensity ratio of the cold rolled specimen, *IRcr*, defined as *IRcr= I{200}β / I{110}β,* changed with cold rolling. Here, *I{200}β* and *I{110}β* are the X-ray peak intensities of the {200}β and {110}β reflections, respectively. The *IRcr* usually increases with the reduction ra‐ tio of cold rolling*.* In addition, as is evident from figure 5.a and table 2, when the alloys were cold rolled by 60%, *IRcr* tended to increase with increasing β-phase stability. In other words, a cold rolling texture was developed in the way that the {200} planes were aligned parallel to the rolling plane preferentially. This texture was formed more readily in the order, Z1>Z2> Z3> Z4, in agreement with the order of the β- phase stability.

the addition of 1% Fe in this A-alloy. However, the O addition seems to be very effective in suppressing the ω- and the α"-phases (in other words, stabilizing β-phase at room tempera‐ ture), as shown in figure 5.b. It is well know that Fe is very strong β-stabilizer. Therefore, the co-addition of O and Fe in A11 alloy was enough to increase the β-phase stability to a level high enough to develop strong {100}<110> rolling texture, same as in Z-alloys with higher β-

Texturing Tendency in β-Type Ti-Alloys http://dx.doi.org/10.5772/53588 127

phase stability such as Z2 and Z3 alloys, as shown in figure 5.b and table 2.

**Figure 5.** XRD profiles of Z-alloys after 60%CR, (a) and A-alloys after 90%CR, (b).

than reduction ratio in developing rolling textures in β-type Ti-alloys.

Figure 7 (a-c) shows {200}, {110} and {112} pole figures obtained from a 98% cold rolled (CR) specimen of the A01 alloy. The center of the pole figures corresponds to the direction normal to the specimen surface (ND). It seems from these pole figures that the {111}<112> rolling texture was developed predominantly compared to the other textures and random grains orientations. The {111}<112> rolling texture developed in A01 after 98%CR is moderate in strength if compared to that {100}<110> rolling texture developed in Z2 alloy after only 90%CR as evidenced from the inverse pole figures shown in Figure 8. This means that the rolling textures in β-type Ti-alloys become stronger with increasing the β-phase stability ir‐ relative to the type of texture developed in them. Also, Z2 with the relatively high β-phase stability shows the{100}<110> texture after 90%CR more remarkably when compared to the lower β-phase stability alloys undergone by the same or even severer cold rolling as report‐ ed in Ref. [18,27,28]. Therefore, it is concluded that the β-phase stability is more effective

Figure 6 (a-c) shows {200}, {110} and {112} pole figures obtained from a 90% cold rolled (CR) specimen of the alloy Z2. The center of the pole figures corresponds to the direction normal to the specimen surface (ND). The right and the top of the pole figures correspond to the rolling direction (RD) and the transverse direction (TD), respectively. It was realized from these pole figures that typically {100}<110> rolling texture with a strength of 97.5 times larg‐ er compared to the random orientation is well developed in the 90% cold rolled (CR) speci‐ men of the alloy Z2, as shown in the {100}<110> texture stereoprojection [26] in figure 6 (d). It is important to mention here that the alloy Z2 with the relatively high β-phase stability shows the{100}<110> texture after 90%CR more remarkably, as compared to the lower βphase stability alloy undergone by the same or even severer cold rolling [18,27,28].

It has been reported that the {100}<110> texture is a main rolling texture formed in the β-type Ti-alloys [17,18,26-28]. Beside this texture, the {211}<110> texture forms in a Ti-35mass %Nb-4mass%Sn alloy [18,26] (location A in figure 3) or the {111}<112> texture forms in a Ti-24Nb-3Al alloy [26] (location B in Figure 3). With a little increase in the β-phase stability, the {100}<110> texture becomes dominant as observed in a Ti-35mass%Nb-7.9mass% Sn al‐ loy [28] (location C in figure 3). As explained in Ref. [20], both Al and Sn work as the β-stabi‐ lizing elements in these β-phase alloys. In the much higher β-phase stability alloys such as Ti-22Nb-6Ta alloy (location D in figure 3), only the {100}<110> texture is developed after 99% cold rolling [27]. This was consistent with the present results that the *IRcr,* which represents {100}<110> texture, increases monotonously with the β-phase stability. Therefore, it can be concluded that the {100}<110> rolling texture is developing in β-type Ti-alloys and its strength is increasing monotonously with increasing β-phase stability. However, in Ti alloys with low β-phase stability, other rolling textures such as {211}<110> and {111}<112> textures are developing more readily than {100}<110> texture in such low β stability alloys.

In the same way for A-alloys, as is evident from figure 5.b and table 2, when the A-alloys were cold rolled by 90%, *IRcr* tended also to increase with increasing β-phase stability. In oth‐ er words, a rolling texture was developed, in the higher β-phase stability alloys, in the way that the {200} planes aligned parallel to the rolling plane preferentially, in accordance with that reported for the above Z-alloys. For example, the *IRcr* of A11 alloy is around 6 times of that of A01 alloy with relatively lower β-phase stability. However*, IRcr* (A01) / *IRcr* (A00) are around 2 times only. This can be interpreted as the addition of 1% O is less efficient to high‐ ly stabilize β-phase and therefore to develop the {100}<110> rolling texture if compared with the addition of 1% Fe in this A-alloy. However, the O addition seems to be very effective in suppressing the ω- and the α"-phases (in other words, stabilizing β-phase at room tempera‐ ture), as shown in figure 5.b. It is well know that Fe is very strong β-stabilizer. Therefore, the co-addition of O and Fe in A11 alloy was enough to increase the β-phase stability to a level high enough to develop strong {100}<110> rolling texture, same as in Z-alloys with higher βphase stability such as Z2 and Z3 alloys, as shown in figure 5.b and table 2.

diffraction patterns taken from the Z-alloys in the 60% cold rolled (60CR) condition. It is known that the cold rolling texture is formed in conventional β- type Ti alloys. As a result,

*I{200}β / I{110}β,* changed with cold rolling. Here, *I{200}β* and *I{110}β* are the X-ray peak intensities of the {200}β and {110}β reflections, respectively. The *IRcr* usually increases with the reduction ra‐ tio of cold rolling*.* In addition, as is evident from figure 5.a and table 2, when the alloys were cold rolled by 60%, *IRcr* tended to increase with increasing β-phase stability. In other words, a cold rolling texture was developed in the way that the {200} planes were aligned parallel to the rolling plane preferentially. This texture was formed more readily in the order, Z1>Z2>

Figure 6 (a-c) shows {200}, {110} and {112} pole figures obtained from a 90% cold rolled (CR) specimen of the alloy Z2. The center of the pole figures corresponds to the direction normal to the specimen surface (ND). The right and the top of the pole figures correspond to the rolling direction (RD) and the transverse direction (TD), respectively. It was realized from these pole figures that typically {100}<110> rolling texture with a strength of 97.5 times larg‐ er compared to the random orientation is well developed in the 90% cold rolled (CR) speci‐ men of the alloy Z2, as shown in the {100}<110> texture stereoprojection [26] in figure 6 (d). It is important to mention here that the alloy Z2 with the relatively high β-phase stability shows the{100}<110> texture after 90%CR more remarkably, as compared to the lower β-

phase stability alloy undergone by the same or even severer cold rolling [18,27,28].

are developing more readily than {100}<110> texture in such low β stability alloys.

In the same way for A-alloys, as is evident from figure 5.b and table 2, when the A-alloys were cold rolled by 90%, *IRcr* tended also to increase with increasing β-phase stability. In oth‐ er words, a rolling texture was developed, in the higher β-phase stability alloys, in the way that the {200} planes aligned parallel to the rolling plane preferentially, in accordance with that reported for the above Z-alloys. For example, the *IRcr* of A11 alloy is around 6 times of that of A01 alloy with relatively lower β-phase stability. However*, IRcr* (A01) / *IRcr* (A00) are around 2 times only. This can be interpreted as the addition of 1% O is less efficient to high‐ ly stabilize β-phase and therefore to develop the {100}<110> rolling texture if compared with

It has been reported that the {100}<110> texture is a main rolling texture formed in the β-type Ti-alloys [17,18,26-28]. Beside this texture, the {211}<110> texture forms in a Ti-35mass %Nb-4mass%Sn alloy [18,26] (location A in figure 3) or the {111}<112> texture forms in a Ti-24Nb-3Al alloy [26] (location B in Figure 3). With a little increase in the β-phase stability, the {100}<110> texture becomes dominant as observed in a Ti-35mass%Nb-7.9mass% Sn al‐ loy [28] (location C in figure 3). As explained in Ref. [20], both Al and Sn work as the β-stabi‐ lizing elements in these β-phase alloys. In the much higher β-phase stability alloys such as Ti-22Nb-6Ta alloy (location D in figure 3), only the {100}<110> texture is developed after 99% cold rolling [27]. This was consistent with the present results that the *IRcr,* which represents {100}<110> texture, increases monotonously with the β-phase stability. Therefore, it can be concluded that the {100}<110> rolling texture is developing in β-type Ti-alloys and its strength is increasing monotonously with increasing β-phase stability. However, in Ti alloys with low β-phase stability, other rolling textures such as {211}<110> and {111}<112> textures

defined as *IRcr=*

the measured X-ray peak intensity ratio of the cold rolled specimen, *IRcr*,

Z3> Z4, in agreement with the order of the β- phase stability.

126 Recent Developments in the Study of Recrystallization

**Figure 5.** XRD profiles of Z-alloys after 60%CR, (a) and A-alloys after 90%CR, (b).

Figure 7 (a-c) shows {200}, {110} and {112} pole figures obtained from a 98% cold rolled (CR) specimen of the A01 alloy. The center of the pole figures corresponds to the direction normal to the specimen surface (ND). It seems from these pole figures that the {111}<112> rolling texture was developed predominantly compared to the other textures and random grains orientations. The {111}<112> rolling texture developed in A01 after 98%CR is moderate in strength if compared to that {100}<110> rolling texture developed in Z2 alloy after only 90%CR as evidenced from the inverse pole figures shown in Figure 8. This means that the rolling textures in β-type Ti-alloys become stronger with increasing the β-phase stability ir‐ relative to the type of texture developed in them. Also, Z2 with the relatively high β-phase stability shows the{100}<110> texture after 90%CR more remarkably when compared to the lower β-phase stability alloys undergone by the same or even severer cold rolling as report‐ ed in Ref. [18,27,28]. Therefore, it is concluded that the β-phase stability is more effective than reduction ratio in developing rolling textures in β-type Ti-alloys.

hence will be controlled by alloying. Also, it is known that the slip/twin boundary is close to

Texturing Tendency in β-Type Ti-Alloys http://dx.doi.org/10.5772/53588 129

As crystal deforms by slip, it undergoes crystal rotations. Such rotations lead to the develop‐ ment of preferred orientations in polycrystalline alloys [32]. The main reason why the {100}<110> texture develops by cold rolling in conventional bcc alloys may be attributable to the glide of dislocations along <111> on {011}, {112}or {123}[33] and the crystal rotations in them. It is well known that C' represents the resistance for the {011}<011> shear and C<sup>44</sup> rep‐ resents the resistance for the {001}<001> shear [34] and both of them decreases with decreas‐ ing β-phase stability [24,35]. The e/a value is kept at 4.24 in the present alloys and C' decreases with decreasing β-phase stability as discussed earlier, so the elastic softening will be enhanced in the order; Z1 < Z2 < Z 3 < Z4. This is a reason why the alloys locating at the β/β+α" boundary showed a very low shear modulus along both <011> on {011} and along <111> on {011}, {112} or {123} as reported in Ref. [36]. So it is expected that as the β-phase stability decreases, secondary slipping systems such as {011}<011> may be activated by de‐ formation beside the main slipping system, i.e., <111> on {011}, {112}or {123}. As a result, a portion of the applied stress is consumed in slipping in such secondary slipping systems, which will make some disturbances in forming the main {100}<110> rolling texture. This can be a reason why the {100}<110> rolling texture was dominant in the high β-phase stability alloys. This was also seen in the steel in which the addition of a ferrite (bcc) stabilizer, Si, is

The textures developed by severe cold rolling will diminish if the specimen is reheated for long time. For example, the *IH(cr)* of a 98%CR A01 specimen was decreased when reheated to different temperatures for 3 hrs, as shown in figure 9. *IH(cr)* is defined as *IHcr=* 1/*IRcr= I{110}β / I{200}β*. Hence, it is expected that the {100}<110> rolling texture is diminishing as a result of this reheating. The decrease in *IHcr* was higher with increasing the reheating temperature till it reach a saturation temperature. This saturation temperature is most probably related to the completion of recrystallization process. Therefore, further increase in temperature will not affect much the value of *IHcr*, as shown in figure 9. The rolling textures will not only dimin‐ ish the heavily cold rolled specimens by reheating, but also recrystallization textures may

In the present designed alloys, the recrystallization textures were developed in different strength in the ST specimen that was solution-treated after 90% cold rolling as explained bellow. In figure 4 (a) and table 2, the X-ray peak intensity and the intensity ratios are shown of the ST specimen solution-treated after 90% cold rolling. Here, *IRst(1)* and *IRst(2)* were defined as *IRst(1)= I{110}β / I{200}β* and *IRst(2)= I{211}β / I{200}β,* respectively. Both of them de‐ creased monotonously with increasing β-phase stability. Namely, in the recrystallization textures the high atomic density planes, i.e., {110} and {211}, aligned parallel to the rolling plane preferentially. This trend further increased with decreasing β- phase stability, as

the β/β+ω+(α") boundary [20].

enhanced to form the rolling texture [32].

develop as a result of this reheating.

shown in table 2 for Z-alloys.

**4.3. Textures developed by recrystallization**

**Figure 6.** EBSD pole figures of alloy Z2 specimen after 90% cold rolling using 200β, (a), 110β, (b) and 112<sup>β</sup> (c) and the stereoprojection of the {100}β<110>β-type texture, (d).

**Figure 7.** XRD pole figures of alloy A01 specimen after 98% cold rolling using 200β, 110β, and 112β.

**Figure 8.** Inverse pole figures in the normal plane of alloy A01 specimen after 98% cold rolling calculated from the XRD ODF, (a), and alloy Z2 specimen after 90% cold rolling calculated from the EBSD ODF.

The β-type Ti-based alloys deform by either slip or twin mechanism [29,30]. The stress-in‐ duced martensitic transformation also takes place in some alloys upon applying external stress to them [30,31]. These phenomena emerge depending on the β-phase stability and hence will be controlled by alloying. Also, it is known that the slip/twin boundary is close to the β/β+ω+(α") boundary [20].

As crystal deforms by slip, it undergoes crystal rotations. Such rotations lead to the develop‐ ment of preferred orientations in polycrystalline alloys [32]. The main reason why the {100}<110> texture develops by cold rolling in conventional bcc alloys may be attributable to the glide of dislocations along <111> on {011}, {112}or {123}[33] and the crystal rotations in them. It is well known that C' represents the resistance for the {011}<011> shear and C<sup>44</sup> rep‐ resents the resistance for the {001}<001> shear [34] and both of them decreases with decreas‐ ing β-phase stability [24,35]. The e/a value is kept at 4.24 in the present alloys and C' decreases with decreasing β-phase stability as discussed earlier, so the elastic softening will be enhanced in the order; Z1 < Z2 < Z 3 < Z4. This is a reason why the alloys locating at the β/β+α" boundary showed a very low shear modulus along both <011> on {011} and along <111> on {011}, {112} or {123} as reported in Ref. [36]. So it is expected that as the β-phase stability decreases, secondary slipping systems such as {011}<011> may be activated by de‐ formation beside the main slipping system, i.e., <111> on {011}, {112}or {123}. As a result, a portion of the applied stress is consumed in slipping in such secondary slipping systems, which will make some disturbances in forming the main {100}<110> rolling texture. This can be a reason why the {100}<110> rolling texture was dominant in the high β-phase stability alloys. This was also seen in the steel in which the addition of a ferrite (bcc) stabilizer, Si, is enhanced to form the rolling texture [32].

#### **4.3. Textures developed by recrystallization**

**Figure 6.** EBSD pole figures of alloy Z2 specimen after 90% cold rolling using 200β, (a), 110β, (b) and 112<sup>β</sup> (c) and the

**Figure 7.** XRD pole figures of alloy A01 specimen after 98% cold rolling using 200β, 110β, and 112β.

**Figure 8.** Inverse pole figures in the normal plane of alloy A01 specimen after 98% cold rolling calculated from the

The β-type Ti-based alloys deform by either slip or twin mechanism [29,30]. The stress-in‐ duced martensitic transformation also takes place in some alloys upon applying external stress to them [30,31]. These phenomena emerge depending on the β-phase stability and

XRD ODF, (a), and alloy Z2 specimen after 90% cold rolling calculated from the EBSD ODF.

stereoprojection of the {100}β<110>β-type texture, (d).

128 Recent Developments in the Study of Recrystallization

The textures developed by severe cold rolling will diminish if the specimen is reheated for long time. For example, the *IH(cr)* of a 98%CR A01 specimen was decreased when reheated to different temperatures for 3 hrs, as shown in figure 9. *IH(cr)* is defined as *IHcr=* 1/*IRcr= I{110}β / I{200}β*. Hence, it is expected that the {100}<110> rolling texture is diminishing as a result of this reheating. The decrease in *IHcr* was higher with increasing the reheating temperature till it reach a saturation temperature. This saturation temperature is most probably related to the completion of recrystallization process. Therefore, further increase in temperature will not affect much the value of *IHcr*, as shown in figure 9. The rolling textures will not only dimin‐ ish the heavily cold rolled specimens by reheating, but also recrystallization textures may develop as a result of this reheating.

In the present designed alloys, the recrystallization textures were developed in different strength in the ST specimen that was solution-treated after 90% cold rolling as explained bellow. In figure 4 (a) and table 2, the X-ray peak intensity and the intensity ratios are shown of the ST specimen solution-treated after 90% cold rolling. Here, *IRst(1)* and *IRst(2)* were defined as *IRst(1)= I{110}β / I{200}β* and *IRst(2)= I{211}β / I{200}β,* respectively. Both of them de‐ creased monotonously with increasing β-phase stability. Namely, in the recrystallization textures the high atomic density planes, i.e., {110} and {211}, aligned parallel to the rolling plane preferentially. This trend further increased with decreasing β- phase stability, as shown in table 2 for Z-alloys.

decrease in the β-phase stability. So, it was likely that the oriented nucleation and/or ori‐ ented growth was enhanced with decreasing β-phase stability, leading to the increase in

Texturing Tendency in β-Type Ti-Alloys http://dx.doi.org/10.5772/53588 131

From these discussions and the data given in table 2, it was concluded that, for the single βphase alloys, the tendency of forming the {100}<110> texture by cold rolling increased with increasing β-phase stability, whereas, for low β-phase stability alloys (such as A00 and Z4 alloys in this study) other rolling textures may develop. On the other hand, the tendency of forming the recrystallization textures increased with decreasing β-phase stability. Thus, the β-phase stability was operating in the completely reverse way between the rolling and the

**Figure 10.** φ sections of the EBSD Euler space plot of alloys Z2-4 specimens solution treated at 1223K for 1.8 Ks after

The co-existence of α"- and/or ω-phase in the low β-phase stability alloys, such as Z4, A00 and A01 alloys in the present study, seems to affect much the deformation process and

**4.4. Textures developed in low β-phase stability alloys**

the strength of the recrystallization texture.

recrystallization textures.

90% cold rolling.

**Figure 9.** Effect of reheating temperature, for 3hrs, on the strength of formerly developed <100>{110} rolling texture after 98%CR, represented by *IHcr*= *1/IRcr*= *I{110} / I{200}*of the A01 alloy.

In figure 10, the EBSD Euler space plots are shown of the ST specimen of the alloys Z2-4 which were subjected to solution treatment at 1223K for 1.8 ks after 90% cold rolling. The Euler space density maps showed that the recrystallization textures were well developed in the alloy Z4, followed by the alloy Z3 and then the alloy Z2. Thus, the tendency of forming the recrystallization texture in the alloys changed in the order, Z4> Z3 > Z2, which was the opposite order of the rolling texture formed by cold rolling.

In the solution treatment after cold rolling, it has been reported that {112}<110> recrystalli‐ zation texture is mainly developed in the low β-phase stability alloys such as Ti-24Nb-3Al alloy, position B in figure 3, [29,37] and Ti-35mass%Nb-4mass% Sn alloy, position A in figure 3, [31,38]. Beside this {112}<110> recrystallization texture, the {110}<211> texture is developed in a little higher β-phase stability alloy, Ti-35mass%Nb-7.9mass%Sn, position C in figure 3, [38]. However, only the {112}<110> recrystallization texture appears in the Ti-22Nb-6Ta alloy, position D in figure 3, [24]. It is important to note here that the {211}<110> recrystallization texture is well developed in the alloy only after severe cold rolling (i.e., 95% and 99%) [24,39] and it tends to diminish with decreasing reduction ratio of cold rolling [24]. Also, recrystallization textures could be controlled by the temperature and time for the heat treatment [39].

Recrystallization is the replacement of deformed grains by the recrystallized grains [40]. The grains with certain crystallographic orientations will be nucleated and grown in the course of annealing [41]. The growth rate of the grains is also 'oriented', because some grains with certain crystallographic orientation will grow faster than others [32]. Other‐ wise, 'oriented' nucleation may control the final texture structure. As discussed earlier, both the elastic softening and the elastic anisotropy becomes more remarkable with the decrease in the β-phase stability. So, it was likely that the oriented nucleation and/or ori‐ ented growth was enhanced with decreasing β-phase stability, leading to the increase in the strength of the recrystallization texture.

From these discussions and the data given in table 2, it was concluded that, for the single βphase alloys, the tendency of forming the {100}<110> texture by cold rolling increased with increasing β-phase stability, whereas, for low β-phase stability alloys (such as A00 and Z4 alloys in this study) other rolling textures may develop. On the other hand, the tendency of forming the recrystallization textures increased with decreasing β-phase stability. Thus, the β-phase stability was operating in the completely reverse way between the rolling and the recrystallization textures.

**Figure 10.** φ sections of the EBSD Euler space plot of alloys Z2-4 specimens solution treated at 1223K for 1.8 Ks after 90% cold rolling.

#### **4.4. Textures developed in low β-phase stability alloys**

**Figure 9.** Effect of reheating temperature, for 3hrs, on the strength of formerly developed <100>{110} rolling texture

In figure 10, the EBSD Euler space plots are shown of the ST specimen of the alloys Z2-4 which were subjected to solution treatment at 1223K for 1.8 ks after 90% cold rolling. The Euler space density maps showed that the recrystallization textures were well developed in the alloy Z4, followed by the alloy Z3 and then the alloy Z2. Thus, the tendency of forming the recrystallization texture in the alloys changed in the order, Z4> Z3 > Z2, which was the

In the solution treatment after cold rolling, it has been reported that {112}<110> recrystalli‐ zation texture is mainly developed in the low β-phase stability alloys such as Ti-24Nb-3Al alloy, position B in figure 3, [29,37] and Ti-35mass%Nb-4mass% Sn alloy, position A in figure 3, [31,38]. Beside this {112}<110> recrystallization texture, the {110}<211> texture is developed in a little higher β-phase stability alloy, Ti-35mass%Nb-7.9mass%Sn, position C in figure 3, [38]. However, only the {112}<110> recrystallization texture appears in the Ti-22Nb-6Ta alloy, position D in figure 3, [24]. It is important to note here that the {211}<110> recrystallization texture is well developed in the alloy only after severe cold rolling (i.e., 95% and 99%) [24,39] and it tends to diminish with decreasing reduction ratio of cold rolling [24]. Also, recrystallization textures could be controlled by the temperature

Recrystallization is the replacement of deformed grains by the recrystallized grains [40]. The grains with certain crystallographic orientations will be nucleated and grown in the course of annealing [41]. The growth rate of the grains is also 'oriented', because some grains with certain crystallographic orientation will grow faster than others [32]. Other‐ wise, 'oriented' nucleation may control the final texture structure. As discussed earlier, both the elastic softening and the elastic anisotropy becomes more remarkable with the

of the A01 alloy.

 */ I{200}*

opposite order of the rolling texture formed by cold rolling.

after 98%CR, represented by *IHcr*= *1/IRcr*= *I{110}*

130 Recent Developments in the Study of Recrystallization

and time for the heat treatment [39].

The co-existence of α"- and/or ω-phase in the low β-phase stability alloys, such as Z4, A00 and A01 alloys in the present study, seems to affect much the deformation process and therefore the rolling texture developed in them when severly deformed. This is because in such low β-phase stability alloys, the deformation by twin and/or stress-induced martensitic transformation mechanisms are predominant. The deformation stress are consumed in forming twinning and inducing ω and/or α" marensite phases. As a result, the grains rota‐ tion is expected to be much less in these low β-phase stability alloys resulting in low strength of the {100}<110> texture. For example, Z4 and A01 alloys show low *Icr* and in the same time the stress induced α" martensitic transformation occurred in them by cold rolling, as evedinced from the two-steps yielding during the tensile test. For further details, refere to Refs. [25,42].

It is interesting to mention here that the low β-phase stability alloys, with α" martensite as the predominant phase, will show also a texture in the α"-phase, as shown in figure 11. The {200}α", {012}α" and {220}α" XRD pole figures obtained from a 98%CR specimen of the low βphase alloy,A00, are shown in Figure 11.a. The center of the pole figures corresponds to the direction normal to the specimen surface (ND). It seems from these pole figures that the α" phase was textured in the specimen by cold rolling in the way that the {220}α" planes aligned parallel to the rolling plane preferentially compared to the other textures and random grains orientations. In such low β-phase stability alloys, cold deformation induces aligned ω and/or α" martensite phases with orientation relationships as explained else where. The three α", ω, and β- phases are co-existed in A01 alloy after deformation as shown in figure 12. However, the ST specimen show less amount of α" martensite and the recrystalliztion texture in the α" –phase is also less as evidenced from figure 11.b for A00 alloy.

**Figure 12.** TEM analysis of a A01 specimen after 98CR. (a) and (b) present, respectively, bright-field and dark-field [im‐ aged on (-113)<sup>β</sup> ] pair of micrographs showing stress induced of the ω and/or α"- phases; (c) shows the corresponding

Texturing Tendency in β-Type Ti-Alloys http://dx.doi.org/10.5772/53588 133

The microstructures shown in figure 13 are of the alloys Z2-4 after 90% cold rolled (90CR) in the RD and TD and the ND cross sections. The stream-like deformation bands were ob‐ served clearly along the RD and TD directions in both the RD and TD cross-sections. As the β-phase stability decreased, the deformation bands seemed to become finer. Also the density of the deformation bands seemed to be higher in the TD cross-section shown in figure 13.a

The microstructure in an enlarged magnification is shown in figure 14 for the 90CR alloys. In the alloy 4, fine shear bands with an angle 42O inclined to the main deformation band were observed in the RD cross-section. However, no such shear bands were observed in the TD cross-section. Also in the alloy Z3, the secondary finer deformation bands with the an‐ gles of 23O and 37O inclined to the main deformation bands were observed in the RD and TD cross-sections, respectively. Neither the shear bands nor the secondary deformation bands were observed in the alloy Z2 in both the RD and TD cross-sections. These results can be interpreted, as above, due to the presence of the texturing systems other than {100}<110> de‐ veloped by the cold rolling and the plastic deformation among these three alloys depending on the β-phase stability. Also, the similarity in the TD and RD microstructure in the alloy Z2 is attributable to the well developed {100}<110> rolling texture in which both the RD and the

composite electron diffraction pattern and its key diagram. The ω-phase is with tow variants.

**4.5. Microstructure change with β-phase stability**

than in the RD cross-section shown in figure 13.b.

*4.5.1. The microstructures after cold rolling*

TD are parallel to <110> [35].

**Figure 11.** XRD pole figures of alloy A00 specimen after 98% cold rolling, (a) and after subsequent solution treatment at 1223K for 1.8 Ks, (b), using 200α", 012α", and 220 α".

**Figure 12.** TEM analysis of a A01 specimen after 98CR. (a) and (b) present, respectively, bright-field and dark-field [im‐ aged on (-113)<sup>β</sup> ] pair of micrographs showing stress induced of the ω and/or α"- phases; (c) shows the corresponding composite electron diffraction pattern and its key diagram. The ω-phase is with tow variants.

#### **4.5. Microstructure change with β-phase stability**

#### *4.5.1. The microstructures after cold rolling*

therefore the rolling texture developed in them when severly deformed. This is because in such low β-phase stability alloys, the deformation by twin and/or stress-induced martensitic transformation mechanisms are predominant. The deformation stress are consumed in forming twinning and inducing ω and/or α" marensite phases. As a result, the grains rota‐ tion is expected to be much less in these low β-phase stability alloys resulting in low strength of the {100}<110> texture. For example, Z4 and A01 alloys show low *Icr* and in the same time the stress induced α" martensitic transformation occurred in them by cold rolling, as evedinced from the two-steps yielding during the tensile test. For further details, refere to

It is interesting to mention here that the low β-phase stability alloys, with α" martensite as the predominant phase, will show also a texture in the α"-phase, as shown in figure 11. The {200}α", {012}α" and {220}α" XRD pole figures obtained from a 98%CR specimen of the low βphase alloy,A00, are shown in Figure 11.a. The center of the pole figures corresponds to the direction normal to the specimen surface (ND). It seems from these pole figures that the α" phase was textured in the specimen by cold rolling in the way that the {220}α" planes aligned parallel to the rolling plane preferentially compared to the other textures and random grains orientations. In such low β-phase stability alloys, cold deformation induces aligned ω and/or α" martensite phases with orientation relationships as explained else where. The three α", ω, and β- phases are co-existed in A01 alloy after deformation as shown in figure 12. However, the ST specimen show less amount of α" martensite and the recrystalliztion texture in the α"

**Figure 11.** XRD pole figures of alloy A00 specimen after 98% cold rolling, (a) and after subsequent solution treatment

–phase is also less as evidenced from figure 11.b for A00 alloy.

at 1223K for 1.8 Ks, (b), using 200α", 012α", and 220 α".

Refs. [25,42].

132 Recent Developments in the Study of Recrystallization

The microstructures shown in figure 13 are of the alloys Z2-4 after 90% cold rolled (90CR) in the RD and TD and the ND cross sections. The stream-like deformation bands were ob‐ served clearly along the RD and TD directions in both the RD and TD cross-sections. As the β-phase stability decreased, the deformation bands seemed to become finer. Also the density of the deformation bands seemed to be higher in the TD cross-section shown in figure 13.a than in the RD cross-section shown in figure 13.b.

The microstructure in an enlarged magnification is shown in figure 14 for the 90CR alloys. In the alloy 4, fine shear bands with an angle 42O inclined to the main deformation band were observed in the RD cross-section. However, no such shear bands were observed in the TD cross-section. Also in the alloy Z3, the secondary finer deformation bands with the an‐ gles of 23O and 37O inclined to the main deformation bands were observed in the RD and TD cross-sections, respectively. Neither the shear bands nor the secondary deformation bands were observed in the alloy Z2 in both the RD and TD cross-sections. These results can be interpreted, as above, due to the presence of the texturing systems other than {100}<110> de‐ veloped by the cold rolling and the plastic deformation among these three alloys depending on the β-phase stability. Also, the similarity in the TD and RD microstructure in the alloy Z2 is attributable to the well developed {100}<110> rolling texture in which both the RD and the TD are parallel to <110> [35].

#### *4.5.2. The micorstructure after solution treatment*

The microstructures of the ST alloys Z2-4 solution treated at 1223K for 1.8ks after 90% cold rolling are shown in figure 15. Shear bands supposedly provide the majority of nucleation sites during recrystallization in the severely cold rolled alloys. In particular, the triple joints between shear bands sites will work as nucleation sites [43]. So, the number of the nuclea‐ tion sites was supposed to increase with decreasing β-phase stability, judging from the for‐ mer deformation microstructures shown in figures 13 and 14. As a result, the grain size in the ST condition seems to decrease as the β-phase stability decreases. As is evident from the slightly defocused micrographs shown in figure 15.a some deformation bands still existed in the TD cross-section, although the recrystallization process seemed to be completed in the rolling plane shown in figure 15.b. However, the equiaxed grains were observed in the TD cross-section of the alloy Z2 in the ST condition in the EBSD image quality map (IQM) as presented in Ref. [25]. Similar deformation bands were observed in a Ti-45Nb alloy in the recrystallized condition after severe cold working [44]. The reason why these deformation bands still remained after the recrystallization treatment is not clear at the moment.

 **Z2 Z3 Z4** 

Texturing Tendency in β-Type Ti-Alloys http://dx.doi.org/10.5772/53588 135

**Figure 14.** SEM micrographs of alloys Z2-4 shown in the enlarged scale of figure 13 (a) and (b).

 **Z2 Z3 Z4** 

**Figure 15.** Micrographs of alloys Z2-4 after solution treatment at 1223K for 1.8 ks after 90% cold rolling in the transverse, TD, cross section, defocused SEM micrographs to show the deformation bands (a), and in the rolling plane, OM (b).

**Figure 13.** SEM micrographs of alloys Z2-4 after 90% cold rolling in the transverse, TD, cross-section (a), in the rolling, (RD), cross-section (b), and in the normal plane (c).

**Figure 14.** SEM micrographs of alloys Z2-4 shown in the enlarged scale of figure 13 (a) and (b).

*4.5.2. The micorstructure after solution treatment*

134 Recent Developments in the Study of Recrystallization

The microstructures of the ST alloys Z2-4 solution treated at 1223K for 1.8ks after 90% cold rolling are shown in figure 15. Shear bands supposedly provide the majority of nucleation sites during recrystallization in the severely cold rolled alloys. In particular, the triple joints between shear bands sites will work as nucleation sites [43]. So, the number of the nuclea‐ tion sites was supposed to increase with decreasing β-phase stability, judging from the for‐ mer deformation microstructures shown in figures 13 and 14. As a result, the grain size in the ST condition seems to decrease as the β-phase stability decreases. As is evident from the slightly defocused micrographs shown in figure 15.a some deformation bands still existed in the TD cross-section, although the recrystallization process seemed to be completed in the rolling plane shown in figure 15.b. However, the equiaxed grains were observed in the TD cross-section of the alloy Z2 in the ST condition in the EBSD image quality map (IQM) as presented in Ref. [25]. Similar deformation bands were observed in a Ti-45Nb alloy in the recrystallized condition after severe cold working [44]. The reason why these deformation

bands still remained after the recrystallization treatment is not clear at the moment.

 **Z2 Z3 Z4** 

**Figure 13.** SEM micrographs of alloys Z2-4 after 90% cold rolling in the transverse, TD, cross-section (a), in the rolling,

(RD), cross-section (b), and in the normal plane (c).

**m** 
