**3. Experimental procedure**

Most of the β-Ti alloys possess good workability. It is possible to fabricate a cold-rolled sheet of the alloys by a reduction ratio higher than 90%. In this case strong deformation tex‐ tures are developed and even recrystallization texture may be developed when subsequent‐ ly heat treated. Therefore, the anisotropy in elastic and plastic properties is induced inevitably to the sheet, resulting in the modification of alloy properties such as the elastic modulus, elastic strain, Poisson's ratio, strength, ductility, toughness, magnetic permeability and the energy of magnetization [21]. In other words, the elastic and plastic properties of the alloy may be improved by using an orientation effect arising from the textures. It is, there‐ fore, important to examine which kind of textures can be developed in the β-Ti alloys under the given conditions of thermo-mechanical treatment, and to investigate the texture effect on

In this chapter, the effect of β-phase stability on the texturing tendency of β-type Ti-alloys are discussed. It is important to high light here that Zr has been known for many decades as neutral element on the stability of β-phase; however, the recent studies have proved that Zr shows a β-stabilizing effect in the β-type alloys [20,22]. Therefore, in this study, it was chos‐ en to study two groups of alloys, one group is Zr-free β-type alloys (referred hereafter as Aalloys) and the other group is high Zr-containing alloys (referred herefater as Z-alloys). The

The two series of alloys were designed across the single β-phase boundary, β/β+ω+(α"),

*A=C44/C'*, is rather high since the value of the elastic shear modulus, *C' = (C11-C12)/2*, is di‐ minishing as the alloy approaches this boundary [23,24], as shown in figure 1. Also, it has been reported that the *(C11-C12)/2* is related to the electronic parameter *e/a* (electron-peratom ratio) and its value approaches zero when the e/a value is about 4.24 [23]. This is the reason why, in this work, the *e/a* value was kept at 4.24 in almost all the designed alloys.

In A-alloys, A00 alloy is designed to be located in the β+α"(+ω) phase zone in the *<sup>B</sup>*¯

gram [20], as shown in figure 3. The composition of A-alloys is listed in table 1. The Fe and O were added to stabilize the β-phase of the alloys. Fe was chosen to stabilize β-phase due to its very strong β-stabilizing effect as obvious from its alloying vector in the *<sup>B</sup>*¯

gram [22]. The oxygen was added to the alloys to suppress the ω and martensite phases, as

+ω+(α") phase boundary in two steps; in the 1st step, four alloys were designed, consist mainly of β-stabilizers, Ta, Nb, Mo, Cr and V. These four alloys are shown across the β/β+ω+

*d* diagram. At this boundary, the elastic anisotropy factor,

*o*-*M*¯

*d*is the average d-orbital energy

*o*-*M*¯ *d*dia‐

*o*-*M*¯ *d*dia‐

*d* diagram across the β/β

the elastic and plastic properties.

122 Recent Developments in the Study of Recrystallization

design of these alloys is explained below.

level (eV) of the elements in the alloy.

*o*- *M*¯

*<sup>o</sup>*is the average bond order between atoms, and *<sup>M</sup>*¯

As for Z-alloys, four alloys were designed with the aid of the *<sup>B</sup>*¯

(α") phase boundary with the dashed arrow in figure 3.

**2. Alloys design**

discussed in ref.[20].

Here, *<sup>B</sup>*¯

with the aid of the *<sup>B</sup>*¯

As explained above, two series of alloys, namely; high Zr-containing [Z-alloys] and Zr-free [A-alloys], were designed across the β/β+ω+(α") phase boundary and the chemical composi‐ tions are listed in table 1. In this chapter, all the compositions are given in atomic percent units unless otherwise noted. These alloys were prepared by the arc-melting of an appropri‐ ate mixture of pure metals (purity: 99.99%) under a high purity argon gas atmosphere. The button-shaped specimens with average 7.5 mm in thickness were cut and homogenized at 1273K for 7.2ks, and then cold rolled to the plate with 4.5 mm thick, followed by the solution treatment at 1223K for 1.8 ks. Subsequently, the specimen was cold rolled by 60%, 90 % or 98% reduction in thickness. The cold rolled specimen is called CR specimen hereafter. The 90%CR specimen was then solution-treated at 1223K for 1.8ks. This finally solution treated specimen is called ST specimen hereafter.

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 acceleration voltage of 20 kV and a tilt angle of 71o . The microstructural characterization was performed using the optical microscope (OM), the scanning electron microscope (SEM) and the transmission electron microscopy (TEM).


**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.

**CR, IRcr= I{200} β /I{110}β** 1.7 1.2 0.8 0.1 0.086 0.194 1.232

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‐

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

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

**ST, IRst(1)= I{110} β /I{200}β** 6.0 8.6 8.5 33.0 **ST, IRst(2)= I{211} β /I{200}β** 1.6 2.0 2.2 3.6

ments to get a β single phase [20,25]. According to the *<sup>B</sup>*¯

Z1 and lowest in alloy Z4.

**4.2. Textures developed by cold rolling**

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

**Z-alloys (60%CR) A-alloys (90%CR)**

**Z1 Z2 Z3 Z4 A00 A01 A11**

*o*-*M*¯

*d* diagram shown in figure 3, the

Texturing Tendency in β-Type Ti-Alloys http://dx.doi.org/10.5772/53588 125

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