**1. Introduction**

## **1.1. Textures**

Preferred orientation of crystal is an intrinsic feature of metals and has an influence on phys‐ ical properties such as strength, electrical conductivity and wave propagation, particularly in the anisotropy of these properties [1]. For example, in the single crystals of many metals it is well known that the main mechanism of plastic deformation, on a microscopic scale, is a simple shear parallel to certain planes and directions. Slip will occur in a certain direction on a crystallographic plane when the shear stress in that direction attains a critical value [2]. So, the observed strength might depend on the loading direction of the crystal. Some other physical properties of the crystal vary depending on the measuring direction.

On the other hand, in a polycrystalline metal, each grain normally has a crystallographic ori‐ entation different from that of its neighbors resulting in isotropy of the properties of the metal. Considering a polycrystalline as a whole, the orientations of the grains may tend to cluster about some particular orientations. Such polycrystalline is said to have a preferred orientation, or texture, which may be defined simply as anon-random distribution of crystal orientations [3]. In this case, a polycrystalline performs in a way close to the single crystal depending on the strength of texture formed in it.

The appearance of preferred orientations (or texture) is very common. The texture produced by forming such as wire drawing and rolling, is called a deformation texture. The grains in a polycrystalline metal tend to rotate during plastic deformation, which results in the texture formation. Each grain undergoes slip and rotation in a complex way by the imposed force and by the restriction of slip and rotation of adjoining grains [3]. The preferred orientation also appear when cold-worked metal (show a deformation texture) is annealed. This is the so-called recrystallization texture (or annealing texture).

© 2013 Gepreel; licensee InTech. This is an open access article distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2013 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

Texture may be influenced by a number of factors that can be divided into two major cate‐ gories, material variables and process variables [4]. The material variables include crystal structure, solute content, second-phase particles, and initial grain size. The process variables include the amount of deformation, strain rate, stress and strain states [5], reduction in thickness and area, intermediate annealing conditions, heating and cooling rates, and an‐ nealing atmosphere [6].

The weakest slip systems in fcc metals are{111}<110>, consisting of 12 symmetrically equiva‐ lent variants. Activity of these systems produce a characteristic texture pattern during roll‐ ing [1]. The large number of slip systems makes it easy to achieve compatibility and the rolling texture can be well explained with the Taylor theory, especially if individual slip sys‐ tems are allowed to harden, according to their activity [8] and if allowance is made for some

Texturing Tendency in β-Type Ti-Alloys http://dx.doi.org/10.5772/53588 119

Also, the bcc metals have 12 variants of slipping system [1] and the most common deforma‐ tion mode in bcc metals is {110}<111> slip, which is a transposition of slip plane and slip di‐ rection with respect to fcc metals. Also, bcc metals slip on other planes than {110} in the <111>direction [10] such as {112}<111>and {123}<111> [3]. Generally, the most preferred ori‐ entations in cold rolled bcc metals consist of two types. The first type is represented by {100}<110> or its rotations around<110> axis, resulting in {hkl}<110>. The other is represented by {111}<112> or its orientations around surface normal axis, resulting in {hkl}<uvw> [3].

The relationship between slip and crystal rotations is straightforward. Other processes such as climb, grain boundary sliding, and diffusion in general may also affect orientation distri‐ butions [1]. Of particular importance is recrystallization. Recrystallization is the replacement of deformed grains by the nucleation and growth of recrystallized grains, both can lead to drastic changes in texture. The nuclei during recrystallization are regions exists in the de‐ formed microstructure and at the same time the recrystallization does not lead to precise ori‐ entation relationship between deformed and recrystallized grains. Although there might be some approximate orientation relationships present in-between the recrystallized and pa‐ rent grains, there has been no report in making quantitative predictions for recrystallization texture based on such approximate orientation relationships. The texture evolution is still highly sensitive to almost the entire spectrum of metallurgical variables and sometimes dif‐ ficult to fully predict. However sometimes it is observed that the recrystallization texture component bear crystallographic relationships to the original deformation textures, which may be described by rotations about simple crystal direction which is often ~25o rotations

In recrystallization, the nuclei by the shear bands competes with the nuclei forming at other nucleation sites. Since stored energy represents the driving force for recrystalliza‐ tion, certain crystallographic orientations will be enhanced during annealing in compari‐ son with others because of more favorable nucleation and/or growth kinetics [12]. Grains with higher stored energies may be consumed through boundary migration by grains with less stored energy (i.e., growth stage). Alternatively, dislocation-free nuclei may form grains with high dislocation density and then grow at expense of others. The texture which finally forms is believed to be controlled either by nucleation or by grain growth. It is possible that grain growth is 'oriented', i.e. for some reasons grains with certain crystal‐ lographic orientation grow faster than others. In that case the grain growth mechanism is likely to control the final texture. Otherwise, 'oriented' nucleation may control the final texture. The texture fresh nuclei may not be random and reflect the final texture. The last

heterogeneity across small grain boundaries [9].

*1.3.2. Recrystallization texture*

around <110> directions in bcc alloys [11].

#### **1.2. Texture characterization techniques**

The most common method of characterizing texture is the presentation of pole figures where X-ray diffraction is used to specify the orientation of the crystallographic planes in space using the stereographic projection. The two different representations of these tex‐ tures are the inverse and direct pole figures [1]. The pole figure is a two-dimensional pro‐ jection of the three-dimensional distribution and represents the probability of finding a pole to a lattice plane (*hkl*) in a certain sample direction. Pole figures are normalized to ex‐ press this probability in multiples of a random distribution. Inverse pole figures are also projections of the orientation distribution factor (ODF), but in this vase the probability of finding a sample direction relative to crystal directions is plotted. In other words, the dif‐ ference between these two representations is their reference frame; inverse pole figures are shown with respect to the specimen reference frame, while direct pole figures refer to the crystal reference frame [1].

Local orientation can also be measured with the Scanning Electron Microscope (SEM). Inter‐ action of the electron beam with the uppermost surface layer of the sample produces elec‐ tron back-scatter diffraction patterns (EBSPs or EBSD) that are analogous to kikuchi patterns in Transmission Electron Microscope (TEM). EBSD are captured on a phosphorus screen and recorded with a low intensity video camera or a CCD device [1].

The texture of sheet is in the most highly developed form, so, most of grains are oriented and the sheet surface becomes roughly parallel to a certain crystallographic plane (hkl). Al‐ so, a certain direction [uvw] in that plane becomes roughly parallel to the direction of defor‐ mation. Such a texture is described by the shorthand notation (hkl)[uvw], and is called an ideal orientation. Most sheet textures, however, have so much scatter in orientations, and hence they may be approximated symbolically only by the sum of number of ideal orienta‐ tions or texture components [3].

### **1.3. Texture types**

## *1.3.1. Deformation texture*

The deformation of a polycrystal is a very complicated heterogeneous process. When an ex‐ ternal stress is applied to the polycrystal, it is transmitted to individual grains. Dislocations move on slip systems, dislocations interact and cause 'hardening', grains change their shape and orientation, thereby interacting with neighbours and creating local stresses that need to be accommodated [1].

The weakest slip systems in fcc metals are{111}<110>, consisting of 12 symmetrically equiva‐ lent variants. Activity of these systems produce a characteristic texture pattern during roll‐ ing [1]. The large number of slip systems makes it easy to achieve compatibility and the rolling texture can be well explained with the Taylor theory, especially if individual slip sys‐ tems are allowed to harden, according to their activity [8] and if allowance is made for some heterogeneity across small grain boundaries [9].

Also, the bcc metals have 12 variants of slipping system [1] and the most common deforma‐ tion mode in bcc metals is {110}<111> slip, which is a transposition of slip plane and slip di‐ rection with respect to fcc metals. Also, bcc metals slip on other planes than {110} in the <111>direction [10] such as {112}<111>and {123}<111> [3]. Generally, the most preferred ori‐ entations in cold rolled bcc metals consist of two types. The first type is represented by {100}<110> or its rotations around<110> axis, resulting in {hkl}<110>. The other is represented by {111}<112> or its orientations around surface normal axis, resulting in {hkl}<uvw> [3].

#### *1.3.2. Recrystallization texture*

Texture may be influenced by a number of factors that can be divided into two major cate‐ gories, material variables and process variables [4]. The material variables include crystal structure, solute content, second-phase particles, and initial grain size. The process variables include the amount of deformation, strain rate, stress and strain states [5], reduction in thickness and area, intermediate annealing conditions, heating and cooling rates, and an‐

The most common method of characterizing texture is the presentation of pole figures where X-ray diffraction is used to specify the orientation of the crystallographic planes in space using the stereographic projection. The two different representations of these tex‐ tures are the inverse and direct pole figures [1]. The pole figure is a two-dimensional pro‐ jection of the three-dimensional distribution and represents the probability of finding a pole to a lattice plane (*hkl*) in a certain sample direction. Pole figures are normalized to ex‐ press this probability in multiples of a random distribution. Inverse pole figures are also projections of the orientation distribution factor (ODF), but in this vase the probability of finding a sample direction relative to crystal directions is plotted. In other words, the dif‐ ference between these two representations is their reference frame; inverse pole figures are shown with respect to the specimen reference frame, while direct pole figures refer to

Local orientation can also be measured with the Scanning Electron Microscope (SEM). Inter‐ action of the electron beam with the uppermost surface layer of the sample produces elec‐ tron back-scatter diffraction patterns (EBSPs or EBSD) that are analogous to kikuchi patterns in Transmission Electron Microscope (TEM). EBSD are captured on a phosphorus screen

The texture of sheet is in the most highly developed form, so, most of grains are oriented and the sheet surface becomes roughly parallel to a certain crystallographic plane (hkl). Al‐ so, a certain direction [uvw] in that plane becomes roughly parallel to the direction of defor‐ mation. Such a texture is described by the shorthand notation (hkl)[uvw], and is called an ideal orientation. Most sheet textures, however, have so much scatter in orientations, and hence they may be approximated symbolically only by the sum of number of ideal orienta‐

The deformation of a polycrystal is a very complicated heterogeneous process. When an ex‐ ternal stress is applied to the polycrystal, it is transmitted to individual grains. Dislocations move on slip systems, dislocations interact and cause 'hardening', grains change their shape and orientation, thereby interacting with neighbours and creating local stresses that need to

and recorded with a low intensity video camera or a CCD device [1].

nealing atmosphere [6].

**1.2. Texture characterization techniques**

118 Recent Developments in the Study of Recrystallization

the crystal reference frame [1].

tions or texture components [3].

**1.3. Texture types**

*1.3.1. Deformation texture*

be accommodated [1].

The relationship between slip and crystal rotations is straightforward. Other processes such as climb, grain boundary sliding, and diffusion in general may also affect orientation distri‐ butions [1]. Of particular importance is recrystallization. Recrystallization is the replacement of deformed grains by the nucleation and growth of recrystallized grains, both can lead to drastic changes in texture. The nuclei during recrystallization are regions exists in the de‐ formed microstructure and at the same time the recrystallization does not lead to precise ori‐ entation relationship between deformed and recrystallized grains. Although there might be some approximate orientation relationships present in-between the recrystallized and pa‐ rent grains, there has been no report in making quantitative predictions for recrystallization texture based on such approximate orientation relationships. The texture evolution is still highly sensitive to almost the entire spectrum of metallurgical variables and sometimes dif‐ ficult to fully predict. However sometimes it is observed that the recrystallization texture component bear crystallographic relationships to the original deformation textures, which may be described by rotations about simple crystal direction which is often ~25o rotations around <110> directions in bcc alloys [11].

In recrystallization, the nuclei by the shear bands competes with the nuclei forming at other nucleation sites. Since stored energy represents the driving force for recrystalliza‐ tion, certain crystallographic orientations will be enhanced during annealing in compari‐ son with others because of more favorable nucleation and/or growth kinetics [12]. Grains with higher stored energies may be consumed through boundary migration by grains with less stored energy (i.e., growth stage). Alternatively, dislocation-free nuclei may form grains with high dislocation density and then grow at expense of others. The texture which finally forms is believed to be controlled either by nucleation or by grain growth. It is possible that grain growth is 'oriented', i.e. for some reasons grains with certain crystal‐ lographic orientation grow faster than others. In that case the grain growth mechanism is likely to control the final texture. Otherwise, 'oriented' nucleation may control the final texture. The texture fresh nuclei may not be random and reflect the final texture. The last word has not been yet said on this problem. Not only orientation distribution, but also grain size distribution are important considerations [1].

**1.5. BCC β-type Ti alloys**

finish temperatures, respectively.

Recently, considerable efforts have been devoted to exploring novel β-titanium alloys for different applications because of their superior properties such as the superelasticity, low Young's modulus, high strength-to-weight ratio, and better formability compared to the α and α + β titanium alloys [17-20]. The mechanical properties of the β-titanium alloys depend strongly on the presence of several phases (e.g., ω phase and martensitic α"-phase) in them. The appearance of these phases could be controlled by either the optimized alloy design [19,20] or the materials processing [18]. The phases present in the alloy are related to the al‐ loying elements and the thermal history of the alloy. The change of the Ti-alloy type de‐

Texturing Tendency in β-Type Ti-Alloys http://dx.doi.org/10.5772/53588 121

**Figure 2.** Schematic phase diagram of titanium alloys with the decomposition products of the β-phase. βc is the critical minimum β stabilizer amount for near β or metastable β alloys to retain β completely on quenching from β phase field and βs is the minimum amount of β stabilizer for stable β alloys; β*trs*, *M*s and *M*f refer to β-transus, martensite start and

pending on the β-stabilizer content is shown schematically in figure 2.

The stored energy is the most important parameter characterizing the influence of micro‐ structure on the recryasallization process which is a driving force of it. The stored energy is proportional to the average critical shear stress for slip in the crystal [13].

#### **1.4. Elastic anisotropy change with β-phase stability**

The elastic anisotropy of bcc metals can be represented by the elastic anisotropy factor, A, which is calculated from elastic constants from the formula, A= C44/C'. C'= (C11-C12)/2 repre‐ sents the resistance for the {110}<110> shear and C44 represents the resistance for the {001}<001> shear [13] and both of them increases with β-phase stability [3,10]. Further more, the ideal strength for tension σmax, which is needed for tensile separation of bcc crystal on {100} plane, is proportional to C' and ideal strength for shear σmax which is needed for plastic shear along <111> on {011}, {112} or {123} [14]. As a result, the deformation mechanism in bcc alloys is relat‐ ed to the β-phase stability as mentioned before [15]. Figure 1 show the change in A and C' with the Cr-content in Ti-Cr binary alloys. The Cr-content is related to β-phase stability [16]. So, it is expected that the anisotropy and texturing tendency are related to the β-phase stability in β Tialloys. So, a correlation between the β-phase stability and the texturing tendency will be pre‐ sented below.

**Figure 1.** Changes in elastic anisotropy factor, *A*, and tetragonal shear constant, C', with the β −phase stability in Ti-Cr binary alloys. *A* shows maximum and C' shows minimum at the β/β+ω phase boundary

#### **1.5. BCC β-type Ti alloys**

word has not been yet said on this problem. Not only orientation distribution, but also

The stored energy is the most important parameter characterizing the influence of micro‐ structure on the recryasallization process which is a driving force of it. The stored energy is

The elastic anisotropy of bcc metals can be represented by the elastic anisotropy factor, A, which is calculated from elastic constants from the formula, A= C44/C'. C'= (C11-C12)/2 repre‐ sents the resistance for the {110}<110> shear and C44 represents the resistance for the {001}<001> shear [13] and both of them increases with β-phase stability [3,10]. Further more, the ideal strength for tension σmax, which is needed for tensile separation of bcc crystal on {100} plane, is proportional to C' and ideal strength for shear σmax which is needed for plastic shear along <111> on {011}, {112} or {123} [14]. As a result, the deformation mechanism in bcc alloys is relat‐ ed to the β-phase stability as mentioned before [15]. Figure 1 show the change in A and C' with the Cr-content in Ti-Cr binary alloys. The Cr-content is related to β-phase stability [16]. So, it is expected that the anisotropy and texturing tendency are related to the β-phase stability in β Tialloys. So, a correlation between the β-phase stability and the texturing tendency will be pre‐

**Figure 1.** Changes in elastic anisotropy factor, *A*, and tetragonal shear constant, C', with the β −phase stability in Ti-Cr

binary alloys. *A* shows maximum and C' shows minimum at the β/β+ω phase boundary

grain size distribution are important considerations [1].

120 Recent Developments in the Study of Recrystallization

**1.4. Elastic anisotropy change with β-phase stability**

sented below.

proportional to the average critical shear stress for slip in the crystal [13].

Recently, considerable efforts have been devoted to exploring novel β-titanium alloys for different applications because of their superior properties such as the superelasticity, low Young's modulus, high strength-to-weight ratio, and better formability compared to the α and α + β titanium alloys [17-20]. The mechanical properties of the β-titanium alloys depend strongly on the presence of several phases (e.g., ω phase and martensitic α"-phase) in them. The appearance of these phases could be controlled by either the optimized alloy design [19,20] or the materials processing [18]. The phases present in the alloy are related to the al‐ loying elements and the thermal history of the alloy. The change of the Ti-alloy type de‐ pending on the β-stabilizer content is shown schematically in figure 2.

**Figure 2.** Schematic phase diagram of titanium alloys with the decomposition products of the β-phase. βc is the critical minimum β stabilizer amount for near β or metastable β alloys to retain β completely on quenching from β phase field and βs is the minimum amount of β stabilizer for stable β alloys; β*trs*, *M*s and *M*f refer to β-transus, martensite start and finish temperatures, respectively.

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 the elastic and plastic properties.

As discussed above, Zr works as β-stabilizer in the β-type Ti alloys and also raises the *<sup>B</sup>*¯

value of the alloy. The 2nd step was to add Zr to the alloys with different amounts varies with the required β-phase stability of each alloy. So, the stability difference between the al‐ loys will be much bigger and the properties difference would be clearer. Considering that

**Figure 3.** Extended *B*¯*o*-*M*¯*<sup>d</sup>* diagram showing the β/β+ω+(α") phase boundary and the location of the designed alloys, A00 and Z1-4. Also, the alloys A, B, C, and D are 35mass%Nb-4mass%Sn, Ti-24Nb-3Al, Ti-35mass%Nb-7.9mass% Sn,

At last, the four designed Z-alloys, namely Z1 - 4, are located across the β/β+ω+(α") phase

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

diagram shown in figure 3, and the chemical compositions are listed in Table 1.

*<sup>o</sup>* region (i.e., high Zr-Containing alloys zone) in the extended *<sup>B</sup>*¯

the e/a value of the alloys did not change after Zr addition.

and Ti-22Nb-6Ta, respectively.

boundary in the high *<sup>B</sup>*¯

**3. Experimental procedure**

specimen is called ST specimen hereafter.

*o*

123

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

> *o*-*M*¯ *d*

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 design of these alloys is explained below.
