**3.2 Microtexture**

442 Recent Trends in Processing and Degradation of Aluminium Alloys

and restoration processes, which occur even at room temperature by large strains (Estrin et al., 1998), the microstructural evolution and flow stress are dependent on the deformation

where Q is the activation energy for operative process and R is the gas constant (Zener &

Fig. 13. Misorientation distribution of the specimens processed by various cycles of ASB

Generally, by increasing this parameter, subgrain size at dynamic recovery and steady state grain size at dynamic recrystallization are decreased and the flow stress is increased (Humphreys & Hatherly, 2004). Tsuji et al. (2003b) have studied the effect of strain rate on microstructural changes in deformation of the ultra-fine grained (UFG) aluminum produced by ARB. They had expected that higher speed of deformation can result in finer grain subdivision, because larger amount of dislocations would operate at higher strain rates. However, their results showed that grain size rather increases with increasing the strain rate, due to heat of deformation. Therefore, they concluded that higher strain rate and simultaneous cooling are favorable for producing finer grain size in SPD. However, as deformations with strain rates higher than 10 s−1 are adiabatic processes (Ryazanov et al., 2003), simultaneous cooling is not effective on deformation temperature and therefore,

Controlled cooling of the specimen in tensile test maintains the normal temperatures of the tests conducted at 10−2 to 10−1 s−1. For higher strain rates, the only way is to perform the test

ε

ε

) and deformation temperature (T) are

exp (Q/RT) (1)

temperature, strain and strain rate. The strain rate (

Z =

comparing to the Mackenzie distribution

conclusion of Tsuji et al. (2003b) cannot be correct.

Hollomon, 1944).

included into the Zener–Hollomon Parameter (Z), which is defined as:

Figure 14 shows φ2 = 45, 65 and 90° sections of the initial orientation distribution functions (ODFs) of both aluminium sheets before the ARB process (fully annealed sheets). It is seen that the initial texture of these samples is mainly composed of a dominant rotated cube {001}<100> which is characteristic of the recrystallisation state. The ODFs corresponding to the samples of Figs. 3 and 4 are illustrated in Figs. 15 and 16 respectively. Only the φ2 = 45 and 90° sections of ODF are represented because these sections contain the most important texture components for ARB processed aluminium sheets. It is seen that the main texture components in the mid layers of both materials are Copper {1 1 2}<1 1 1> located at φ1, Φ, φ<sup>2</sup> ≡ 45, 65 and 90° and Dillamore {4 4 11}<11 11 8> at φ1, Φ, φ2 ≡ 90, 27 and 45° sections and there is no evidence of shear components. The ODFs of these ARB processed samples were generally similar and the overall texture intensity and the concentration of the contour lines increase continuously with increasing strain. However, it should be noted that in comparison with the AA3003 alloy, a more severe texture was observed in AA1100 alloy during the ARB process and the strength of the overall textures developed at the eighth cycle was 18.4 and 28.2 times the random levels, respectively.

The presence of second phase particles can also change the microtextural evolution in the ARB processed aluminium sheets. The comparison of the ODFs presented in Figs. 15 and 16 indicates that the developed texture in the AA1100 alloy is much stronger than the particle containing alloy. The evolution of a strong texture during the ARB process is attributed to the cyclic nature of this process. Heason and Prangnell (2002a) investigated the texture evolution during the ARB process of the AA1100 aluminium sheets by using a model based on the FC-Taylor approach. According to their model, the surface shear texture which deforms in plane strain compression (in the mid layer) rotates towards the copper orientation during the next cycle. In contrast, this model predicts that all the rolling

Nanostructure, Texture Evolution and Mechanical Properties of

overall texture intensity will be much weaker than the AA1100 alloy.

Aluminum Alloys Processed by Severe Plastic Deformation 445

components (copper, S and brass) rotate towards the shear orientation when they come to the surface of the sheet. This can also explain the increase in the intensity of copper and Dillamore components (which are close to each other) as the main texture components observed in this study. These authors have also reported that this strong texture leads to the development of coarse unrefined bands and prevents the full grain refinement to a submicrometre scale during the ARB process of a single phase aluminium alloy. Whereas, the presence of the second phase particles in the AA3003 alloy results in large local lattice rotation and can cause different textures compared with the rest of the matrix. Thus, the

Fig. 16. φ2=45 and 90° sections of ODFs corresponding to AA3003 samples processed by

Kim et al. (2005) reported similar results in ARB processed AA8011 aluminium alloy. According to their study, the textural change from a shear texture to a rolling texture at the sheet centre during the ARB process contributed to an increase in the fraction of high angle boundaries. Also, a large number of second phase particles in the AA8011 alloy sheets weakened the texture intensity which is due to the inhomogeneous deformation around the second phase particles. A weak and more spread texture caused by the presence of the

various ARB cycles

Fig. 14. φ2=45, 65 and 90° sections of ODF of fully annealed (a) AA3003 and (b) AA1100 aluminium sheets before ARB process: standard texture components are indicated with symbols: Cu: cube; R.Cu: rotated cube; C: copper; B: brass; G: Goss; D: Dillamore

Fig. 15. φ2=45 and 90° sections of ODFs corresponding to AA1100 samples processed by various ARB cycles: these ODFs were calculated assuming orthorhombic sample symmetry

Fig. 14. φ2=45, 65 and 90° sections of ODF of fully annealed (a) AA3003 and (b) AA1100 aluminium sheets before ARB process: standard texture components are indicated with symbols: Cu: cube; R.Cu: rotated cube; C: copper; B: brass; G: Goss; D: Dillamore

Fig. 15. φ2=45 and 90° sections of ODFs corresponding to AA1100 samples processed by various ARB cycles: these ODFs were calculated assuming orthorhombic sample symmetry components (copper, S and brass) rotate towards the shear orientation when they come to the surface of the sheet. This can also explain the increase in the intensity of copper and Dillamore components (which are close to each other) as the main texture components observed in this study. These authors have also reported that this strong texture leads to the development of coarse unrefined bands and prevents the full grain refinement to a submicrometre scale during the ARB process of a single phase aluminium alloy. Whereas, the presence of the second phase particles in the AA3003 alloy results in large local lattice rotation and can cause different textures compared with the rest of the matrix. Thus, the overall texture intensity will be much weaker than the AA1100 alloy.

Fig. 16. φ2=45 and 90° sections of ODFs corresponding to AA3003 samples processed by various ARB cycles

Kim et al. (2005) reported similar results in ARB processed AA8011 aluminium alloy. According to their study, the textural change from a shear texture to a rolling texture at the sheet centre during the ARB process contributed to an increase in the fraction of high angle boundaries. Also, a large number of second phase particles in the AA8011 alloy sheets weakened the texture intensity which is due to the inhomogeneous deformation around the second phase particles. A weak and more spread texture caused by the presence of the

Nanostructure, Texture Evolution and Mechanical Properties of

ARB process is characterized by the shear coefficient K (K= εxz/εzz).

with the experimentally measured ODFs (Figs. 16 and 18).

component in the middle layers of the ARB sample.

responsible for the unique property.

**3.2.1 Modeling of texture evolution** 

Aluminum Alloys Processed by Severe Plastic Deformation 447

It shows that in the early stages of the ARB process the intensity of the β fiber, running from copper over S to the brass orientation, is almost homogeneous, but with increasing the number of cycles the intensity of the copper component gradually grows at the expense of the brass and S components. The τ fiber is presented in Fig. 17c. It is observed that in this material the texture evolution is mainly centered around the copper and Dillamore components, and with an increasing number of cycles these components further intensify.

In order to interpret the texture change involved in the ARB process, the texture development was simulated with the Alamel model (Van Houtte et al., 2005, 2006). The deformation during ARB processing is approximated by a two-dimensional velocity gradient tensor which is composed of a compressive strain (εzz, εyy = 0 and εzz =−εxx) added with a simple shear strain (εxz). The texture simulations have been performed on the initial texture of Fig. 14 for different strain modes with various ratios of the simple shear component (εxz) to the plane strain compression components. The imposed shear during the

The texture simulation for the subsurface region has been performed with a large value of the shear coefficient (K= 2.4) due to localization of shear deformation in this region. The texture prediction for the mid-thickness part has been performed under the deformation condition with K= 0.1. It is assumed here that even in the middle layers of the sheet there will still be a small amount of frictional shear as the ARB experiment was carried out in totally dry conditions (i.e. without lubrication). Due to the specific stack-and-roll geometry of the ARB process, the subsurface texture after the first ARB pass appears in the central layer of the composite sample in the second ARB pass. Hence, in order to simulate the texture evolution in the middle layer in the ARB pass ni the Alamel model was applied on the subsurface from previous pass n(i−1) with a deformation mode that was characterized by K= 0.1. Fig. 18 presents results of the texture predictions for second, fourth, sixth and eighth ARB cycles. The calculated textures are in very good qualitative and quantitative agreement

The model calculation has produced the textures which exhibit a strong rolling β-fiber. It is important to notice, though, that the Dillamore and S components are far more important than the brass orientation, which corresponds very well to the experimentally observed texture, cf. Fig. 16. Furthermore, intensities of both S and brass orientations increase insignificantly whereas a considerable strengthening of the Dillamore orientation is observed during ARB process. Hence, it is shown here, on the basis of crystal plasticity modeling, that the sequence of rolling, cutting and stacking, which is characteristic for the ARB process eventually triggers a mechanism which leads to strengthening of the Dillamore

Rotation of different texture components during ARB cycles, i.e. shear to copper and Dillamore components in the midsection under the plane strain compression mode and copper, S and brass to the shear component in the surface layer due to the shear deformation, is the unique feature of ARBed aluminum sheets (Heason & Prangnell, 2002a; Kim et al., 2005). This can be considered as the result of introduction of surface layer to the midsection during the next cycle and it is also associated with the increase in the number of interfaces. In fact, these interfaces induce additional strain during rolling and are

second phase particles contributes to the transformation from low angle subgrain boundaries to high angle grain boundaries by increasing the misorientation between two adjacent subgrains and removes the unrefined bands from the microstructure. For a better understanding of the texture evolution, the main FCC fibers were calculated and plotted in Fig. 17. Generally, with increasing the number of cycles, the intensity of the α fiber decreased while the intensities of the β and τ fibers increased. The α fiber, running from Goss to brass in Fig. 17a, indicates how the brass component decreased during the ARB process. It is seen that during all ARB cycles the intensity of Goss component is negligible. Fig. 17b shows the intensity of the β fibers.

Fig. 17. Intensity of the FCC fibers in the ARB processed aluminum sheets: (a) β-fiber, (b) τfiber, and (c) α-fiber

It shows that in the early stages of the ARB process the intensity of the β fiber, running from copper over S to the brass orientation, is almost homogeneous, but with increasing the number of cycles the intensity of the copper component gradually grows at the expense of the brass and S components. The τ fiber is presented in Fig. 17c. It is observed that in this material the texture evolution is mainly centered around the copper and Dillamore components, and with an increasing number of cycles these components further intensify.
