**Nanostructure, Texture Evolution and Mechanical Properties of Aluminum Alloys Processed by Severe Plastic Deformation**

Abbas Akbarzadeh

*Department of Materials Science and Engineering, Sharif University of Technology, Tehran, Iran* 

**1. Introduction** 

Various research works have been conducted to replace heavy steel body constructions with lighter aluminum ones to achieve stronger energy consumption and environmental standards. The most important technical obstacle to this goal is the inferior ductility of most aluminum sheet alloys. It has been reported that control of the microstructure and the texture of materials is essential for improvement of their mechanical properties (Lee et al., 2002). Reducing the grain size of polycrystalline metallic materials to the nanosize (*d* < 100 nm, nanocrystalline) or submicron levels (100 nm< *d* <1 μm, ultra-fine grain) is an effective and relatively economic way of improving mechanical properties such as strength, toughness, or wear resistance in structural materials (Kim et al., 2006; Prangnell et al., 2001) which even can give rise to superplastic behavior under appropriate loading conditions (Pérez-Prado et al., 2004). Since it is practically difficult to reduce the grain size of many metallic materials such as aluminum alloys below 5 μm by a conventional cold working and recrystallization process, several new methods are developed to manufacture ultrafine grained (UFG) materials (Kim et al., 2006). These methods can be classified into two main groups namely bottom-up and top-down processes. In the bottom-up procedures, such as rapid solidification, vapor deposition and mechanical alloying, an ultra-fine microstructures is configured from the smallest possible constituents which are prohibited to grow into the micrometer domain (Pérez-Prado et al., 2004). In the top-down procedures, on the other hand, an existing microscale microstructure is refined to the submicrometer scale, e.g. by a process such as severe plastic deformation (SPD) (Pérez-Prado et al., 2004; Saito et al., 1999). The ancient Persian swords are the interesting examples of severe upset forging for development of fine microstructures (Sherby and Wadsworth, 2001).

By now, various SPD processes such as accumulative roll bonding (ARB) (Saito et al., 1999), cyclic extrusion compression (CEC) (Richert J. & Richert M., 1986), equal channel angular pressing (ECAP) (Valiev et al., 1991), and high pressure torsion (HPT) (Horita et al., 1996) have been proposed and successfully applied to various materials. The common feature of these techniques is that the net shape of the sample during processing is approximately constant, so that there is no geometric limitation on the applied strain (Prangnell et al., 2001). Among these processes, accumulative roll bonding has some unique features. Firstly,

Nanostructure, Texture Evolution and Mechanical Properties of

2004; Park et al., 2001; Saito et al., 1998; Tsuji et al, 2002b).

deformed materials at present (Tsuji et al., 2002a).

Aluminum Alloys Processed by Severe Plastic Deformation 427

increase in strength and hardness, more than two or three times of the values of starting materials can be achieved by ARB process in aluminum alloys and IF steel sheets (Lee et al.,

Microstructural evolution during ARB has been studied by several researchers. TEM investigation and local crystallography were used (Huang et al., 2003) to study the microstructural evolution of AA1100 up to eight cycles of ARB. Their results show that at large strains, almost a homogenous lamellar structure is formed across the thickness. This microstructure is subdivided by high angle boundaries and low angle dislocation boundaries. The spacing of dislocation boundaries decrease and their misorientations increase by accumulating strain. EBSD analysis illustrates three mechanisms of UFGs development at different levels of strains (Pirgazi et al., 2008b). It was shown that development of subgrains is the major mechanism during the first two cycles of ARB. This mechanism is followed by strain induced transition of low angle grain boundaries to high angles and formation of a thin lamellar structure at medium levels of strain. Fragmentation of these thin lamellar structures into more equi-axed grains is considered as the dominant mechanism after sixth cycle of ARB. A similar microstructural evolution during ARB of AA6061 has been reported (Park et al., 2001). In addition to fabrication of UFGed and nanostructured materials for metallic sheets, some other unique features for ARB process

are applied to different sheets to fabricate multi-layer composites (Min et al., 2006).

The texture analysis is a powerful tool to investigate the microstructural and substructural evolution of plastically deformed materials because it provides information on the fragmentation behavior of grains. When orientation contrast microscopy is used, the microstructural data can be linked with local textural information. It is commonly known that in materials produced by severe plastic deformation processes, the band contrast of the corresponding Kikuchi-lines decreases as a result of the presence of lattice defects, such as regions with high dislocation density, subgrains and grain boundaries. In the TSL® software, the contrast of the Kikuchi lines is quantified by the image quality (IQ) parameter. The IQ parameter is sensitive to a wide variety of additional material and instrumental factors which makes it almost impossible to deconvolute the IQ signal in order to convert it to univocal quantitative information on the local microstrain. Nevertheless, the EBSD technique is the only possible tool for nanoscale analysis of relatively large areas in severely

It has been reported (Chowdhury et al., 2006a) that the texture development in an AA3103 alloy during accumulative roll bonding process shows symmetry at all stages and the major components can be characterized as the Dillamore {4411} <11 11 8> component along with the S component with a scatter around the brass component. It has been also reported that in AA8011 aluminum alloy sheets processed by ARB, the deformation texture is dominated by the Dillamore component and the shear texture was developed near the surface of the sheets. This surface shear texture disappears rapidly as the surface area of the material reappears in the center of the composite sample in the next ARB pass (Kim et al, 2002, 2005). Similar results have been published by (Heason & Prangnell, 2002a). They have reported that in AA1100 alloy processed by the ARB, coarse unrefined bands can be retained even at very high strains. They have also proposed that the strong texture developed during ARB processing may lead to this inhomogeneity (Heason & Prangnell, 2002a). It has also been reported that the textures in Ti-IF steel produced by ARB display a conventional cross sectional gradient with typical shear component in the subsurface planes and plane strain compression components in the midsection (Reis & Kestens, 2005; Kolahi et al., 2009).

unlike the ECAP, CEC and HPT processes which require forming machines with large capacity and expensive dies, the ARB process can be performed by a conventional rolling mill without any special die. Secondly, in comparison to the other methods, the productivity of the ARB process is relatively high because this process implies the potential of industrial up scaling to a continuous production of UFGed metallic sheets or plates (Saito et al., 1999). In principle at least, the use of heavy deformation in metal processing, with the objective of producing metal alloys with superior properties, has a long history which may be traced back to the early metal-working of ancient China [16], the blacksmith's production of highquality Damascus steel in the Middle East [17] and the fabrication of the legendary Wootz steel in ancient India [18].

The ARB process is a repetitive procedure of cutting, stacking and roll bonding of similar sheets for a desired number of cycles, Fig. 1, which is practically limited by technological constraints such as the occurrence of edge cracks. In this process the thickness of the sheet varies between fixed limits and by repeating the procedure very high strains can be accumulated in material, and as a result significant structural refinement can be achieved (Huang et al., 2003; Saito et al., 1999; Tsuji et al., 2003a).

Fig. 1. Schematic illustration showing the procedure of the ARB process (Pirgazi et al., 2008a)

The first scientific paper on the ARB was published in 1998 and afterwards extensive studies have been conducted regarding the microstructural evolution and mechanical properties of various materials processed in different ARB cycles. Pure aluminum (Huang et al., 2003; Saito et al., 1998; Tsuji et al., 2003a), AA5083 (Saito et al., 1999), AA6061 (Lee et al., 2002; Park et al., 2001), AA8011 (Kim et al., 2002, 2005), AA3103 (Chowdhury et al., 2006a) and AA8090 (Chowdhury et al., 2006b) aluminum alloys, magnesium alloys (del Valle et al., 2005; Pérez-Prado et al., 2004), and Ti-IF steel sheets (Reis & Kestens, 2005, Reis et al., 2005; Tsuji et al., 2002a) are the most important materials which have already been successfully produced by the ARB process. The results of these investigations mainly indicate that during the first stages of ARB, ultra-fine grains with diameter less than 1 μm partially form in the sheets, and the volume fraction of these grains increases with increasing the number of cycles, so that after high levels of strain the sample is completely covered with ultra-fine lamellar grains which are not equi-axed and represent an aspect ratio bigger than one (Huang et al., 2003; Park et al., 2001; Saito et al., 1998; Tsuji et al., 2003a). The investigations on the mechanical properties by previous researchers show that the ARB is a promising process for improving this feature of metallic sheets. It has been reported that a significant

unlike the ECAP, CEC and HPT processes which require forming machines with large capacity and expensive dies, the ARB process can be performed by a conventional rolling mill without any special die. Secondly, in comparison to the other methods, the productivity of the ARB process is relatively high because this process implies the potential of industrial up scaling to a continuous production of UFGed metallic sheets or plates (Saito et al., 1999). In principle at least, the use of heavy deformation in metal processing, with the objective of producing metal alloys with superior properties, has a long history which may be traced back to the early metal-working of ancient China [16], the blacksmith's production of high-

The ARB process is a repetitive procedure of cutting, stacking and roll bonding of similar sheets for a desired number of cycles, Fig. 1, which is practically limited by technological constraints such as the occurrence of edge cracks. In this process the thickness of the sheet varies between fixed limits and by repeating the procedure very high strains can be accumulated in material, and as a result significant structural refinement can be achieved

Fig. 1. Schematic illustration showing the procedure of the ARB process (Pirgazi et al., 2008a) The first scientific paper on the ARB was published in 1998 and afterwards extensive studies have been conducted regarding the microstructural evolution and mechanical properties of various materials processed in different ARB cycles. Pure aluminum (Huang et al., 2003; Saito et al., 1998; Tsuji et al., 2003a), AA5083 (Saito et al., 1999), AA6061 (Lee et al., 2002; Park et al., 2001), AA8011 (Kim et al., 2002, 2005), AA3103 (Chowdhury et al., 2006a) and AA8090 (Chowdhury et al., 2006b) aluminum alloys, magnesium alloys (del Valle et al., 2005; Pérez-Prado et al., 2004), and Ti-IF steel sheets (Reis & Kestens, 2005, Reis et al., 2005; Tsuji et al., 2002a) are the most important materials which have already been successfully produced by the ARB process. The results of these investigations mainly indicate that during the first stages of ARB, ultra-fine grains with diameter less than 1 μm partially form in the sheets, and the volume fraction of these grains increases with increasing the number of cycles, so that after high levels of strain the sample is completely covered with ultra-fine lamellar grains which are not equi-axed and represent an aspect ratio bigger than one (Huang et al., 2003; Park et al., 2001; Saito et al., 1998; Tsuji et al., 2003a). The investigations on the mechanical properties by previous researchers show that the ARB is a promising process for improving this feature of metallic sheets. It has been reported that a significant

quality Damascus steel in the Middle East [17] and the fabrication of the legendary

Wootz steel in ancient India [18].

(Huang et al., 2003; Saito et al., 1999; Tsuji et al., 2003a).

increase in strength and hardness, more than two or three times of the values of starting materials can be achieved by ARB process in aluminum alloys and IF steel sheets (Lee et al., 2004; Park et al., 2001; Saito et al., 1998; Tsuji et al, 2002b).

Microstructural evolution during ARB has been studied by several researchers. TEM investigation and local crystallography were used (Huang et al., 2003) to study the microstructural evolution of AA1100 up to eight cycles of ARB. Their results show that at large strains, almost a homogenous lamellar structure is formed across the thickness. This microstructure is subdivided by high angle boundaries and low angle dislocation boundaries. The spacing of dislocation boundaries decrease and their misorientations increase by accumulating strain. EBSD analysis illustrates three mechanisms of UFGs development at different levels of strains (Pirgazi et al., 2008b). It was shown that development of subgrains is the major mechanism during the first two cycles of ARB. This mechanism is followed by strain induced transition of low angle grain boundaries to high angles and formation of a thin lamellar structure at medium levels of strain. Fragmentation of these thin lamellar structures into more equi-axed grains is considered as the dominant mechanism after sixth cycle of ARB. A similar microstructural evolution during ARB of AA6061 has been reported (Park et al., 2001). In addition to fabrication of UFGed and nanostructured materials for metallic sheets, some other unique features for ARB process are applied to different sheets to fabricate multi-layer composites (Min et al., 2006).

The texture analysis is a powerful tool to investigate the microstructural and substructural evolution of plastically deformed materials because it provides information on the fragmentation behavior of grains. When orientation contrast microscopy is used, the microstructural data can be linked with local textural information. It is commonly known that in materials produced by severe plastic deformation processes, the band contrast of the corresponding Kikuchi-lines decreases as a result of the presence of lattice defects, such as regions with high dislocation density, subgrains and grain boundaries. In the TSL® software, the contrast of the Kikuchi lines is quantified by the image quality (IQ) parameter. The IQ parameter is sensitive to a wide variety of additional material and instrumental factors which makes it almost impossible to deconvolute the IQ signal in order to convert it to univocal quantitative information on the local microstrain. Nevertheless, the EBSD technique is the only possible tool for nanoscale analysis of relatively large areas in severely deformed materials at present (Tsuji et al., 2002a).

It has been reported (Chowdhury et al., 2006a) that the texture development in an AA3103 alloy during accumulative roll bonding process shows symmetry at all stages and the major components can be characterized as the Dillamore {4411} <11 11 8> component along with the S component with a scatter around the brass component. It has been also reported that in AA8011 aluminum alloy sheets processed by ARB, the deformation texture is dominated by the Dillamore component and the shear texture was developed near the surface of the sheets. This surface shear texture disappears rapidly as the surface area of the material reappears in the center of the composite sample in the next ARB pass (Kim et al, 2002, 2005). Similar results have been published by (Heason & Prangnell, 2002a). They have reported that in AA1100 alloy processed by the ARB, coarse unrefined bands can be retained even at very high strains. They have also proposed that the strong texture developed during ARB processing may lead to this inhomogeneity (Heason & Prangnell, 2002a). It has also been reported that the textures in Ti-IF steel produced by ARB display a conventional cross sectional gradient with typical shear component in the subsurface planes and plane strain compression components in the midsection (Reis & Kestens, 2005; Kolahi et al., 2009).

Nanostructure, Texture Evolution and Mechanical Properties of

**2. Experimental methods** 

and post-processing procedures.

accumulated equivalent strain of 6.4 was achieved.

spinning process for manufacturing of clad tubes (Mohebbi, 2009a).

Aluminum Alloys Processed by Severe Plastic Deformation 429

bonding of two or more sandwich sheets (Zhan et al., 2006). Although it is possible to fabricate seamed composite tubes by sandwich sheets, it is ideal to plan a method with advantages of roll-bonding to produce seamless thin-walled tubes and cylinders. The spinbonding method in cold-bonding of cylinders was also proposed by utilizing the tube

Tube spinning is an incremental and localized deformation in which material between the roller and the mandrel gradually deforms to the final thickness. Actually the deformation area is limited to a part of work piece which is in contact to the roller, so that the deformation is constrained strongly by surrounding metals. This is why the stress and strain have such a complicated distribution in this incremental process. There is always an inhomogeneous material flow due to the local deformation between the roller and the preform surface with a high strain rate (Mohebbi & Akbarzadeh, 2010b, 2010c, 2009b).

The materials used to study the ARB process were fully annealed sheets of Al-Mn alloy (AA3003) and commercial purity aluminum (AA1100) with initial grain sizes of 40 and 34 µm. Thickness of the sheets was respectively 0.4 and 0.5 mm, and the chemical compositions are presented in Table 1. Two sheets of 150×50 mm2 were degreased (in acetone) and wire brushed (by a stainless steel brush with wires of 0.4 mm in diameter). After the surface treatment, the two sheets were stacked on top of each other and preheated to a temperature of 250 °C for 5 minutes. The plane strain rolling was performed along the longest dimension by 50% reduction in thickness at 200 °C without any lubrication and the mean strain rate was 51 sec-1. Afterwards the roll bonded sample was cut into two sheets of approximately the initial dimension and the procedure was repeated up to a total of 8 cycles so that an

Alloy Fe Cu Si Ti Mg Mn Al

AA1100 0.57 0.12 0.13 0.03 0.02 0.013 Bal.

AA3003 0.436 0.216 0.18 0.009 0.005 1.076 Bal.

The microtexture and microstructure measurements were performed on the section which is perpendicular to the transverse direction of rolling and is located at the mid-thickness of the sheets. The analysis was carried out by employing an Orientation Image Microscopy (OIM) attached on a Philips XL30 ESEM microscope equipped with a LaB6 filament. The electron back-scattering diffraction (EBSD) mappings were carried out with step sizes in the range of 0.12 μm to 80 nm and the OIM software developed by TSL® was used for data acquisition

Since the ARBed samples were severely deformed, acquiring good Kikuchi-patterns was not evident. By means of electro polishing a very smooth surface was produced and by adjusting the operating parameters of the SEM and the software (OIM Data Collection) a

Table 1. Chemical composition of the aluminium alloys (Pirgazi & Akbarzadeh, 2009)

It has been reported that in AA8011 aluminium alloy sheets processed by ARB, the main deformation texture orientation is the Dillamore {4 4 11}<11 11 8> component and the shear texture is developed on the surface of the sheets (Kim et al., 2002). This surface shear texture disappears rapidly as the surface area of the material reappears in the centre of the composite sample during the next ARB cycle. The role of second phase particles on the grain size reduction of aluminium sheets during the accumulative roll bonding process has been investigated by comparing the microstructure and texture of a single phase (AA1100) and a particle containing aluminium alloy (AA3003) during various ARB cycles (Pirgazi & Akbarzadeh, 2009).

Circular shaped hollow sections like tubes and cylinders, as a category of engineering components, are also expected to achieve high strengths by nanostructure. Tube spinning is used as a common tube fabricating process (Wong et al., 2003), based on which a coldbonding process titled "spin-bonding" with the advantages of ARB of sheets has been recently proposed (Mohebbi & Akbarzadeh, 2010a, 2010b) for manufacturing of high strength thin-walled tubes and cylinders. The SPD process proposed in that work is, in fact, repeatedly spin-bonding of layered tubes (accumulative spin-bonding, ASB) to induce large plastic strain on tubes similar to the ARB of sheets, Fig. 2 (Mohebbi & Akbarzadeh, 2010a).

Fig. 2. Schematic illustration of the spin-bonding (Mohebbi & Akbarzadeh, 2010a)

In many cases, the inner and outer surfaces of hollow sections are exposed to different environments, and different characteristics are required inside and outside. In these cases, various bimetallic or clad tubes of stainless steels or high-alloy steels and super alloys clad onto carbon or low alloy steels, are utilized in boilers, heat exchangers, nuclear power plants and petroleum and chemical industries (Chen et al., 2003). So far, several methods have been used to produce composite tubes and cylinders. While centrifugal casting (Sponseller et al., 1998) and extrusion (Chen et al., 2003) can be used to fabricate the thick-walled bimetallic tubes, explosive bonding (Berski et al., 2006), ball attrition (Zhan et al., 2006), thermohydraulic fit method and hydraulic expansion method (Wang et al., 2005) are capable of manufacturing thin-walled cylinders. The most widely used cladding process is rollbonding of two or more sandwich sheets (Zhan et al., 2006). Although it is possible to fabricate seamed composite tubes by sandwich sheets, it is ideal to plan a method with advantages of roll-bonding to produce seamless thin-walled tubes and cylinders. The spinbonding method in cold-bonding of cylinders was also proposed by utilizing the tube spinning process for manufacturing of clad tubes (Mohebbi, 2009a).

Tube spinning is an incremental and localized deformation in which material between the roller and the mandrel gradually deforms to the final thickness. Actually the deformation area is limited to a part of work piece which is in contact to the roller, so that the deformation is constrained strongly by surrounding metals. This is why the stress and strain have such a complicated distribution in this incremental process. There is always an inhomogeneous material flow due to the local deformation between the roller and the preform surface with a high strain rate (Mohebbi & Akbarzadeh, 2010b, 2010c, 2009b).
