Microstructure, texture and interface integrity in sheets processed by Asymmetric Accumulative Roll-Bonding

Materials Science & Engineering A xxx (xxxx) xxx

Contents lists available at ScienceDirect

Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea

Microstructure, texture and interface integrity in sheets processed by Asymmetric Accumulative Roll-Bonding ~es a, *, Martina Avalos b, Renan Pereira de Godoi a, Danielle Cristina Camilo Magalha b a Raul Eduardo Bolmaro , Vitor Luiz Sordi , Andrea Madeira Kliauga a a b

Department of Materials Engineering, Federal University of S~ ao Carlos (UFSCar), 13565-905, S~ ao Carlos, Brazil Institute of Physics Rosario (IFIR) FCEIA-UNR-CONICET, Bv. 27 de Febrero 210 Bis, S2000EZP, Rosario, Argentina

A R T I C L E I N F O

A B S T R A C T

Keywords: Aluminum alloy Asymmetric accumulative roll-bonding (AARB) Diffusion bonding Microstructure Texture Mechanical properties

Accumulative Roll-Bonding (ARB) and Asymmetric Rolling (AR) techniques were combined to produce ultrafinegrained aluminum sheets with the mechanical characteristics of a Severe Plastic Deformation (SPD) process. Temperature and number of bonding cycles were varied to promote grain refinement, texture randomization and high-quality sheet bonding. Finite element simulation for a single pass was performed to clarify the strain dis­ tribution differences between symmetric and asymmetric roll -bonding. The microstructure and crystallographic texture were measured by Electron Backscatter Diffraction (EBSD) and X-ray diffraction. Hardness and tensile tests characterized strain distribution and bonding efficiency. A fine grain structure with a mean grain size of 1.0 μm was achieved at 350 � C, whereas a coarser grain structure was obtained at 400 � C. The grain size and shape distribution were linked to enhancing the mechanical strength in a transversal direction. During repeated bonding cycles at both temperatures, extra shear in the interfacial region yielded favorable homogeneous strain distribution and a weak shear texture across the sheet. Rotated-cube orientation was the strongest component in both processing temperatures. To increase the interfacial strength, mainly on the last bond interface, an extra 50% reduction step was added. This improved the adhesion in the last bonding interface, and thus enhanced the ductility. These findings helped to provide a basis for determining the processing conditions for aluminum alloys.

1. Introduction Intense grain refinement produced by Severe Plastic Deformation (SPD) processes has recently attracted growing interest. Accumulative Roll-Bonding (ARB) is one of the SPD processes, which successfully produces alloy sheets and metal matrix composites with submicrometric or nanometric grain sizes, thus increasing yield strength, hardness, fatigue resistance and strain rate sensitivity [1–7]. An addi­ tional advantage of ARB is the potential it has for continuously pro­ ducing large bulky materials using existing industrial structures [8]. This process consists of repeatedly roll-bonding two or more plates at a 50% minimum reduction ratio, as determined by Li et al. [9]. The bond is achieved from frictional forces acting upon the interface and may be enhanced by diffusion. Thus, intermediate annealing treatments and warm or hot rolling within a temperature range of 150–500 � C are normally applied in the ARB process for aluminum alloys to improve interfacial adhesion [9–11]. However, Chekhonin et al. [12] reported

that sheets can be produced with high interfacial bonding quality of high and technical purity aluminum at room temperature, using up to 10 cycles, without any intermediate annealing steps. ARB has also been used to produce laminate metallic composites and it is applicable for processing several metals and alloys, including those that cannot be bonded by traditional bonding methods [13]. To achieve adhesion and maintain workability of heat treatable aluminum alloys (AA6061, AA7075, AA2219) bonded with non-heat treatable (AA1050, AA5086) ones [14–16], these multilayered composites must be pro­ cessed above or near the solubilization temperature of the heat treatable material. Particularly for high Stacking Fault Energy (SFE) materials, such as aluminum and its alloys, it is also well established that ultrafine-grained (UFG) structures may be produced by warm to hot deformation [17]. Therefore, it is important to know the microstructural evolution of the different layers at higher temperatures. The preferential orientation of grains, i.e., crystallographic texture, plays a crucial role in the final properties of metals and alloys. The ARB

* Corresponding author. E-mail address: [email protected] (D.C. Camilo Magalh~ aes). https://doi.org/10.1016/j.msea.2019.138634 Received 26 July 2019; Received in revised form 3 November 2019; Accepted 5 November 2019 Available online 9 November 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.

Please cite this article as: Renan Pereira de Godoi, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2019.138634

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

process, similar to conventional rolling, has a conservative deformation path, which results in pancake-like grain structures with high texture indices, regardless of the type of alloy processed [18,19]. ARB also in­ troduces some redundant shear near the interface regions, which con­ tributes to the grain orientations being scattered. However, the most notable grain orientations are the traditional rolling texture components [20,21]. An alternative modification of this feature to achieve more homogeneous properties is the association with the Asymmetric Rolling (AR) technique. During AR, the circumferential velocities of the top and bottom rolls are different. Additionally, considerable shear and rigid body rotation are introduced, leading to more random orientations. The AR technique was applied to improve the press formability of aluminum alloys [22–24]. One of the key features of the AR method is the impo­ sition of a high friction coefficient between rolls and sheet surfaces. Without sufficient friction, compressive strains dominate the process [25,26]. Combining the AR and ARB techniques creates the Asymmetric Accumulative Roll-Bonding (AARB) technique, which was first described by Wang and Shi [27] in an attempt to achieve an ultrafine grain structure in pure copper. Roll asymmetry applied to the ARB technique collaborate to improve adhesion because the presence of shear at the layer interfaces assist in the emergence of surface oxides. A critical success factor for sheets produced using AARB is interface integrity, which is evaluated by the bonded area, mechanical strength and various other methods detailed in earlier literature [9,28]. In fact, the same disadvantages that apply to ARB are also common for the AARB process: the existence of poorly bonded regions between layers and the need to establish the process conditions to minimize these imperfect areas. In addition, another disadvantage is that AARB requires changes in the configuration of the conventional rolling mill. In recent years, mechanisms for grain refinement and the mechanical behavior of sheets after AARB processing have been studied for different metals and alloys, focusing mainly on pure copper [29–31] and Al-based alloys [32–34]. Until now, there has not been a detailed investigation into crystallographic texture and the enhancement of mechanical strength after AARB for commercially pure aluminum. Taking into account the existing literature on AR [22–26], the AARB process should be more effective for both grain refinement and texture randomization, i.e., to decrease the texture intensities by introducing shear strains. In this study, experimental conditions that promote grain refinement and high-quality bonding of commercially pure aluminum (AA1050) were evaluated using the AARB process at 350 � C and 400 � C. The texture and microstructure of the Al alloy were analyzed by means of X-ray diffraction and Electron Backscatter Diffraction (EBSD) to better understand the role of AARB on interface integrity, grain refinement and texture components. Tensile tests were performed to evaluate strength and ductility in the rolling and transverse directions. The mechanisms controlling the evolution of the microstructures, texture and interface formation during AARB processing are discussed.

high friction coefficient. First, AR was performed to reduce the initial thickness from 7 mm to 1 mm. Additionally, the material was annealed at 350 � C for 2 h. Before the AARB process, the samples were wirebrushed to produce virgin surfaces, which were then rinsed and degreased with ethanol. The sheet was folded in half so that the roughened surfaces faced each other. The folded stacks were pre-heated for 5 min at 350 � C or 400 � C, then AR was performed without a lubri­ cant using a 50% thickness reduction (equivalent strain equal to 0.9) in a single pass, thus producing a primary stack. The mill was kept at room temperature and the deformation was performed with a cooling gradient. After each cycle, the sheet was quenched in water in order to control the temperature profile in the AARB cycles. The roll-bonded sheet was wire-brushed, cleaned and folded again to produce a new stack, and the process was repeated for up to 10 cycles. The rolling di­ rection of the sheets was maintained throughout all the rolling passes. The process is described schematically in Fig. 1. A second AARB set was produced at 350 � C with an initial sheet thickness of 2 mm, maintaining the same reduction rate. For these samples, an additional AR pass, using a 50% thickness reduction was performed in order to improve the interfacial quality. The final thickness of both sets was 1 mm. Table 2 summarizes these experimental conditions. 2.2. Microstructural characterization After AARB, microstructural characterization was performed by conventional optical metallography on samples anodized with Barker’s reagent (1.5% HBF4 acid in aqueous solution) at 20 V for 5 min and observed under polarized light. At the LNNano Synchrotron National Laboratory in Campinas, Brazil, the crystallographic texture of the AARB sheets was determined through X-ray diffraction on a Philips X’Pert MPD diffractometer using Co-Kα radiation with the Schulz method. The Orientation Distribution Func­ tions (ODFs) were calculated from the three incomplete pole figures adopting the Series Expansion Method from Bunge [35] using JTex software [36]. The pole figures of all samples were measured at the center and at the sub-surface of the sheets. At the IFIR – CONICET in Argentina, mesotexture was measured by means of EBSD using a FEI Quanta 200 FEG-SEM equipped with a detector from EDAX/TSL. The EBSD scans were taken on the transverse plane (RD–ND plane). Before acquiring the texture and EBSD measurements, the sample’s surface was electro-polished using a polishing solution with 160 mL distilled water, 800 mL ethanol and 60 mL perchloric acid at 10 � C and using a voltage of 25 V for 20 s, which provided a current density of 1.5–2.0 A/cm2.

2. Experimental procedure 2.1. Material and AARB processing Table 1 describes the chemical composition of the presently studied AA1050 Al alloy, which was determined by optical emission spectroscopy. The deformation process was carried out using a rolling mill, which operated at 24 rpm. The roll diameters were modified to produce a roll diameter ratio of 1.5 and the surfaces were grooved in order to ensure a Table 1 AA1050 Al chemical composition (wt.%). Al

Si

Fe

Cu

Mn

Mg

Zn

Ti

Balance

0.081

0.185

0.012

0.003

0.003

0.002

0.012

Fig. 1. Schematic illustration of the first AARB cycle. ND ¼ normal direction, RD ¼ rolling direction and TD ¼ transversal direction. 2

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

Table 2 Summary of processing conditions and specimen nomenclatures. Specimen

Processing conditions

350_4 350_6 350_10 350_6 þ 1 350_10 þ 1 400_4 400_6 400_10

AARB AARB AARB AARB AARB AARB AARB AARB

at 350 � C, at 350 � C, at 350 � C, at 350 � C, at 350 � C, at 400 � C, at 400 � C, at 400 � C,

4 cycles 6 cycles 10 cycles 6 cycles, plus one additional AR pass 10 cycles, plus one additional AR pass 4 cycles 6 cycles 10 cycles

2.3. Finite element (FE) simulation of the AARB process The FE simulation was carried out using the DEFORM 3D™ software package, which deals with large strain problems inherent to most in­ dustrial processes. The same geometrical parameters of the experimental rolling mill were used in the numerical simulations. The initial dimen­ sion of the sheet was 20 � 35 mm with variable thicknesses; initially a symmetry restriction in the middle RD x ND plane was imposed. The thickness was subdivided into eight element nodes. Sheets were considered as elastoplastic and the rolls as perfect rigid bodies, respec­ tively. The material behavior was described by the interpolation of stress-strain curves from the software databank. A triaxial deformation model and a shear friction coefficient of 0.9 were applied at the rollsheet interface at the AARB and 0.45 for the ARB since the last pro­ cess is usually performed in a lubricated condition. The shear friction coefficient at the central interface was also assumed to be 0.9. From the simulations, discrete values for shear strain εxz and compression strain εzz were calculated. 2.4. Mechanical properties and fracture surface analysis

Fig. 2. Microstructural characterization of homogenized material: (a) Scanning electron image in backscattered mode showing intermetallic particles; and (b) optical image in polarized mode showing the coarse grain structure.

Vickers microhardness profile measurements were carried out to evaluate the strain distribution across the thickness of the sheets using a Future Tech microhardness tester, under a load of 100 gf and maintained for 15 s. Hot compression tests were carried out at 350 � C and 400 � C in an INSTRON testing machine, both with an initial strain rate of 1.0 � 10 3 s 1 on specimens with a height of 9 mm and a diameter of 6 mm. Tensile tests were performed on sub-size samples with a gauge length of 7 mm and a cross section of 3 � 2 mm2. Elongation was monitored with an optical extensometer during the tests. All tensile tests were performed at room temperature with an initial strain rate of 1.0 � 10 3 s 1. After the tensile test, a FEI-Inspect S50 scanning electron microscope was used to characterize the fracture surfaces to investigate the fracture behavior of the roll-bonded sheets. 3. Results 3.1. Microstructure and mechanical behavior of AA1050 prior to AARB Fig. 2 presents the initial microstructure of the AA1050 Al. The backscattered electron image in Fig. 2-a highlights the presence of Fe and Si rich intermetallic particles that are common in this alloy grade (respective EDX analysis is shown in the inset table). The volume frac­ tion of these dispersoids was estimated to be equal to 0.3%. The optical microscope image in Fig. 2-b shows the coarse grain structure with coarse grains with a mean size of 85.0 � 10.0 μm. To provide a better understanding of the mechanisms that control the microstructural evolution during AARB, hot compression tests were performed. Fig. 3 shows the typical stress-strain curves obtained during the uniaxial hot compression tests at 350 � C and 400 � C of the AA1050 Al alloy. The hot compression tests show a softening behavior of the AA1050 Al alloy at the process temperatures. The true stress-true strain curves present a peak and saturation stress. At 300 � C, these stress values were 35 MPa and 30 MPa, respectively. In addition, at 450 � C the

Fig. 3. True stress-strain curves from uniaxial hot compression tests performed at 1 � 10 3 s 1, and temperatures of 350 � C and 400 � C.

saturation stress was 25 MPa and 20 MPa, respectively. The strain at the peak stress for both temperatures was approximately 0.2. Each AARB cycle applied a strain of 0.9 and should, therefore, promote either dy­ namic recrystallization (DRX) or continuous dynamic recrystallization (cDRX). 3.2. Strain distribution and microstructures after AARB Fig. 4 presents the microstructures of the AA1050 aluminum alloy after AARB processing and shows the evolution of the adherence quality as the number of cycles increased. It can also be seen that at 350 � C, a 3

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

Fig. 4. Microstructures of the AA1050 after AARB in different conditions: (a) 350_4, (b) 350_6, (c) 400_4, and (d) 400_6. Arrows indicate reminiscent interface regions. ND ¼ normal direction and RD ¼ rolling direction.

much finer microstructure was obtained compared to the samples pro­ cessed at 400 � C. Furthermore, more homogeneous distribution of grain size across the sample thickness is perceived at 350 � C, compared to 400 � C. Hardness profiles for the AA1050 alloy before and after AARB are shown in Fig. 5. The hardness measurements were performed to deter­ mine strain distribution along the AARB sheets (Fig. 5-a) and to evaluate the influence of the intermediate annealing (Fig. 5-b) in the samples processed at 350 � C. Fig. 5-a provides evidence of the influence that the temperature has on the microstructures shown in Fig. 4. At 350 � C, the hardness saturated at 62 HV and in the 400 � C samples, the hardness saturated at 54 HV. After the first cycle, the sheet center hardness was smaller than at the roll/sheet interfaces, but the mean hardness and strain homogeneity increased in correlation with the increasing number of cycles. During the intermediate annealing (see Fig. 5-b), the hardness of the samples processed at 350 � C decreased from 64 HV to 40 HV, indicating a recovery process. In view of these hardness profiles, it is evident that after ten cycles there was a better homogenization of me­ chanical properties across sample thicknesses. Deformation was not accumulative, but rather consisted of deformation and recovery cyclic steps. The saturation level of the hardness at 350 � C was similar to the one obtained by Equal-Channel Angular Pressing (ECAP) and ECAP plus rolling in the same material at room temperature [37]. Fig. 6 presents the finite element simulation results for AARB at 350 � C and 400 � C and for the ARB process at 400 � C. For both processes, the compression component is higher than the shear component. The calculated compressive strains had a constant value of 0.7 for ARB and varied from 0.7 at the upper roll to 0.6 to the bottom roll in the AARB process. The major difference is in the shear strain induced by the asymmetry, as shown in Fig. 6-b. For the ARB process, a shear of low intensity (0.1) is predicted close to the sheet/roll interface. For the AARB process, a strong shear gradient is predicted from the top (0.35) to the bottom roll (0.1). The friction between the layers and temperature significantly influence the shear transmission from the top to the bottom layer: at 400 � C, a stronger discontinuity was predicted at the interfacial region, whereas a more continuous gradient was predicted at 350 � C.

Fig. 5. Hardness profiles of the samples: (a) processed at 350 � C and 400 � C; (b) after intermediate annealing at 350 � C (letter A following the sample label means intermediate annealing).

4

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

the center and surface of the sheets. The typical texture components are listed in Table 4. Fig. 9 shows that the crystallographic orientation distributions are very similar, after 10 cycles for both temperatures, with an indication of {111} plane positioning on the surface of the sheets and the formation of rotated-cube orientations (100)[011], which were already illustrated in Figs. 7 and 8. There was no significant variation of component intensity between the center and the surface. The relative intensities of the shear and the rolling components are shown in Fig. 9-e. Due to asymmetry, the AARB generated shear textures with a weaker intensity than the com­ ponents of a conventional rolling technique with 50% thickness reduc­ tion, which were also plotted for the sake of comparison. The distributions of cube/rotated-cube and γ-fiber texture (B, Bs) components on the microstructure of 350_4 and 400_ 10 sheets in the analyzed region are presented in Fig. 10, because these are the most characteristic orientations yielded by shear. In sample 350_4, the rotated-cube component had a more intense orientation than the other samples and this orientation is located at the interface regions, while the {111}//ND orientations are dispersed in the bulk and follow the shear direction. Sample 400_10 had a much larger grain size than the layer thickness and both orientations are well distributed along the sheet thickness, keeping in mind that the {111}//ND orientations are asso­ ciated with shear bands. The cube/rotated-cube texture components of sample 350_4 (see Fig. 10-a), in red, occupy about 19% of the analyzed area, while the γ-fiber texture component (green areas in Fig. 10-b) covers approximately 16%. It was observed that while the temperature increases, the cube and rotated-cube texture components exert more influence on the sheet texture, as shown in Fig. 10-c, which occupy 20% of the analyzed area. However, this value is lower than the 350_4 sample, indicating that AR is capable of decreasing the cube texture components in aluminum. In addition, the γ-fiber texture component for the 400_10 sample covers up to 15% of the analyzed area, as presented in Fig. 10-d. Thus, the results of the current study affirm that the AARB process produces a typical texture of shear strain, combined with intense grain refinement and more random orientation.

Fig. 6. Results from the FE simulation: (a) equivalent strain map for the AARB process at 350 � C; and (b) calculated compression (εzz) and shear (εxz) strain components along the sheet thickness for AARB and ARB.

This shows that the interfacial friction will greatly influence the shear strain distribution of the upper and bottom layers. Because of interme­ diate annealing steps, it is reasonable to assume that these strain dis­ tributions were repeated at each AARB cycle. EBSD maps were used to obtain more detailed information of orientation and the spatial distribution of the grains. Orientation maps and grain boundary maps for 4, 6, and 10 cycles at 350 � C and 400 � C are presented in Figs. 7 and 8. Grain size, proportion of Low-Angle Grain Boundaries (LAGBs) and High-Angle Grain Boundaries (HAGBs) are summarized in Table 3. At 350 � C, the grains are finer and the structure is characterized by a high fraction of HAGBs. The average grain size varies in the range of 1–2 μm, which is in the range of the minimum grain size obtained by SPD in ECAP or ECAP followed by rolling at room temperature [37]. After 10 cycles the layer thickness is smaller than the actual grain size, which shows that it is grain growth that has promoted the effective bonding. In sample 350_10 (see Fig. 7-e), regions close to the interme­ tallic particles show evidence of particle stimulated nucleation. The orientation maps and pole figures at both temperatures (Fig. 7-a, 7-c, 7-e, 8-a, 8-c and 8-e) show that a rotation of the conventional rolling texture occurred at positioning {111} planes parallel to the sheet sur­ face, and also that the rotated-cube (100)[011] component is one of the major components in these samples, which are common to shear deformation with a rotation axis parallel to the transverse direction. At 400 � C, a coarser grain structure was formed, the larger grains have a thickness of approximately 15 μm, which is the same magnitude of layer thickness for 6 cycles, showing that the layers were consolidated during the repeated annealing steps. The mean grain size is 8 μm for all samples. A fine grain structure is observed around the HAGBs (see Fig. 8e and 8-f), indicating the occurrence of partial dynamic recrystallization at this temperature. At the layer interface, a formation process of new grains can also be seen in Fig. 8-c and 8-d. This indicates that shear acting between the layers at the interface can induce nucleation of new grains, which is one of the likely bonding mechanisms.

3.4. Mechanical response after AARB and fracture surface analysis The average values of main tensile properties of the sheets processed by AARB at 350 � C and 400 � C are summarized in Table 5. As expected, the mechanical strength after 10 cycles decreases and elongation is enhanced, which corroborates with average grain size distribution observed in these samples. For both temperatures, 6 cycles were the condition that increased the sheet’s mechanical strength. After different amounts of AARB cycles on the AA1050, the interface integrity improves, as shown in Fig. 11. With increasing accumulated strain, interface bonding becomes stronger and decreases the relative areas of debonding on the fractured surface. It also improves elongation. Fig. 11-c and 11-d show that the temperature enhances the sheet bonding, especially when comparing the 4-cycle samples in Fig. 11-a and 11-d, which promote a microstructural homogeneity along the sheet’s thickness. The 400 � C route promotes better bonding than the 350 � C route, however softening and grain growth can be observed after 10 cycles. This consequent softening behavior shows that control over the intermediate annealing time and temperature is a critical parameter in the AARB process. Additionally, the imperfections related to the last AARB cycle are present. The weakest bonded area has always been the last interface. In general, roll-bonding of metals and alloys is affected by several factors, including the amount of deformation, rolling temperature, bonding time, surface conditions, among others. In the present study, one extra AR pass was performed without sheet bonding, while maintaining the 50% thickness reduction in order to improve adhesion between the sheets. The 350_6 and 350_10 sample sheets were selected because they showed a better balance between strength and ductility. The results of the fracture surface analysis after the extra AR pass are shown in Fig. 12.

3.3. Deformation textures To better understand the microstructural evolution after AARB, the texture was evaluated through X-ray diffraction measurement, which is appropriate to describe the orientation at a higher volume of material on 5

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

Fig. 7. EBSD Results from samples processed at 350 � C. Orientation maps, and respective {100} and {111} pole figures for A1 ¼ RD and A2 ¼ TD: (a) 4 cycles; (c) 6 cycles and (e) 10 cycles. Grain boundary maps showing HAGBs (in black) and LAGBs (in red): (b) 4 cycles, (d) 6 cycles and (f) 10 cycles. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

6

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

Fig. 8. EBSD results from samples processed at 400 � C. Orientation maps and respective {100} and {111} pole figures for A1 ¼ RD and A2 ¼ TD: (a) 4 cycles; (c) 6 cycles and (e) 10 cycles. Respective grain boundary maps showing HAGBs (in black) and LAGBs (in red): (b) 4 cycles, (d) 6 cycles and (f) 10 cycles. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

7

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

sheets produced by AARB were investigated in detail. For the AA1050 Al alloy, the hardness saturation value after SPD at room temperature reported in recent studies was approximately 55–80 HV [37]. This stress saturation is obtained at equivalent strain levels greater than four when a grain size saturation of approximately 1.0 μm is reached. In the present work, this level of structure refinement, with a high area fraction of 70% HAGBs, was achieved after a nominal strain of two at 350 � C. At this temperature, the AARB cycle led to a saturation of equiaxed grains of about 1.5 μm, as shown in Table 3, similar to that obtained through the ECAP process at room temperature [37,39]. It is worth noting that the asymmetry of the rolls clearly enhances the grain subdivision by shear bands and the rotational field in the process ac­ celerates the HAGB formation. Each intermediate annealing induced recovery and grain growth, which reduced hardness from 60 to 40 HV, as demonstrated in Fig. 4-b. A hardening/softening cycle was estab­ lished and the grain saturation size was independent of the nominal applied strain. Furthermore, a high volumetric fraction of 70% HAGBs was observed and new fine grains were found at the interfaces between layers and around intermetallic particles. The fluctuation of average grain size (see Table 3) shows that the actual amount of accumulated strain depends on the control of this interactive thermo-mechanical process. Recrystallization is a key phenomenon on the microstructural evo­ lution control, and thus contributes to optimizing the mechanical strength. The cDRX in high SFE materials, such as aluminum and its alloys, is promoted through continuous accumulation of dislocations at sub-grain boundaries and through gradual rotation until they turn into HABGs, and this mechanism is also predominant in any temperature range [38]. In contrast, the discontinuous dynamic recrystallization (dDRX) will happen at low strain rates and high temperatures, as re­ ported by Sakai et al. [17]. In Fig. 3, the softening of the AA1050 Al alloy in hot compression can be observed, and Figs. 7 and 8 show evidence of cDRX at both testing temperatures. During the AARB process, the microstructure changes due to compression and shear deformation and cDRX. The fluctuation of the average grain size (see Table 3) shows that the actual amount of accumulated strain depends on the control of this interactive thermo-mechanical process. This evidence supports the idea reported in previous studies [17,40], in which cDRX should occur in AARB at the applied temperatures rather than from the accumulative deformation. In addition, it is clear that the adhesion between interfaces improves with an increasing number of cycles. Therefore, optimal pro­ cessing conditions should balance these characteristics. It is well-established that the amount of deformation required to achieve a homogeneous fine-grained structure increases with the tem­ perature [17]. At 400 � C, due to the necklacing mechanism in the AARB, there is a large deviation of measurements and the average grain size is approximately constant regardless of the number of cycles (see Table 3). A more developed necklace substructure was observed after 10 cycles, as can be seen in Fig. 8. This reveals that the cDRX was incipient. The resulting distribution of a high volumetric fraction of LAGBs in the grain interior and a necklace of fine grain structures closer to the grain boundaries with HAGBs was due to the greater accumulation of misorientation close to the parent grain boundaries. This misorientation is enhanced by the shear strain and the rigid body rotation imposed by the AR process. It is interesting to note that new fine grains were found at the in­ terfaces between layers around intermetallic particles, as shown in Fig. 7. Particle stimulated nucleation in aluminum alloys is observed at room temperature during SPD [37,41] and in warm ARB [42]. The presence of particles yields larger lattice rotation and the rapid refine­ ment of grains in their surroundings. When the volume fraction is as high as 3–5%, the affected volume will be large enough to reduce the texture intensity [42]. In the present work, the volume fraction of intermetallic particles is low (~0.3%) and the effect on the texture is not relevant. Nevertheless, in Fig. 7-e and 7-f, it is shown that they

Table 3 Theoretical layer thickness compared to grain size: LAGB and HAGB fractions obtained from the EBSD measurements. Specimen 350_4 350_6 350_10 400_4 400_6 400_10

Layer thickness (μm) 62.50 15.62 0.97 62.50 15.62 0.97

Grain size (μm) 0.8 � 0.4 1.1 � 0.5 2.4 � 0.8 8.8 � 6.0 8.5 � 6.1 8.3 � 6.0

HAGBs (%) 77 89 69 37 59 53

LAGBs (%) 23 11 31 63 41 47

Table 4 Main ideal orientations along the simple shear and usual rolling orientations given by the Miller indices {hkl}. Notation

Texture component

RC (rotatedcube) A, A1, A2 B, Bb C A, Ab Cube Copper Goss (Gx) S1, S2, S3

{100}<110>

Shear Components

{112}<211>,{112}<110> {111}<110>, {111}<112> {011}<011>, {011}<011> {110}<112> {100}<001> {112}<111> {110}<001> {124}<211>, {123}<412>,{123}< 634> {110}<112>

Rolling Components

Brass (Bs)

While comparing the surface fractures before the extra AR pass (see Fig. 11), it can be seen that the debonded fraction area has decreased. The tensile properties were also measured in the Rolling (RD) and Transverse Directions (TD), as demonstrated in Fig. 13. At 350 � C, there were no significant property differences (yield, ultimate tensile strength and elongation). However, at 400 � C, a difference was perceived when the tensile axis changed. Due to the texture similarities in these samples, this difference was not caused by texture orientations, but rather by the different grain size perceptions in the RD and TD. Both yield and ulti­ mate tensile stresses slightly decreased for the 350_6 þ 1 sample, compared to 350_6, as presented in Fig. 13. Moreover, the extra AR pass improved the elongation for both the 350_6 þ 1 and 350_10 þ 1 samples. A more detailed investigation of interface adhesion was performed using orientation mapping in the 350_6 þ 1 and 350_10 þ 1 samples, presented in Fig. 14. The fine grain distribution was maintained and total interface union was achieved. 4. Discussion 4.1. Microstructural and texture development in AARB processing The AARB technique was applied to the AA1050 Al alloy as a pro­ cessing route to obtain a weak crystallographic texture with a finegrained microstructure. The first studied parameter was the integrity of the bonded interface, which was affected by several factors, such as the rolling temperature, equivalent strain and annealing time. In the present investigation, the temperature range was chosen in order to improve the bonding mechanisms. In fact, sheets can be produced by ARB at lower temperatures than in this work. However, solid-state bonding is a consequence of both mechanical and diffusion processes. Mechanical bonding is achieved in a very short period during asym­ metric rolling, while diffusion bonding involves temperature and pres­ sure and it requires a considerable time. Hence, it could conceivably be hypothesized that although the process is driven by mechanical bonding, warm or hot rolling may improve the interfacial adhesion as diffusion is enhanced at this temperature range. In addition, the resul­ tant crystallographic texture, microstructural development, and the mechanical response in the longitudinal and transversal directions of the 8

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

Fig. 9. {100} and {111} pole figures for samples: (a) 350_10 at the sheet surface and (b) center of the same samples; (c) 400_10 at sheet surface and (d) center of the same sample; (e) intensity of texture components for the 350_10 and 400_10 samples, in addition to the conventional rolled sample with a 50% thickness reduction (R50%). The letters C and S mean center and surface, respectively.

contribute to a localized grain refinement and texture orientation spread. The orientation maps also show that the distribution of orientations is very different from the ones commonly obtained in symmetric rolling. ARB at room temperature or at higher temperatures elongated grains with similar orientation and high texture intensity is the most charac­ teristic feature [43]. On the other hand, for the AARB results presented here, the spatial distribution of the texture components typical to that of shear strain, γ-fiber {111}//ND and rotated-cube as seen in Fig. 10, are more randomly distributed. In the AR process, the shear component is mostly concentrated close to the rolling surfaces, whereas in the sheet center, the rigid body rotation was predominate [44–48]. The subdivi­ sion of grains leads to scattering around the ideal texture components, which are considerably evident in the present study. This is the main effect of the shear strain introduced by the asymmetry. Fig. 6 shows that a shear gradient forms from the upper to the bottom roll and that the sliding at the interface may decrease the shear

transmission. Fig. 8-d shows that at the middle interface new grain nucleation is stimulated by these increments of localized sliding. Moreover, Fig. 9 shows that at the center, the shear texture is as intense as on the upper surface. Thus, the numerical simulation and the microstructural analysis indicate that the asymmetry enhances oxide film break out due to friction, and grain nucleation and growth at the interface, which are the main bonding mechanisms. Fig. 9 clearly shows a decrease in the texture components relative to the conventional rolling process (S1, S2, S3, Cu, Bs, G), and the modi­ fication of the orientation towards a {111}//ND fiber (γ-fiber) and rotated-cube. This is due to the shear component being oriented in the TD direction and is induced among the roll/sample and at the interfaces of the layers during the process. The increase of rotated-cube texture components, commonly observed after AR processing can be explained by the displacement of the copper (112)[111] and cube components [49] due to the shear. Fig. 7 shows the EBSD scan of a large area in the 350_4 sample and 9

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

Fig. 10. Texture component distributions for the sheets processed by AARB at 350 � C, 4 cycles for 350_4: (a) cube/rotated-cube, (b) γ-fiber; and for 400_10: (c) cube/ rotated-cube and (d) γ-fiber.

process yielded an intense grain refinement at both tested temperatures. At 400 � C, the grains are elongated along RD, but have a less oriented distribution.

Table 5 Average values of main tensile properties of the AA1050 Al alloy sheets pro­ cessed by AARB. Specimen

Yield stress (MPa)

Ultimate tensile stress (MPa)

Total elongation (%)

350_4 350_6 350_10 400_4 400_6 400_10

140 � 8 157 � 2 133 � 7 134 � 3 138 � 4 106 � 15

151 � 5 169 � 7 152 � 7 151 � 5 149 � 3 130 � 11

7.5 � 5.0 5.0 � 2.0 7.0 � 4.0 11.0 � 5.0 8.0 � 2.0 13.0 � 5.0

4.2. Tensile behavior of the sheets processed by AARB and fracture surface analysis To better understand the observed mechanical behavior, an inves­ tigation of fracture mechanisms at the interface is useful. For all speci­ mens tested, a ductile fracture was observed. At 350 � C, delamination indicates that even though a good cohesion was achieved, the area fraction of non-adherent areas was higher (see Fig. 11-a). After 10 AARB cycles, delamination did not occur throughout the entire sheet due to the presence of numerous well-bonded points that did not allow for the propagation of delamination cracks, as illustrated in Fig. 11-b. Another possible explanation for the lack of delamination is the layer thickness at 10 cycles, which was close to the grain size limit for this material. The grain size limit is about 1.5 μm (see Table 3) and, at that point, no layer boundary was observed.

provides additional evidence that the rotated-cube orientation is asso­ ciated with the shear strain at the interface of the layers. The rotated-cube and Dillamore orientations have also previously been observed in ARB processed aluminum due to interfacial sliding [18,42]. Thus, it is determined that the AARB process can be applied to reach a combination between UFG structure, a more random orientation and a typical crystallographic orientation of the shear. The above reported 10

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

Fig. 11. Fracture surfaces after the tensile tests of AARB performed at: (a, b) 350 � C, and (c, d) 400 � C.

Fig. 12. Fracture surfaces of the samples with an extra AR pass: (a) 350_6 þ 1 and (b) 350_10 þ 1.

The most obvious finding from the surface fracture analysis was the reduction in delamination among sheets processed at 400 � C, as seen in Fig. 11-c and 11-d. This result can be explained by the fact that diffusion mechanisms are also activated, which increase diffusion bonding as well as recrystallization and grain growth processes into neighboring layers

Fig. 13. Tensile results in the rolling and transverse directions of the AARB samples: (a) Yield Stress (YS) and Ultimate Tensile Stress (UTS), and (b) total elongation.

11

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

mainly at the interface zone, and promote the work softening behavior at higher strain levels. The results of the present study also indicated that there is a slight increase in grain size with the increasing strain by AARB. In Table 3, it can be observed that 350_6 and 350_10 samples present an average grain size of 1.1 and 2.4, respectively. These results are a strong indicative of the occurrence of the softening mechanisms, namely cDRX and static and dynamic recovery, are in good agreement with the literature [17, 54]. Hence, the strength of the sheets did not necessarily enhance with increasing strain after several cycles of AARB; the final strength is dictated by a combination among the average grain size, a fraction of HAGBs and the number of AARB cycles. Thus, a minimum of four AARB cycles at 350 � C is needed to obtain an equiaxed and ultrafine-grained structure in AA1050 Al alloy, with a relatively high fraction of the HAGBs. In addition, six cycles represent essentially an upper limit to avoid any problems associated with cDRX, static and/or dynamic re­ covery and partial grain growth in the sheets, with the highest strength level. At higher temperature ranges, it can be observed for the 400_6 sample that the grain size is close to 8.5 � 6.1 μm and the HAGBs covers 59% of the area fraction. In comparison to 400_10, the grain sizes are very similar, but the HAGB fraction decreased to 53%. It can be related also to the smaller strength level after ten cycles at 400 � C compared to the six cycles at same processing temperature. Once more, the cDRX, static and/or dynamic recovery and partial grain growth are responsible for the mechanical behavior observed. Obviously, the magnitude and kinetics of these phenomena are greater due to the temperature range compared to processing at 350 � C. In addition, an unexpected outcome for the samples submitted up to 10 AARB cycles can be seen in Table 3: the average grain sizes are larger than the theoretical layer thickness. From the initial sheets with 1 mm of thickness, after 10 cycles it is expected that this value will achieve approximately 0.97 μm, estimated by the expression 1/2N, where N is the number of cycles. Surprisingly, the average grain size in 350_10 and 400_10 are 2.4 � 0.8 μm and 8.3 � 6.0 μm, respectively. This is strong evidence of good adhesion between adjacent layers, so much so that grain growth has surpassed “parent interfaces”. Basically, there are four main theories to explain the bonding mechanisms: film theory, energy barrier, recrystallization and diffusion bonding. In the present case, the bonding is clearly consolidated be­ tween AA1050 layers by a combination among mechanical bonding (from the compression and shear in the rolling), diffusion bonding (enhanced by the temperature) and recrystallization phenomena. In the latter, it is suggested that grain growth during recrystallization may eliminate some oxide film in the surface that acts as a non-metallic barrier to the bonding [55]. Moreover, it is worth noting that the asymmetric rolling introduces an extra sliding at the central interface, which induces new grain nucleation and may be one of the bonding mechanisms; it assists the film oxide breakout in the surfaces (film theory). This phenomenon added to the grain growth in the intermediate annealing steps improves the interfacial integrity and confirms that cDRX and static/dynamic recovery are important for both interface consolidation and tensile properties. In fact, thin oxide layers at the surface or even contamination can be related to enhancement of the bond strength and allow further grain refinement of the microstructure around the interface [56]. However, in the present investigation using AA1050 Al alloy, oxide fragments were not observed at the magnification level (OM and SEM). In addition, very thin and stable layers of Al2O3 oxide are formed in aluminum alloys, so after many cycles of deformation in AARB, this surface contamination will hardly be observed within the recrystallized grains.

Fig. 14. Grain boundary maps (HAGBs in black and LAGBs in red) and grain size distribution after AARB processing and the extra AR pass samples: (a) 350_6 þ 1 and (b) 350_10 þ 1. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

to consolidate in the solid bulk. At 400 � C, adhesion was more efficient, but the last interface remained the weakest interface. Kamikawa and Tsuji [50] and Su et al. [51] have also reported similar bonding weak­ ness during the last ARB cycle. The main tensile properties, summarized in Table 5, demonstrated a strength enhancement and a decrease in the total elongation values after AARB processing. The increase in yield strength is directly related to smaller grain size, which leads to the accumulation of dislocations and results in a mechanical strength increase after ARB, as discussed by Scharnweber et al. [52]. The evolution of the bonding interface, which resulted in an increase of yield strength was also reported by other studies [51,53]. From Table 5, it also can be observed that the highest strength values were achieved for the 350_6 and 400_6 sheets. A possible explanation for this might be related to the grain size and the fraction of HAGBs. In the former, the average grain size is 1.1 � 0.5 μm with a very high fraction of HAGBs, close to 89% (see Table 3). This microstructure is typical of the SPD process and presents an ultrafine-grained structure. In contrast, the sample processed by 10 cycles at 350 � C presents a slight decrease in the strength and the HAGB fraction decreased to 69%, which is related to the extension of the cDRX and partial grain growth phenomena. Xing et al. [54] observed a work softening behavior – in which the ultimate tensile strength decreased and elongation increased – at accumulated strains higher than 4.8 in ARB of an AA8011 Al alloy at room temper­ ature. This behavior was intensified when the temperature was raised from room temperature to 200 � C and the highest strength was achieved at an accumulated strain of 2.4, which was associated to the cDRX and recovery processes (both static and dynamic) [54]. In a similar way, the cDRX phenomenon takes place in the AA1050 Al alloy submitted to AARB processing at a higher temperature range after several cycles. This may be the most important reason for the highest levels of the yield stress and ultimate tensile strength in the 350_6 and 400_6 sheets, and the softening behavior observed in the 350_10 and 400_10 samples. In addition, the asymmetric rolling combined with ARB definitely im­ proves the grain subdivision by shear in the earlier cycles and the rotational field in this process may accelerate the HAGB formation,

4.3. Effect of the extra AR pass on the mechanical response and microstructure In an attempt to improve the quality of adhesion between sheets 12

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx

during AARB, an extra AR pass was applied with a 50% thickness reduction. Upon closer inspection, the extra AR pass may be considered as another cycle. The grain size increased slightly to approximately 1.5 μm after 11 cycles at 350 � C. After the sixth cycle, the fraction of HAGBs decreased from 89% to the range of 65–75%. Therefore, these results confirm the trend of increasing the softening phenomena after 6 cycles. The most important result is in the fracture surface observations in Fig. 12-a and 12-b. They show the improvement of the sheet adhesion. The extra AR pass promoted greater bonding homogeneity along the sheet thickness in response to the reduction of the failure areas related to previous cycles. Despite of a strength decrease, Fig. 13-b shows the enhancement of total elongation. Additionally, Fig. 13-a demonstrates that the mechanical strength was slightly enhanced in the TD when compared to the RD. An important feature observed at 400 � C was an increase in the UTS and YS in the TD due to grain size distribution and necklace structure. These mechanical response differences observed among AARBed sheets rolled at 350 � C and 400 � C may be associated with the grain shapes – elongated and pancake-like at 400 � C – as the crystallographic texture and the intensity of the texture components were similar.

Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgments ~o Paulo We gratefully acknowledge the financial support of the Sa Research Foundation - FAPESP, Brazil (Grant number 2016/10997-0) and the National Council for Scientific and Technological Development CNPq, Brazil (Grant numbers 449009/2014-9 and 153585/2018-8). The authors would also like to thank the Brazilian Nanotechnology National Laboratory – LNNano for the use of their X-ray diffractometer. References [1] Y. Saito, N. Tsuji, H. Utsunomiya, T. Sakai, Novel ultra-high straining process for bulk materials-development of the accumulative roll-bonding (ARB) process, Acta Mater. 47 (1999) 579–583, https://doi.org/10.1016/S1359-6454(98)00365-6. [2] N. Tsuji, Y. Saito, S.H. Lee, Y. Minamino, ARB (Accumulative Roll-Bonding) and other new techniques to produce bulk ultrafine grained materials, Adv. Eng. Mater. 5 (2003) 338–344, https://doi.org/10.1002/adem.200310077. [3] T.C. Lowe, R.Z. Valiev, Producing nanoscale microstructures through severe plastic deformation, JOM-US 52 (2000) 27–28, https://doi.org/10.1007/s11837-0000127-8. [4] R.Z. Valiev, I.V. Alexandrov, Y.T. Zhu, T.C. Lowe, Paradox of strength and ductility in metals processed by severe plastic deformation, J. Mater. Res. 17 (2002) 5–8, https://doi.org/10.1557/JMR.2002.0002. [5] H.W. H€ oppel, R.Z. Valiev, On the possibilities to enhance the fatigue properties of ultrafine-grained metals, Z. Metallkd. 93 (2002) 641–648, https://doi.org/ 10.3139/146.020641. [6] H.W. H€ oppel, J. May, M. G€ oken, Enhanced strength and ductility in ultrafinegrained aluminium produced by accumulative roll bonding, Adv. Eng. Mater. 6 (2004) 781–784, https://doi.org/10.1002/adem.200300582. [7] J. May, H.W. H€ oppel, M. G€ oken, Strain rate sensitivity of ultrafine-grained aluminium processed by severe plastic deformation, Scr. Mater. 53 (2005) 189–194, https://doi.org/10.1016/j.scriptamat.2005.03.043. [8] M. Ruppert, W. B€ ohm, H. Nguyen, H.W. H€ oppel, M. Merklein, M. G€ oken, Influence of upscaling accumulative roll bonding on the homogeneity and mechanical properties of AA1050A, J. Mater. Sci. 48 (24) (2013) 8377–8385, https://doi.org/ 10.1007/s10853-013-7648-3. [9] L. Li, K. Nagai, F. Yin, Progress in cold roll bonding of metals, Sci. Technol. Adv. Mater. 9 (2) (2008), 023001, https://doi.org/10.1088/1468-6996/9/2/023001. [10] L. Liensh€ oft, P. Chekhonin, D. Z€ ollner, J. Scharnweber, T. Marr, T. Krauter, H. W. H€ oppel, W. Skrotzki, Static recrystallization and grain growth of accumulative roll bonded aluminum laminates, J. Mater. Res. 32 (24) (2017) 4503–4513, https://doi.org/10.1557/jmr.2017.386. [11] R. Jaamati, M.R. Toroghinejad, Cold roll bonding bond strength: review, Mater. Sci. Technol. 27 (7) (2011) 1011–1108, https://doi.org/10.1179/ 026708310X12815992418256. [12] P. Chekhonin, B. Beausir, J. Scharnweber, C.-G. Oertel, T. Haus€ ol, H.W. H€ oppel, H.G. Brokmeier, W. Skrotzki, Confined recrystallization of high-purity aluminium during accumulative roll bonding of aluminium laminates, Acta Mater. 60 (2012) 4661–4671, https://doi.org/10.1016/j.actamat.2012.04.004. [13] S.M. Ghalehbandi, M. Malaki, M. Gupta, Accumulative roll bonding – a review, Appl. Sci. 9 (2019) 3627, https://doi.org/10.3390/app9173627. [14] L. Su, C. Lu, A.K. Tieu, G. Deng, X. Sun, Ultrafine grained AA1050/AA6061 composite produced by accumulative roll bonding, Mater. Sci. Eng. A 559 (2013) 345–351, https://doi.org/10.1016/j.msea.2012.08.109. [15] G. Anne, M.R. Ramesh, H.S. Nayaka, S.B. Arya, S. Sahu, Microstructure evolution and mechanical and corrosion behavior of accumulative roll bonded Mg-2%Zn/Al7075 multilayered composite, J. Mater. Eng. Perform. 4 (2017) 1726–1734, https://doi.org/10.1007/s11665-017-2576-z. [16] S. Roy, B.R. Nataraj, S. Suwas, S. Kumar, K. Chattopadhyay, Microstructure and texture evolution during accumulative roll bonding of aluminum alloys AA2219/ AA5086 composite laminates, J. Mater. Sci. 47 (2012) 6402–6419, https://doi. org/10.1007/s10853-012-6567-z. [17] T. Sakai, A. Belyakov, R. Kaibyshev, H. Miura, J.J. Jonas, Dynamic and postdynamic recrystallization under hot, cold and severe plastic deformation conditions, Prog. Mater. Sci. 60 (2014) 130–207, https://doi.org/10.1016/j. pmatsci.2013.09.002. [18] B. Beausir, J. Scharnweber, J. Jaschinski, H.-G. Brokmeier, C.-G. Oertel, W. Skrotzki, Plastic anisotropy of ultrafine grained aluminium alloys produced by accumulative roll bonding, Mater. Sci. Eng. A 527 (2010) 3271–3278, https://doi. org/10.1016/j.msea.2010.02.006. [19] R. Jamaati, M.R. Toroghinejad, Effect of stacking fault energy on deformation texture development of nanostructured materials produced by the ARB process, Mater. Sci. Eng. A 598 (2014) 263–276, https://doi.org/10.1016/j. msea.2014.01.048.

5. Conclusions The ARB and AR techniques were combined to produce ultrafinegrained aluminum sheets with the random crystallographic texture and high-quality interface bonding at temperatures of 350 � C and 400 � C. Based on the experimental evidence presented in this study, the following conclusions were drawn: a. One of the most significant findings from this study was that when the bonding process is activated by diffusion concomitant with friction and shear strain between interfaces, high-quality interfacial layer sheets were obtained at both the studied temperatures. Two main factors improved the quality of the junction and can be enumerated: i) there is an optimal number of AARB cycles (between 4 and 6 for the AA1050) and the achievement of the optimal number increases the strain distribution, the bonding among previous cycles and attainment of grain size saturation; ii) the adoption of an extra AR pass improved interface bonding; b. The final microstructure after the AARB process was very refined. At 350 � C, the average grain size was about 0.6–1.0 μm, with 70% of HAGBs and the refinement process was controlled by cDRX. Alter­ natively, at 400 � C the continuous dynamic recrystallization was incipient forming a necklace structure of finer grains surrounding recovered areas; the average grain size was 8 μm; c. The method induces enough shear components to modify the texture and randomize the orientation distribution. At both AARB temper­ atures, the shear texture component distributions were more ho­ mogenous throughout the sheets compared to conventional rolling. Additionally, γ-fiber {111}//ND and rotated-cube components were the dominant feature in the samples; d. Increasing the amount of deformation resulted in the yield stress reaching a saturation level of around 160 MPa. This was due to grain refinement and improved bonding quality. In addition, the elonga­ tion was slightly enhanced after the extra AR pass. At 400 � C, the grain size and shape distribution were linked to the enhancement of mechanical strength in the TD. Data availability The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations.

13

R.P. de Godoi et al.

Materials Science & Engineering A xxx (xxxx) xxx [39] M. Kawasaki, Z. Horita, T.G. Langdon, Microstructural evolution in high purity aluminum processed by ECAP, Mater. Sci. Eng. A524 (2009) 143–150, https://doi. org/10.1016/j.msea.2009.06.032. [40] M.Z. Quadir, M. Ferry, O. Al-Buhamad, P.R. Munroe, Shear banding and recrystallization texture development in a multilayered Al alloy sheet produced by accumulative roll bonding, Acta Mater. 57 (2009) 29–40, https://doi.org/ 10.1016/j.actamat.2008.08.056. [41] H. Pirgazi, A. Akbarzadeh, Role of second phase particles on microstructure and texture evolution of ARB processed aluminum sheets, Mater. Sci. Technol. 25 (5) (2009) 625–631, https://doi.org/10.1179/174328408X332771. [42] E. Bagherpour, M. Reihanian, H. Miyamoto, Tailoring particle distribution nonuniformity and grain refinement in nanostructured metal matrix composites fabricated by severe plastic deformation (SPD): a correlation with flow stress, J. Mater. Sci. 52 (2017) 3436–3446, https://doi.org/10.1007/s10853-016-0632-y. [43] I. Kim, S. Akramov, H.B. Jeong, Texture and formability development of asymmetry rolled AA3003 Al alloy sheet, Int. J. Mod. Phys. B 22 (2008) 523–528, https://doi.org/10.1142/9789814261579_0082. [44] S. Wronski, B. Ghilianu, T. Chauveau, B. Bacroix, Analysis of textures heterogeneity in cold and warm asymmetrically rolled aluminum, Mater. Char. 62 (2011) 22–34, https://doi.org/10.1016/j.matchar.2010.10.002. [45] H. Jin, D.J. Lloyd, The different effects of asymmetric rolling and surface friction on formation of shear texture in aluminium alloy AA5754, Mater. Sci. Technol. 26 (2010) 754–760, https://doi.org/10.1179/174328409X405634. [46] S. Li, L. Yang, N. Qin, Development of through-thickness texture gradient and persistence of shear-type textures during annealing of commercial purity aluminium sheet processed by accumulative roll-bonding, J. Mater. Sci. Technol. 34 (2018) 821–831, https://doi.org/10.1016/j.jmst.2017.04.019. [47] J. Sidor, A. Miroux, R. Petrov, L. Kestens, Microstructural and crystallographic aspects of conventional and asymmetric rolling processes, Acta Mater. 56 (2008) 2495–2507, https://doi.org/10.1016/j.actamat.2008.01.042. [48] H. Jin, D.J. Lloyd, Evolution of texture in AA6111 aluminum alloy after asymmetric rolling with various velocity ratios between top and bottom rolls, Mater. Sci. Eng. A465 (2007) 267–273, https://doi.org/10.1016/j. msea.2007.02.128. [49] C.H. Choi, J.W. Kwon, K.H. Oh, D.N. Lee, Analysis of deformation texture inhomogeneity and stability condition of shear components in fcc metals, Acta Mater. 45 (1997) 5119–5128, https://doi.org/10.1016/S1359-6454(97)00169-9. [50] N. Kamikawa, N. Tsuji, Microstructure and mechanical properties of ARB processed aluminium with different purities, Mater. Trans. 57 (10) (2016) 1720–1728, https://doi.org/10.2320/matertrans.MH201519. [51] L. Su, C. Lu, H. Li, G. Deng, K. Tieu, Investigation of ultrafine grained AA1050 fabricated by accumulative roll bonding, Mater. Sci. Eng. A614 (2014) 148–155, https://doi.org/10.1016/j.msea.2014.07.032. [52] J. Scharnweber, W. Skrotzki, C.G. Oertel, H.G. Brokmeier, H.W. H€ oppel, I. Topic, J. Jaschinski, Texture, microsctructure and mechanical properties of ultrafine grained aluminum produced by accumulative roll bonding, Adv. Eng. Mater. 12 (10) (2010) 989–994, https://doi.org/10.1002/adem.201000067. [53] M. Eizadjou, H.D. Manesh, K. Janghoraban, Microstructure and mechanical properties of ultra-fine grains (UFGs) aluminum strips produced by ARB process, J. Alloy. Comp. 474 (2009) 406–415, https://doi.org/10.1016/j. jallcom.2008.06.161. [54] Z. Xing, S. Kang, H. Kim, Microstructural evolution and mechanical properties of the AA8011 alloy during the accumulative roll-bonding process, Metall. Mater. Trans. A 33 (2002) 1521–1530, https://doi.org/10.1007/s11661-002-0074-9. [55] J.M. Parks, Recrystallization welding, Weld. J. (1953) 209–222. [56] L. Su, C. Lu, A.K. Tieu, G. Deng, X. Sun, Ultrafine grained AA1050/AA6061 composite produced by accumulative roll-bonding, Mater. Sci. Eng. A559 (2013) 345–351, https://doi.org/10.1016/j.msea.2012.08.109.

[20] M. Naseri, E. Reohanian, E. Borhani, A new strategy to simultaneous increase in the strength and ductility of AA2024 alloy via accumulative roll bonding (ARB), Mater. Sci. Eng. A656 (2016) 12–20, https://doi.org/10.1016/j.msea.2016.01.020. [21] M. Alvand, M. Naseri, E. Borhani, H. Abdollh-Pur, Nano/ultrafine grained AA2024 alloy processed by accumulative roll bonding: a study of microstructure, deformation texture and mechanical properties, J. Alloy. Comp. 712 (2017) 517–525, https://doi.org/10.1016/j.jallcom.2017.04.117. [22] H. Jin, D.J. Lloyd, The reduction of planar anisotropy by texture modification through asymmetric rolling and annealing in AA5754, Mater. Sci. Eng. A399 (2005) 358–367, https://doi.org/10.1016/j.msea.2005.04.027. [23] S.B. Kang, B.K. Min, H.W. Kim, D.S. Wilkinson, J. Kang, Effect of asymmetric rolling on the texture and mechanical properties of AA6111-aluminum sheet, Metall. Mater. Trans. A A36 (2005) 3141–3149, https://doi.org/10.1007/s11661005-0085-4. [24] D.C.C. Magalh~ aes, A.M. Kliauga, M. Ferrante, V.L. Sordi, Asymmetric cryorolling of AA6061: strain distribution, texture and hardening behavior, Mater. Sci. Eng. A736 (2018) 53–60, https://doi.org/10.1016/j.msea.2018.08.075. [25] H. Jin, D.J. Lloyd, The different effects of asymmetric rolling and surface friction on formation of shear texture in aluminium alloy AA5754, Mater. Sci. Technol. 26 (2010) 754–760, https://doi.org/10.1179/174328409X405634. [26] S. Wronski, B. Ghilianu, T. Chauveau, B. Bacroix, Analysis of textures heterogeneity in cold and warm asymmetrically rolled aluminium, Mater. Char. 62 (2011) 22–34, https://doi.org/10.1016/j.matchar.2010.10.002. [27] J.L. Wang, Q.N. Shi, Orientation evolution and nucleation mechanism of UFGcopper prepared by SD-AARB with heat treatment, Appl. Mech. Mater. 33 (2010) 63–167. https://doi.org/10.4028/www.scientific.net/AMM.33.163. [28] A. Mendes, I. Timokhina, A. Molotnikov, P.D. Hodgson, R. Lapovok, Role of shear in interface formation of aluminum-steel multilayered composite sheet, Mater. Sci. Eng. A705 (2017) 142–152, https://doi.org/10.1016/j.msea.2017.08.025. [29] L. Chen, Q. Shi, D. Chen, S. Zhou, J. Wang, X. Luo, Research of textures of ultrafine grains pure copper produced by accumulative-roll bonding, Mater. Sci. Eng. A508 (2009) 37–42, https://doi.org/10.1016/j.msea.2008.12.018. [30] J.L. Wang, Q.N. Shi, T.C. Qian, S.H. Wang, X.K. Yang, Recrystallized microstructural evolution of UFG copper prepared by asymmetrical accumulative roll bonding process, Trans. Nonferrous Metals Soc. China 20 (2010) 559–563, https://doi.org/10.1016/S1003-6326(09)60178-3. [31] H. Mi, Q.N. Shi, J.L. Wang, L. Zhou, L.W. Chen, Effect of asymmetrical accumulative roll-bonding and heat treatment on the formation of ultra-fine twin copper, Mater. Sci. Forum 667–669 (2011) 611–616. https://doi.org/10. 4028/www.scientific.net/MSF.667-669.611. [32] S.H. Lee, T. Sakai, D.H. Shin, Fabrication of ultrafine grained aluminum by differential speed accumulative roll-bonding (ARB) process, Mater. Trans. 44 (2003) 1382–1385, https://doi.org/10.2320/matertrans.44.1382. [33] H.P. Ng, T. Przybilla, C. Schmidt, R. Lapovok, D. Orlov, H.W. H€ oppel, M. G€ oken, Asymmetric accumulative roll bonding of aluminium-titanium composite sheets, Mater. Sci. Eng. A 576 (2013) 306–315, https://doi.org/10.1016/j. msea.2013.04.027. [34] A. Wierzba, S. Mroz, P. Szota, A. Stefanik, R. Mola, The influence of the asymmetric ARB process on the properties of Al-Mg-Al multi-layer sheets, Arch. Metall. Mater. 60 (4) (2015) 2821–2825, https://doi.org/10.1515/amm-2015-0450. [35] H.J. Bunge, Texture Analysis in Materials Science, Butterworts, London, 1982. [36] J.J. Fundenberger, B. Beausir, Universit�e de Lorraine – Metz, JTEX – software for TextureAnalysis. http://jtex-software.eu/, 2015. [37] M.C.V. Vega, R.E. Bolmaro, M. Ferrante, V.L. Sordi, A.M. Kliauga, The influence of deformation path on strain characteristics of AA1050 aluminum processed by Equal-Channel angular pressing followed by rolling, Mater. Sci. Eng. A646 (2015) 154–162, https://doi.org/10.1016/j.msea.2015.07.083. [38] B. Bay, N. Hansen, D.A. Hughes, D. Kuhlmann-Wilsdorf, Evolution of FCC deformation structures in polyslip, Acta Metall. Mater. 40 (2) (1992) 205–219.

14