Evolution of microstructure and mechanical properties in a hypoeutectic Al–Si–Mg alloy processed by accumulative back extrusion

Author’s Accepted Manuscript Evolution of microstructure and mechanical properties in a hypoeutectic Al-Si-Mg alloy processed by accumulative back extrusion N. Haghdadi, A. Zarei-Hanzaki, H.R. Abedi, D. Abou-Ras, M. Kawasaki, A.P. Zhilyaev www.elsevier.com/locate/msea

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S0921-5093(15)30529-3 http://dx.doi.org/10.1016/j.msea.2015.10.066 MSA32917

To appear in: Materials Science & Engineering A Received date: 3 August 2015 Revised date: 17 October 2015 Accepted date: 17 October 2015 Cite this article as: N. Haghdadi, A. Zarei-Hanzaki, H.R. Abedi, D. Abou-Ras, M. Kawasaki and A.P. Zhilyaev, Evolution of microstructure and mechanical properties in a hypoeutectic Al-Si-Mg alloy processed by accumulative back e x t r u s i o n , Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.10.066 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Evolution of microstructure and mechanical properties in a hypoeutectic Al-Si-Mg alloy processed by accumulative back extrusion N. Haghdadi1, A. Zarei-Hanzaki1, H.R. Abedi1, D. Abou-Ras2, M. Kawasaki3,4, A.P. Zhilyaev5,6 1.

School of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, Tehran, Iran Helmholtz-Zentrum Berlin für Materialien und Energie GmbH, Hahn-Meitner-Platz 1, 14109 Berlin, Germany 3. Division of Materials Science and Engineering, College of Engineering, Hanyang University, Seoul 133-791, South Korea 4. Departments of Aerospace & Mechanical Engineering and Materials Science, University of Southern California, Los Angeles, USA 5. Institute for Metals Superplasticity Problems, Khalturina 39, Ufa 450001, Russia 6. Research Laboratory for Mechanics of New Nanomaterials, St. Petersburg State Polytechnical University, St. Petersburg 195251, Russia 2.

Abstract This study demonstrates the evolution of microstructure and mechanical properties of a hypoeutectic Al-7Si-0.4Mg (A356) alloy processed by accumulative back extrusion (ABE) at temperatures ranging from 200 to 500 ºC. ABE processing is one of the new severe plastic deformation techniques enabling one to produce relatively large ultrafine-grained materials in a cylindrical shape. One complete pass of ABE was estimated to introduce a reasonably homogeneous effective strain of ~3 as calculated by finite element analysis. Microstructural observation showed that globular α-Al primary phase was subdivided into fine substructures and Si particles having a fibrous shape were fragmented and spheroidized within the eutectic constituent through ABE processing. There was no evidence of homogeneous distribution of the fine Si particles in the α-Al phase after ABE. Mechanical testing at room temperature showed that both yield strength and ultimate tensile strength of the A356 alloy dramatically increased through ABE, especially at lower processing temperatures, as compared with the as-cast condition whereas there was no significant reduction in ductility at all processing temperatures. The experimental results were discussed with emphasis on the microstructure evolution involving dynamic recrystallization and deformation behavior including strengthening mechanisms and strain hardening in the Al-Si alloy. Keywords: Accumulative back extrusion; Al-Si alloy; grain refinement; mechanical properties; severe plastic deformation

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Corresponding Author. Tel.: +98 21 61114167; Fax: +98 21 88006076; E-mail address: [email protected]

1. Introduction The great interest in environmental conservation and fuel efficiency by light-weight components has attracted attention to aluminum alloys as promising substitutes for conventional ferrous materials [1]. In particular, due to their excellent casting characteristics, high corrosion resistance, good weldability and low thermal expansion, aluminum-silicon (Al-Si) alloys are widely used in the fields of automotive and aerospace industries. Although ductility and fatigue generally decrease with increasing Si content [2], Al-7Si (wt.%) alloy (about half-way to the eutectic composition) holds an optimum combination of microstructural uniformity and mechanical properties including good wear resistance [3].

The Al–7Si hypoeutectic alloy

consists of two phases: primary α-Al and a eutectic phase containing coral-like Si particles. Modification of the shape and size of Si particles together with refinement in the Al-matrix is the key to improve the mechanical performance of this alloy. The application of severe plastic deformation (SPD) is a promising method for producing ultrafine-grained (UFG) materials where grain sizes of a few micrometers to tens of nanometers can be successfully produced in relatively large bulk materials [4]. A number of SPD techniques have been developed in the last two decades including equal-channel angular pressing (ECAP) [5], high-pressure torsion (HPT) [6], accumulative roll bonding (ARB) [7], and friction stir processing (FSP) [8]. In particular, SPD is considered as a main post-casting approach for improving mechanical properties by microstructural refinement in Al-Si alloys and there have been a number of studies evaluating the microstructural evolution and/or improved mechanical properties in Al-Si alloys after processing through ECAP [9-15], HPT [14,18-19], ARB [20], and FSP [13,21].

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There have been many studies demonstrating new SPD methods in recent years [22] due to the considerable potential of UFG metals processed by SPD techniques. Besides those mentioned methods, accumulative back extrusion (ABE) is a newly developed SPD technique initially proposed by Zarei-Hanzaki et al. in 2009 [23]. In this method, a material in the shape of a cylinder is subjected to severe shear deformation under backward extrusion followed by twodimensional constrained back-pressing utilizing an innovative twin punch setup explained elsewhere [24,25]. This technique has a great potential for producing bulk UFG samples with its capability of being a continuous operation. ABE processing was successfully applied at elevated temperatures to demonstrate microstructural refinement and homogeneity, evolution of shear bands and development of texture in AZ31 [23-28] and AZ81 [29-30] magnesium alloys. Moreover, detailed microstructural and hardness homogeneity were investigated by increasing numbers of ABE passes at room temperature on pure Al [31] and pure Cu [32] and an alternative route of ABE reverse processing was demonstrated to provide excellent hardness homogeneity in pure Al [31].

Moreover, due to the evolution of microstructure through ABE, enhanced

mechanical properties were recorded in pure Al [33] and 2124 aluminum alloy [34]. There is a single report to date on an Al-Si-Mg alloy demonstrating the evolution of microstructure after ABE at 300 °C followed by semisolid isothermal treatment [35]. Nevertheless, there is lack of fundamental information on the microstructure evolution and mechanical properties of Al-Si alloys after ABE processing.

Accordingly, the present

investigation, as an extension of earlier work, was to study Al-Si-Mg alloy was processed by ABE under a wide range of temperatures with extensive measurements to study detailed microstructure evolution and room temperature tensile properties of this alloy.

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2. Experimental material and procedure The material used in this study was a Thixocast A356 aluminum alloy with the chemical composition (in wt.%) of 7.5 Si, 0.4 Mg, 0.15 Fe, 0.03 Cu, 0.03 Mn, 0.2 Ti, 0.05 Sr, remainder Al. Cylindrical work-pieces 18 mm diameter by 8 mm tall were machined from the as-received material. Processing by ABE was conducted over a wide range of temperatures between 200-500 °C. The earlier reports explained the conventional procedure and principles of ABE processing [23] and illustrated the processing facility with a twin punch setup [24]. Specifically, two separate straining stages constitute one pass ABE process: first the sample is back-extruded into the gap between the inner punch and the die, and subsequently the back-extruded part is compressed back until it reaches its initial cylindrical shape by inserting the outer punch as a hollow ram between the inner punch and the die. In the present work, ABE processing was carried out using a 300 kN servo-electric universal testing machine under a constant ram speed of 5 mm/min. Graphite spray was applied to reduce the friction between the work-piece and the tool surfaces. Temperature was controlled using a K-type thermocouple fixed in the die wall. In order to better understand the distribution of effective strain across the A356 alloy processed by ABE, 3D finite element analysis (FEA) was performed by a commercial FEA code, DEFORM TM-3D V10.0. The data of the hot compression stress-strain curves obtained for the same material [36] were employed for the analysis and the other critical parameters and material data used for the FEA software are given in Table1. Figure 1 describes the initial finite element mesh for analyzing the plastic deformation through ABE processing as well as the initial die, punches and specimen positions prior to running the simulations. Global automatic remeshing was utilized to accommodate the large strains during the simulations.

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Each work-piece was cut along the diameter and vertical cross-sections for microstructural analysis. For optical microscopy (OM) and scanning electron microscopy (SEM), the sample surfaces were chemically etched using Keller’s solution. For EBSD, the vertical cross-sections were polished with diamond paste and OP-U (Struers). The EBSD analysis was performed using a Zeiss UltraPlus scanning electron microscope equipped with an Oxford Instrument NordlysNano EBSD and an 80 mm2 X-Max X-ray detector. The applied acceleration voltage was 15 kV. The EBSD patterns were acquired and evaluated using the Oxford Instruments AZtec software package. The EBSD maps were recorded with point-to-point distances of 40 nm. Mechanical properties of the A356 alloy after ABE were examined using a tensile testing facility equipped with a small-scale sample holder using specimens with gage dimensions of 10 mm length, 1.5 mm width and 0.7 mm thickness. The tensile samples were prepared by wirecutting to have the gauge length along the diameters at the mid sectional-plane of the cylindrical samples both in as-received and as-processed conditions. The typical cylindrical sample after ABE processing and the machined tensile specimen were illustrated in an earlier report [33]. All of the tensile tests were conducted at room temperature and at an initial strain rate of 10-3 s-1 using a Gotech-AI7000 servo-controlled electronic universal testing machine. 3- Experimental results 3-1 Evaluation of effective strain by FEA The distribution of effective strain was estimated by FEA for the A356 alloy through ABE for 1 pass at a processing temperature of 300 °C. Results of this simulation in cross section are shown in Fig. 2: (a) very early stage of ABE where the inner punch moves one-half of the initial height of the sample, (b) the last stage of back extruding the sample by the inner punch, (c)

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the middle stage of back-compression, and (d) the processed condition through one complete pass of ABE. It should be noted that the scales for the estimated strain are lower in the early stages of ABE processing due to the lower levels of accumulated strain in the earlier processing stages; compare Fig. 2(a) and (b) with Fig. 2(c) and (d). It is apparent in Fig. 2(a) that the region at the inward wall which is in contact with the inner punch receives severe strain in the very early stage of processing whereas almost no strain is observed below the inner punch. With moving the punch down as shown in Fig. 2(b), the inward wall of the sample keeps receiving high strain continuously and the material between the punch and the bottom of the sample are under compression so that the bottom of the cylindrical sample receives high strain in this stage. Thus, almost all area within the sample receives effective strain except for the upper region of the outer side wall. During deformation by backcompression as shown in Fig. 2(c) and (d), the formerly localized strain is eliminated due to straining in the opposite direction from the one during back extrusion; instead, relatively homogeneous distribution of effective strain is achieved throughout the sample height except for the region close to the outer wall. However, as analyzed earlier [31], the outer wall region is anticipated to receive very high shear strain during ABE. The present FEA calculated an average imposed effective strain of ~3 in the A356 alloy after one complete pass of ABE at 300 °C. The detailed shear straining and effective strain predictions were demonstrated through FEA on an AZ31 magnesium alloy at different processing stages of one ABE pass [25,27].

3-2 Microstructure evolution 3-2-1 Overall microstructure examined by OM/SEM

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Figure 3 shows the microstructure of the A356 alloy in an as-cast condition where (a) is the general microstructure taken by optical microscopy, and (b) shows a closer view taken at a eutectic region by SEM where the sample was prepared by polishing followed by deep etching prior to analysis.

The alloy consists of discontinuous primary α-Al phase and continuous

eutectic region as seen in Fig. 3(a) where the sizes of the α globules vary in a wide range of ~20100 µm. Detailed measurements showed the microstructure contains 76 vol.% of primary α-Al with semi-globular morphology and 24 vol.% of eutectic constituent. Figure 3(b) reveals that most Si particles demonstrate a fibrous structure with average cross-sectional dimensions of ~0.5 µm and lengths of ~5 µm. Figure 4 shows the representative microstructures of the A356 alloy processed through one pass ABE at different temperatures of (a) 200 °C, (b) 300 °C and (c) 400 °C through one pass where each image is vertical to the sample height and the images have been taken from the center of the cross-sections. Despite the high strain imposed into the material through ABE, a constant globular appearance of α-Al occurred with the globules size being slightly refined at the lower processing temperatures. Image analysis revealed that the variation of the average globule size was between 23 to 40 µm in the processing conditions used and these sizes fell in the range of globule sizes in the initial microstructure. The evolution of microstructure without any geometrical redistribution of the eutectic Si particles after one complete pass by ABE is equivalent to the microstructure changes by processing through ECAP using routes C or BC for up to 8 passes at room temperature [9-13]. However, this inhomogeneous distribution of the primary alpha phase and Si particles in ABE and routes C or BC of ECAP contrasts with the processed microstructure after FSP [13,21] and HPT [14,16-18] where the primary and eutectic constituents are no longer apparent and the Si

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particles are dispersed in the primary Al phase during processing. The recent paper suggested that the as-cast structure with non-uniform distribution of Si particles remains constant even after imposing the von Mises effective strain of >9.0 [13]. 3-2-2 Detailed microstructures measured by EBSD and SEM EBSD analysis was conducted on the centers of the vertical cross-sections of the alloy after ABE for one pass at different temperatures, and a set of OIM micrographs were taken over a wide measurement area. A map demonstrating the grain boundary misorientation angles at the consistent area and an OIM micrograph taken at the eutectic constituent are shown in Fig. 5 for (a)-(c) the sample processed by ABE at 200 °C and (d)-(f) the sample processed at 500°C. In Fig. 5(b) and (e), the yellow lines denote low-angle grain boundaries having misorientations in the range of 2º to 15º and violet lines denote high-angle grain boundaries having misorientations of >15º. The EBSD results demonstrate several significant features of both the primary alpha and eutectic regions after ABE for one pass at these two different temperatures. First, the globules of the Al primary phase consist of cell structures after ABE at 200 °C as shown in Fig. 5(a) and these cell boundaries have misorientation angles of less than 15° which is apparent in Fig. 5(b), thereby leading to the formation of subgrains in the globules through ABE for 1 pass. The measurements gave an average subgrain size of ~1-2 µm. Second, ABE processing introduced significant grain refinement in the eutectic region where most grain boundaries demonstrated high angles of misorientations as shown in Fig. 5(b)-(c). The measurements showed an average grain size of ~0.44 µm in the eutectic region. Third, there is a similar microstructural changes after ABE at 500 °C, although there is a slight increase in (sub)grains size in both the primary and eutectic constituents as shown in Figs. 5(c)-(e). In practice, the measurements showed

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subgrain sizes of ~2-3 µm in the primary phase and a grain size of ~0.47 µm in the eutectic region in the sample processed at 500 °C. The EBSD results were further analyzed for evaluating the texture after ABE at these two temperatures. The texture maps for the consistent regions displayed in Fig. 5 (a) and (d) are shown in Fig. 6 (a) and (c) and the corresponding (111) pole figures are shown in Fig. 6 (b) and (d) for the samples processed at 200 °C and at 500 °C, respectively. The different colors in the maps and pole figures correspond to different texture: blue denotes shear C-component and red denotes recrystallization component [37]. The analysis demonstrated that ABE at 200 °C induced a mostly recrystallized microstructure, while processing at 500 °C introduced a predominant shear texture. This may be explained by the delay of strain driven dynamic recrystallization (DRX) observed also in an Al7% Si alloy processed by HPT at high temperature [17]. The study demonstrated that, at equivalent strains up to ~7, a cycle of “grain refinement-DRX-grain refinement” was observed in the sample processed at room temperature whereas the stage of “grain refinement– recrystallization” was delayed for strains of ~2 and not fully completed in the high temperature sample [17]. In the present experiments, the specimens taken at the centers of the processed cylindrical samples receive equivalent strains of ~3 after 1 pass as shown in Fig. 2 so that the sample after ABE at 200 °C may be able to show the recrystallized texture while simple deformed texture with delay in DRX is observed in the high temperature sample. Figure 7 shows the SEM micrographs taken at the center of the cross-sections in the specimens processed by ABE for one pass under different processing temperatures of (a) 200, (b) 300, and (c) 500 °C. It is apparent that Si particles are significantly fragmented and spheroidized through ABE processing for one pass at all processing temperatures. However, these Si particles

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remain within the eutectic phase and further observations showed no evidence of Si particles dispersion or nucleation in the Al matrix phase after ABE at any temperature. Quantitative analysis conducted to measure the average diameter and roundness coefficient of the Si particles in each processed sample are listed in Table 2 where a circular particle has a roundness coefficient of 100% and lower values represent more deviations from a circular configuration.

The measurements show that both the particle size and roundness

coefficient increase with increasing processing temperatures due to coarsening. Fig. 7(c) clearly shows the sample processed at the highest temperature to exhibit the largest Si particles size with wide inter-particle spacing in the eutectic region. The fragmentation of Si particles in the eutectic phase was demonstrated earlier on Al-Si alloys after 8 passes ECAP through route BC and C [10-14] and at the central area in the disk sample after 5 turns by HPT [14]. 3-3 Mechanical properties The mechanical properties of Al-Si alloys are dependent upon the α-Al phase characteristics as well as the size and distribution of Si particles which are influenced by the ABE processing temperature.

Thus, tensile testing was conducted at room temperature to

evaluate yield strength (YS), ultimate tensile strength (UTS) and elongation to failure using a miniaturized tensile testing method. The variations of YS and UTS are shown in Fig. 8 for the A356 alloy after processing by ABE at different temperatures ranging from 200 to 500 ºC. For the sake of comparison, the values of strength in the as-cast condition are also included in the chart. It is apparent that all samples after ABE processing demonstrated higher YS and UTS values compared with the ascast sample which is attributed to significant changes in microstructure including the formation of subgrains in the Al-α phase and fragmentation of Si particles in the eutectic region through

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ABE processing. This observation is supported by the Hall–Petch relationship [38,39], Bailey– Hirsch relationship [40] and the fundamentals of particle strengthening mechanisms. Especially, specimens processed at lower temperatures as shown in Figs 4(a)-(b) and 5 (a)-(c) are anticipated to receive these influences more significantly. Subsequently, the specimens after ABE at lower temperatures of 200-300 °C exhibited higher YS and UTS values which are about 1.5 and 2 times, respectively, of those in the as-cast condition. The improvement in both YS and UTS through ABE processing was compared with other SPD techniques [9,12,19-21] and it is shown in Table 3 where the results show that strengthening caused by ABE is comparable to other processing methods. In practice, a trend of higher strengthening at lower processing temperature is apparent during ABE. A consistent trend is observed in the alloy after HPT for 5-10 turns at RT and at 172 °C [19]. The corresponding elongations to fracture are given in Fig. 9 for the alloy in the as-cast condition and after processing at different temperatures ranging from 200 to 500 ºC. Regardless of processing temperature, there was no significant change in ductility after ABE processing. In practice, the processed samples record elongations of 7-8% while the as-cast sample demonstrates an elongation of ~9%. Although it may not be perfect to compare the values of elongation when there is a difference in the geometry of the tensile specimens especially in the thickness to gage length ratio, more significant drop in ductility was observed in Al-7% Si alloys after ECAP through route BC and C for 8 passes [11-13] while there is a large improvement in elongations to failure when processed by ECAP through route A and BA [11-13], by FSP [13] and HPT for 5-10 turns [19]. These reports suggest that the difference in ductility is closely affected by the degree of microstructural homogenization though processing. 4. Discussion

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4-1 Mechanisms of microstructural evolution during ABE In the present study, A356 Al alloy was processed through ABE for one pass at elevated temperatures.

The FEA results shown in Fig. 2 demonstrated that an overall reasonably

homogeneous effective strain is achieved along the sample height direction through the different straining stages each producing localized strain in specific locations by normal and shear straining [25,27]. Figs 4-7 show the significant changes in microstructure observed in the alloy after ABE at different temperatures with continuous DRX being the main mechanism for theses microstructural changes. The key driver for this is the severe shear straining that occurs during ABE, which generates sub-boundaries with dislocations wherever a strain gradient is introduced. In fact, geometrically necessary dislocations (GNDs) are required to accommodate the lattice curvature, and the GND density increases by increasing the strain gradient up to a steady state level [41]. These GNDs and dense dislocation walls then form cell blocks because less energy is required for deformation if a grain is split into cells, as cells deform through fewer than the five independent slip systems required for constrained deformation [42]. The formation of fine (sub)grains in the vicinity of Si particles in the eutectic region may be explained by the particles stimulated nucleation (PSN) mechanism [43]. One investigation shows that, due to intense strain incompatibility within the microstructure during plastic deformation, a deformation zone with a high dislocation density and large lattice misorientation are observed near non-deformable Si particles in Al–Si alloys [44]. Thus, the same recrystallization scheme provides a strong driving force for grain refinement within the eutectic constituent during ABE processing. Recrystallization through PSN was also reported during ECAP on an Al-7% Si alloy in an earlier report [13]. As suggested in Fig. 6, it is worth noting

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that there is delay in the occurrence of grain refinement and DRX when the sample is processed at high temperatures. Considering that the Si particles undergo fragmentation, spheroidization and ultimately coarsening in the thermally activated condition, the refinement in size of the Si particles may be explained by thermal and mechanical effects. Thermal disintegration of Si plates is a welldefined phenomenon which begins at local crystal defects such as terminations and kinks as well as holes and fissures in the Si plates (i.e., Rayleigh’s criteria) [45]. Moreover, the coefficient of thermal expansion of Si (2.6 μm/K) is much lower than that of Al (22.7 μm/K) which results in the generation of thermal stresses at Al/Si interfaces that, during heating, leads to brittle fracture at Si particles in the eutectic region [46]. According to the aforementioned mechanisms, higher processing temperature should lead to significant refinement of Si particles but this shows discrepancy with the results presented in Fig. 7. On the other hand, significant mechanical straining through severe plastic deformation can describe the fragmentation of Si particles in the Al-Si alloys. Under mechanical straining, significant strain incompatibility develops at the interface between the Si particles and the matrix due to a difference in their elastic constants and deformation behavior [47] more pronounced at lower processing temperatures. Additionally, a stress state with high stress triaxiality leads to an increase in the damage rate of the Si particles. Lower processing temperatures exacerbate the cracking and fracture of the Si particles due to the higher strength and hardening rate at these processing temperatures which is consistent with the present experimental results. Earlier studies have also reported that Si in an A356 alloy was refined during hot tensile and compression loading and that the refinement was more significant with decreasing deformation temperature [48,49]. This contrasts with the studies of Ogris [45] and Colley [50] who have reported an

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increase in a Si fragmentation rate with increasing temperature, but during the silicon spheroidization treatment. The absence of mechanical stress during spheroidization heat treatment, compared with severe plastic deformation that occurs during thermomechanical processing and ABE best explains the discrepancy between these results. Spheroidization of the eutectic Si particles is essentially driven by a decrease in interfacial and elastic strain energy [45] with the capillary forces due to curvature triggering mass transport from the sharp edges to the round corners. Higher temperatures will thus favor spheroidization due to faster migration of Si atoms through the aluminum matrix [45]. The higher kinetics of strain-induced diffusion further accelerates spheroidization as observed in the present study. 4-2 Improved mechanical properties after ABE The general mechanical behavior of UFG materials involves high strength as a consequence of the Hall-Petch relationship, although this is generally associated with lower overall ductility [51,52]. However, the present experiments demonstrated that the Al-Si alloy has potential for demonstrating significantly improved strength without losing much ductility through microstructural refinement during ABE. In terms of improving strength in the Al-Si alloy, as was well explained earlier [13], there may be contributions of strain hardening, grain refinement, solid solution hardening, dispersion strengthening, and precipitation hardening. Since there was no evidence of precipitation in microstructure in the present study, it is expected there are other hardening factors to contribute to increasing strength noted. The FEA analysis shown in Fig. 2, and the earlier detailed modeling [25,27], explain the feasibility of strain hardening after ABE for one pass, which provides a high amount of strain, and thus high dislocation density, within the cylindrical sample. From the EBSD analysis shown

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in Fig. 5, there was formation of substructure with high numbers of subgrains divided by lowangle grain boundaries within the Al-rich globules after one pass of ABE. The higher strength of this alloy at the lower ABE processing temperatures is due to the significantly refined eutectic constituent, and strengthening by dislocations being pinned by the homogeneously distributed fine Si particles. With increasing the processing temperature, grain refinement through ABE is moderate as shown in Figs 4-6 and strengthening by microstructural refinement becomes less effective. Higher processing temperatures also reduce fragmentation of Si particles; so less hardening occurs due to the coarser Si particles present at these processing temperatures. The earlier study by Estey et al. [53] has shown that some Si atoms start to dissolve in the Al-rich phase at the temperatures of about 500 °C. It is thus reasonable to assume that there is an increase in the Si content in the Al-rich phase after ABE processing at 500°C compared with that at lower processing temperatures. Therefore, solute strengthening in the alloy is more significant when processing at ~500 °C which is reflected as a slight increase in UTS at processing temperatures from 450 to 500 °C in the alloy as shown in Fig. 8. On the contrary, the ductility in the Al-Si alloy is significantly reflected by the degree of microstructural homogeneity. Earlier reports suggested that there is an apparent difference in mechanical properties when the samples receive different levels of microstructural homogeneity after ECAP by applying different processing routes and after FSP [11-13]. In addition, the microstructural measurements at the fracture tips of the processed samples demonstrated that the fracture propagation takes place interdendritically through the eutectic constituent in the nonhomogeneous microstructure of the samples after ECAP using route BC and C [12]. Accordingly, poor ductility was observed in these processed samples and it is consistent with the

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low ductility in the as-cast sample having coarse non-homogeneous microstructure but differs from the recorded high elongations to failure in the samples after ECAP using route A and BA which provided reasonably homogeneous microstructure in the alloy. Thus, the samples after ABE was anticipated to exhibit general UFG behavior of improved strength but instead poor ductility due to the refined but inhomogeneous microstructure after processing. Nevertheless, elongations to failure as shown in Fig. 8 presented that the reduction in ductility was very limited in the Al-Si samples after ABE compared with the as-cast sample at all testing temperatures. Figure 10 shows the variation of the degree of strain hardening by simply calculating the ratio of UTS to YS, σUTS/σYS, for the A365 Al alloy after ABE at different processing temperatures. Compared with the ratio of UTS to YS of 2.0 for the as-cast microstructure, Fig. 8, the strain hardening increases greatly to about 2.5-2.6 when processing at the lower temperatures between 200-300 °C, but decreases to about 2.2 with increasing processing temperature. The ductility of all ABE-processed samples were similarly independent of processing temperature and slightly lower than the as-cast material. However, more important is that ABE processing leads to significant strengthening especially at the lower processing temperatures without much loss in ductility compared to the as-cast state due in part to the high degree of strain hardening capacity. In practice, ABE processing refines the Si particles and induces a subgrain structure within the Al-rich phase, as shown in Fig. 7, which accommodates the dislocations generated during deformation including significant numbers of low-angle grain boundaries that form. This microstructure evolution leads to high strain hardening which delays localized deformation leading to the enhanced UTS in the A365 Al alloy. This dislocation intragranular activity and strain hardening is more pronounced within the grains of samples processed at lower temperatures of 200-300°C which is considered the

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preferred temperature range for processing. The difference in the improved UTS depending on the ABE processing temperature is apparent in Fig. 8 and Table 3.

5. Summary and conclusions In the present study, the microstructural evolution and improved mechanical properties have been investigated for an A356 aluminum alloy processed by ABE for one pass at 200500ºC. The following conclusions are suggested based on the present results and analysis. 1. Finite element analysis demonstrated that ABE processing of an A356 alloy introduces a high strain of ~3 with one pass. This results in a refined microstructure due to the formation of a cell structure in the globular α-Al primary phase and fragmentation and spheroidization of the Si particles, but without intermixing of the two phases that occurs, for example, in FSP processing. 2. These microstructural and texture changes are apparent especially at low processing temperatures compared with a delay in the cycle of grain refinement and DRX during higher temperature deformation. 3. Increasing the yield and ultimate tensile strengths is attributed to the significant microstructural refinement achieved by only one pass of the ABE process for all temperatures between 200-500 ºC. These higher strengths are achieved with minimal loss of elongation compared with the as-cast condition due to high degree of strain hardening introduced by evolved microstructure.

Acknowledgements

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The authors are grateful to C. Förster at Helmholtz-Zentrum Berlin, and A. Marandi, A.R. Khalesian and M.H. Maghsoudi at University of Tehran for their kind help during the experiments. This work was supported in part by the National Research Foundation of Korea funded by Ministry of Education under Grant No. NRF-2014R1A1A2057697 (MK) and in part by the Russian Science Foundation under Grant No. 14-29-00199 (APZ).

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Figure captions

Fig. 1

The initial finite element mesh as well as the initial die, punches and workpiece positions prior to running the finite element simulations.

21

Fig. 2 Distribution of effective strain at the cross-sections of the A356 alloy estimated by FEA for (a) the very early stage of ABE where the inner punch moves down a half of the initial height of the sample, (b) the last stage of back extruding the sample by the inner punch, (c) the middle stage of back-compression and (d) the processed condition through one compete pass of ABE at 300 °C.

(a)

(b)

Fig. 3 Microstructure of the A356 alloy in the as-cast condition where (a) shows the general microstructure taken by OM and (b) shows the close view taken at a eutectic acquired by SEM.

22

(b)

(a)

(c)

Fig. 4 The representative microstructures of A356 alloy processed by ABE at different temperatures of (a) 200, (b) 300 and (c) 400 °C through one pass.

(a)

(b)

(c) 23

(d)

(e)

(f)

Fig. 5 OIM images taken over the wide measurement area and photos demonstrating the grain boundary misorientation angles at the consistent area and OIM micrographa taken at the eutectic constituent for the samples processed by ABE at (a-c) 200 and (d-f) 500°C. In (b) and (e), the yellow lines denote low-angle grain boundaries having misorientations in the range from 2º to 15º and violet lines high-angle grain boundaries having misorientations of larger than 15º.

(b)

24

(d)

Fig. 6 The texture maps and (111) pole figures for the alloy after ABE at (a)-(b) 200 and (c)-(d) 500°C. In the EBSD-texture maps, the recrystallization texture component is marked in red and the shear C-component is marked in blue.

Fig. 7 SEM micrographs taken at the center of the cross-sectional planes in the specimens processed by ABE for one pass at different processing temperatures of (a) 200, (b) 300 and (c) 500°C.

25

500

YS 402

385

Strength (MPa)

400

291

300

280

270

199

200 100

UTS

157

154

128

98

125

127

0 As-received

200

300 400 450 o ABE Temperature ( C)

500

Elongation (%)

Fig. 8 Variations of YS and UTS for the A356 alloy after processing by ABE at different temperatures ranging from 200 to 500 ºC.

11 10 9 8 7 6 5 4 3 2 1 0

9.1 8.1

8

As-received

7.1

7.2

7.3

200

300 400 450 o ABE Temperature ( C)

500

Fig. 9 Variation of elongations to fracture at room temperature for the A356 alloy in an as-cast condition and after processing at different temperatures ranging from 200 to 500 ºC. 26

3

2.8

σUTS/σYS

2.6 2.4 2.2 2

1.8 1.6 100

200

300 400 ABE Temperature (o C)

500

600

Fig. 10. Variation of the degree of strain hardening by calculating the ration of UTS to YS, σUTS/σYS , for the A365 Al alloy after ABE at different temperatures.

Table 1. The utilized data for numerical simulation

Simulation data Friction coefficient

0.2

Punch speed (mm/min)

5

Initial temperature (°C)

300

Number/type of elements

45000/Tetrahedral

Relative remeshing interference depth 0.7 Material data Thermal conductivity (W m−1 K−1)

151

Specific heat capacity (J g−1 °C−1)

0.963

Young's modulus (GPa)

71

Poisson ratio

0.33

27

Table 2. Variations of Si particles size and roundness with ABE temperature

SPD technique

ABE

Properties in an as-cast condition YS (MPa)

UTS (MPa)

98

199

Processing conditions

Testing strain rate (s-1)

Properties after processing

Improvement (%)

YS (MPa)

UTS (MPa)

YS

UTS

1 pass at 200 °C

1.0 × 10-3

157

402

60

101

1 pass at 500 °C

-3

127

280

29

40

1.0 × 10

022°C

Average diameter 0.4

Roundness coefficient (%) 75

022°C

2.5

16

022°C

1.3

16

072°C

1.4

10

722°C

1.6

51

ABE temperature

28

Reference

this work

ARB

-

103

5 cycles at RT

1.67 × 10-4

-

269

-

161

[20]

ECAP

90

130

10 passes at RT

4.0 × 10-4

245

270

172

108

[9]

1.0 × 10

-3

200

245

122

56

[12]

1 pass at RT*

1.0 × 10

-3

171

251

29

48

[21]

5 turns under 6.0 GPa at RT

1.0 × 10-3

230

285

130

114

ECAP FSP

HPT

90 132

100

157 169

133

8 passes at RT

[19] 5 turns under 6.0 GPa -3 1.0 × 10 190 235 90 77 at 172 °C Table 3. The capability of different SPD techniques in improving yield strength (YS) and ultimate tensile strength (UTS) of Al-7% Si hypoeutectic alloys in room temperature (RT) tensile testing.

*There is an expected temperature rise of ~ 400–500 °C during FSP.

29