Microstructural evolution in Al–Mg–Sc alloy (AA5024): Effect of thermal treatment, compression deformation and friction stir welding

Journal Pre-proof Microstructural evolution in Al–Mg–Sc alloy (AA5024): Effect of thermal treatment, compression deformation and friction stir welding Sumit Chhangani, Suresh Kumar Masa, Rohit T. Mathew, M.J.N.V. Prasad, M. Sujata PII:

S0921-5093(19)31575-8

DOI:

https://doi.org/10.1016/j.msea.2019.138790

Reference:

MSA 138790

To appear in:

Materials Science & Engineering A

Received Date: 25 July 2019 Revised Date:

5 December 2019

Accepted Date: 7 December 2019

Please cite this article as: S. Chhangani, S.K. Masa, R.T. Mathew, M.J.N.V. Prasad, M. Sujata, Microstructural evolution in Al–Mg–Sc alloy (AA5024): Effect of thermal treatment, compression deformation and friction stir welding, Materials Science & Engineering A (2020), doi: https:// doi.org/10.1016/j.msea.2019.138790. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. 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. © 2019 Published by Elsevier B.V.

Microstructural evolution in Al-Mg-Sc alloy (AA5024): Effect of thermal treatment, compression deformation and friction stir welding Sumit Chhangani1, Suresh Kumar Masa2*, Rohit T Mathew1, M.J.N.V. Prasad1, Sujata M2 1

Department of Metallurgical Engineering and Materials Science, Indian Institute of Technology

Bombay, Mumbai-400 076, INDIA 2

Materials Science Division, CSIR-National Aerospace Laboratories, Bengaluru-560 017, INDIA

Abstract

Effect of thermal annealing and deformation in both compression and friction stir welding conditions on microstructural evolution in the cold rolled Al-4.36Mg-0.26Sc-0.09Zr (wt.%) alloy was investigated. To evaluate the thermal stability of the alloy, differential scanning calorimetry and static annealing experiments were carried out as a function of temperature. Microhardness measurements and quasi-static compression testing were performed on the as-received and annealed alloy samples. Friction stir welding was carried out on the as-received alloy sheet in butt configuration at two different tool traverse speeds of 250 and 500 mm/minute. Upon annealing, the alloy showed continuous recrystallization with transformation of the elongated grain structure possessing strong rolling texture to coarse equiaxed microstructure with random orientation. The annealed alloy exhibited reduced hardness and compressive strength at room temperature. Detailed microstructural investigation of hot compression deformed and friction stir welded samples revealed formation of subgrain structure and followed by fine recrystallized grains with nearly random orientation. The analysis suggests that continuous dynamic recrystallization involving progressive subgrain rotation is the possible mechanism for microstructural changes occurring during hot deformation of Al-Mg-Sc alloy. Keywords: Al-Mg-Sc alloy; Hot compression; Friction stir welding; Microstructure; Microtexture

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1. Introduction

Al-Mg-Sc alloys are becoming promising materials for many applications including aerospace and ship building owing to their unique combination of desired properties like low density, moderate tensile strength, good fatigue strength, better corrosion resistance and excellent weldability [1]. There have been several reports on the alloy development, the structure-property correlations and the superplastic forming of the Al-Mg-Sc based alloys. Willey [2] investigated the effect of scandium on mechanical properties of Al-alloys and reported that addition of small amount of Sc to Al-Mg alloy increases the strength significantly. This is attributed to the formation of fine, homogeneous and coherent Al3Sc precipitates. The favorable microstructures developing in the alloy results in an attractive combination of strength and ductility [3]. Although dynamic recrystallization (DRX) occurs through strain-induced grain boundary bulging [4,5], the alloy exhibits high thermal stability [6] and creep properties [7] as compared to other Al-alloys due to stability of Al3(Sc,Zr) precipitates. The grain refinement by severe plastic deformation and cryorolling has been reported by several authors [8-10]. Avatokratova and coworkers [11,12] demonstrated high strain rate superplasticity with elongations beyond 2000% in Al-Mg-Sc alloy. Integral structures can be fabricated via many emerging technologies like high speed machining, laser beam welding, friction stir welding (FSW) etc. Among them, FSW, a solid state welding technique, has received a greater attention especially in aerospace industry because it offers lower distortion and greater welding efficiency [13]. However, severe plastic deformation and metal flow during FSW process affect the local microstructure significantly in the weld zone. In general, the microstructure of FSW joint comprises of three zones namely weld nugget zone (stir zone, SZ in friction stir processing, FSP), thermo-mechanically affected zone (TMAZ) and base metal (BM) [14]. The grains within the stir zone are generally found to be fine equiaxed and

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the observed grain refinement in this zone is attributed to dynamic recrystallization [4,5]. However, it is not yet clear from previous reports that whether the grain refinement in Al-Mg-Sc alloys is a result of discontinuous DRX, continuous DRX (extended recovery), geometric DRX or particle stimulated nucleation (PSN) that contributes to the grain refinement in Al-Mg-Sc alloys [15]. A few efforts have also been made to simulate the recrystallized microstructure in FSW using plane-strain hot compression test, shear compression specimen test, torque specimen test etc. [16, 17]. Limited reports are available in open literature on FSW of Al-Mg-Sc alloys. Lapasset and coworkers [18] studied the FSW microstructure and its influence on the property, and reported that decrement in strength of FSW joint is due to occurrence of recrystallization at the expense of fine unrecrystallized microstructure of the base material. Munoz and coworkers [19] identified that limited dissolution of Al3Sc precipitates at FSW temperatures results in lesser drop in strength values as compared to fusion welding. With regard to the dynamic mechanical properties, Besel et al. [20] reported that fatigue cracks are initiating around the stir zone and adjacent TMAZ due to complex local microstructure that exists in the weld zone. Mishra and coworkers [21-23] have conducted several investigations related to microstructure evolution, thermal stability and mechanical properties (tensile and fatigue) of AA5024 alloy subjected to FSP. It is known that the AA5024 alloy in H116 temper condition exhibits excellent microstructural stability under thermal and mechanical deformation conditions. Conversely, there are a limited number of studies available on the effect of heat treatment temperature and time on microstructure and texture evolution in AA5024-H116 alloy. It is well established that the material flow during FSW occurs under a combination of significant temperature rise and complicated stress state with a major component of compressive deformation [24]. However,

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there is a scarcity of literature on microstructure and texture evolution under monotonic compressive deformation as well as during FSW of AA5024 alloy. In view of this, the present study was undertaken with a specific objective to investigate systematically the effect of thermal treatment and plastic deformation (under monotonic and complicated stress state) on microstructure (including texture) evolution in a commercially available Al-Mg-Sc (AA5024) alloy.

2. Material and Experimental Methods

Commercially available rolled Al-Mg-Sc alloy (AA5024) sheets of 7 mm and 3.2 mm thickness having a nominal composition (in weight percent) of Mg: 4.36, Mn: 0.35, Fe: 0.34, Si: 0.32, Sc: 0.26, Zr: 0.09 and remaining Al in H116 (work hardened and special corrosion resistant) temper condition were used for the present investigation. In order to assess the thermal stability of the as-received alloy, differential scanning calorimetry (DSC) studies were conducted on the as-received material. A sample of ~50 mg was heated continuously in aluminum pans covered with lids under argon atmosphere from room temperature to 903 K at a rate of 10 K/minute in TA DSC instrument. A second DSC heating scan was performed on the same sample under identical conditions to obtain the baseline. In addition, the as-received alloy samples were annealed isothermally in a tubular furnace for 24 h at different temperatures in the range of 623-873 K for evaluating the microstructural evolution. Cylindrical specimens of ~4.5 mm diameter and ~7 mm long were prepared by electricdischarge machining from the 7 mm thick as-received alloy sheets. These were used for examining the effect of compression deformation on microstructure evolution in both rolled and annealed conditions and to correlate with friction stir processed/welded microstructure of the as-

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received alloy. The specimens were cut such that the compression axis was parallel to the short transverse direction of the initial material. Compression tests were performed up to a true strain of 0.5 on each specimen at different constant cross-head velocities using a Tinius-Olsen instrument equipped with a load cell capacity of 25 kN. The as-received alloy and the annealed specimens were deformed at room temperature with an initial strain rate of 1 × 10-3 s-1. To examine the test temperature effect on flow behavior, the as-received alloy specimens were deformed in compression with an initial strain rate of 1 × 10-3 s-1 at different test temperatures of 473 K, 523 K, 623 K and 723 K. Also, the as-received specimens were deformed at different strain rates of 1 × 10-3 - 1 × 10-1 s-1 at a few selected test temperatures. To assess the effect of microstructural changes with annealing, Vickers microhardness measurements were made using a peak load of 5 N and 10 s dwell time on three faces (long, L; long transverse, LT; short transverse, ST) of the as-received alloy sheet and the annealed samples. A single pass friction stir welding (FSW) was performed on the as-received alloy sheets having 3.2 mm thickness with butt joint configuration to investigate microstructural evolution in and away of the stir zones of samples processed under two different tool transverse speeds of 250 mm/minute and 500 mm/minute. For FSW, a high speed steel cylindrical tool with probe diameter of 5 mm and shoulder diameter of 15 mm was used. The FSW was carried out at a tool rotation rate of 800 revolutions per minute with the tool tilt angle of 2° and the plunge depth of ~3 mm. The FSW was done on L-LT plane with direction perpendicular to the rolling direction and the tool bit axis was in normal direction of the alloy sheets. For examining the microstructure and microtexture of the as-received alloy, annealed samples, selected compression deformed specimens and FSW samples, orientation imaging microscopy (OIM) was performed at different locations by using field emission gun scanning

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electron microscope (FEG-SEM) attached with EDAX-OIMTM electron backscattered diffraction (EBSD) detector operated at 20 kV. The OIM scans were obtained with a step size in the range of 60-200 nm depending on the microstructural feature size. More than 1000 grains were considered for calculating the average grain size, and the equivalent circular diameter method was used for evaluating the grain size from the OIM data. The inverse pole figures (IPFs) and the pole figures were generated from the OIM data using TSL-OIMTM v7.2 software. A threshold value of confidence index (C.I) greater than 0.1 was used for analyzing the OIM data to generate pole figures and inverse pole figures. Some selected samples were examined by Carl Zeiss Auriga Compact 45-58 dual beam focused ion beam (FIB)-scanning electron microscope (SEM) to obtain microstructures by ion channel contrast imaging. For OIM and FIB examination, samples were polished mechanically using SiC abrasive papers, followed by colloidal silica polishing and electropolishing in a solution made of 20% perchloric acid and 80% methanol solution at a voltage of 20 V and a temperature of ~263 K. Transmission electron microscopy was performed on selected samples using JEOL 2100F field emission gun transmission electron microscope (FEG-TEM) operated at 200 kV for the analysis of grain structure and precipitates/dispersoids. For TEM examination samples were ground to a thickness of ~100 µm and then ion-milled by precision ion polishing system (PIPS) from GATANTM instrument for obtaining the electron transparent foils.

3. Results

Figure 1a shows the triplanar colored orientation imaging microscope-inverse pole figure (OIM-IPF) images taken on three different faces of the alloy sheet in as-received condition. The as-received alloy displayed a typical cold rolling texture of Al i.e. majorly copper {112}<111>

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texture component [25] on the L(RD)-LT(TD) face as seen in the pole figure shown in Fig. 1b. Figure 1c reveals the pole figure showing the corresponding deformation texture observed on the L(RD)-ST (ND) plane. Owing to heavy cold rolling, the deformation texture of the as-received alloy on L-ST plane appeared to be as <111> fiber texture. Bright-field and dark-field TEM micrographs along with the selected area electron diffraction (SAED) pattern of the as-received sample in L-ST plane are presented in Fig. 1d and 1e. The presence of heavily deformed microstructure with elongated grains of ~200-300 nm wide confirmed that the as-received alloy is in work hardened temper condition. Further, the alloy contained a very fine particles (<20 nm, as highlighted by arrows in Fig. 1d) within grains and along grain boundaries as well. It is known fact that the 5024 alloy contains numerous Al3Sc or Al3(Sc1-x, Zrx) precipitates which play a very important role in microstructural stability of the alloy [26]. Figure 2 shows the DSC thermograms obtained by heating and reheating the as-received alloy within same temperature range at a constant heating rate of 10 K min-1. The base line corrected thermogram obtained by subtracting the second heating DSC scan from the first heating scan is also presented in Fig. 2. The as-received alloy displayed a significant change in slope in the thermogram after 623 K and exhibited two exothermic peaks at ~713 K and ~813 K and one strong endothermic peak at ~867 K with an onset temperature of ~853 K. Figure 3a-c shows the microstructure evolution in three different planes of the alloy samples upon soaking the as-received alloy for 24 h at three different temperatures of 623 K, 793 K and 873 K, respectively. At 623 K, the alloy has just started showing some noticeable changes (recovery) of microstructure, especially in LT-ST plane (Fig. 3a). The alloy still exhibited the elongated microstructure, however the grains were wider and nearly strain-free with little change in texture after heat treating at 793 K (Fig. 3b). Upon thermal treatment at 873 K, the alloy

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experienced significant changes in microstructure and texture, especially in L-ST plane (Fig. 3c). The elongated grain structure replaced with nearly equiaxed and coarse grains of ~50-100 µm. The initial strong deformation texture present in the as-received condition was almost weakened. Figure 3d shows the secondary electron micrograph on LT-ST plane of the sample heat treated at 873 K for 24 h. It is noted that there was substantial coarsening of precipitates (>100 nm) and they were present not only along grain boundaries but also inside grains. The room temperature mechanical characteristics of the as-received and the heat treated samples were evaluated by Vickers microindentation and compression testing. Figure 4a illustrates the effect of thermal treatment on variation in hardness of the alloy on all three different planes. As anticipated, the alloy exhibited highest hardness in the as-received condition and its hardness decreased gradually with increasing the heat treatment temperature. It is noted that the alloy hardness is relatively higher in L-ST plane as compared to that in other planes. Figure 4b shows the true stress-true strain curves obtained in compression up to a true strain (ε) of 0.5 at strain rate of 1 × 10-3 s-1 for the as-received and the heat treated specimens. The asreceived alloy exhibited a compressive strength of ~400 MPa, whereas the specimen heat treated to 623 K for 24 h shows a slight reduction in the strength up to ~370 MPa. Further heat treatment at higher temperatures led to significant reduction in strength of the alloy. This corroborates with the Vickers microindentation measurements. Figures 5a shows the true stress-true strain curves obtained by deforming the as-received alloy in compression at five different temperatures but at a fixed initial strain rate of 1 × 10-3 s-1. The alloy exhibited typical hot deformation flow behavior characteristic features. The flow curves displayed initially a slight strain hardening with a peak stress, followed by flow softening and then steady state flow behavior. With increasing the test temperature, there was significant

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drop in flow stress and the strain value corresponding to peak stress also decreased. The lowest stress value of ~4 MPa was noted at ε~0.1-0.3 in flow curve obtained at 723 K. Although the asreceived alloy was in heavily deformed condition, it showed minor work hardening during initial stages of deformation at room temperature (RT) and 473 K. In order to study the effect of strain rate, the as-received alloy was deformed in compression at 623 K under two additional strain rates (1 × 10-2 s-1 & 1 × 10-1 s-1). Figure 5b shows the flow curves of the alloy as a function of strain rate. It is noticed that there was a drop in true stress after a peak stress value, followed a saturated stress up to a certain strain and then a slight raise in true stress. Furthermore, the alloy exhibited a very high positive strain rate sensitivity i.e. significant increase in flow stress with increasing strain rate. Microstructural investigation of the deformed samples can provide more insights for the observed deformation characteristics of the alloy. Figure 6 shows the microstructures of the asreceived samples after compressive deformation at a strain rate of 1 × 10-3 s-1 but at two different temperatures of 523 K (Fig. 6a-c) and 623 K (Fig. 6d-f). Figures 6a, 6b and 6d are OIM-IPF maps, Fig. 6c and 6f are secondary electron micrographs and Fig. 6e is FIB image. It is noted that the grain flow under compression is wavy in nature; however, lesser waviness at higher temperature which indicates easy metal flow with increasing temperature. It is interesting to note that angle of waviness is nearly 45-50° to the compression axis. The wavy flow behavior, especially at low temperature of 523 K can be due to severe shear deformation of prior elongated grain structure. Figure 6b is the higher magnification OIM-IPF map of a small selected region of Fig. 6a obtained under finer EBSD step size of 60 nm. It reveals fragmentation of prior elongated grain structure and thereby formation of fine grained structure during low temperature deformation. There are some regions of dynamically recovered (elongated) grain structure in

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sample deformed at 523 K (Fig. 6c). The sample deformed at higher temperature of 623 K displayed extended recovery with formation of subgrain within prior elongated grain structure and some recrystallized grains as well. Moreover, the initial rolling texture of the as-received alloy has started weakening during hot compression, especially at temperatures where the recrystallized grain structure start forming. Figure 7 illustrates the effect of strain rate on microstructural evolution in compression at 623 K. The triplanar OIM-IPF images (Fig. 7a) and FIB image on L-LT plane (Fig. 7b) of the sample deformed at 1 × 10-1 s-1 revealed that the alloy underwent substantial amount of extended recovery in all planes. It is observed that the extent of shear band formation, subgrain formation and recrystallization, and randomization of crystallographic orientation increased with increasing strain rate (Fig. 6 and Fig. 7). Figures 8 and 9 illustrate the evolution of microstructure and texture in AA5024 alloy sample subjected to friction stir welding at a traverse speed of 250 mm/minute. The OIM-IPF maps were obtained on L-ST plane of the alloy sheet at five different locations -three within friction stir processed (FSP) region and two near seam/base material (BM) interface but one each near to the advancing and the retreating sides of FSW, respectively as shown in Fig. 8. Figure 9 presents (111) pole figures (Fig. 9a-d) and [001] inverse pole figures (Fig. 9e-h) obtained from OIM analysis of the as-received alloy sample and at three different locations (marked as 1-3 in Fig. 8) of the FSW sample. There are three important observations: (i) the alloy exhibited equiaxed grain structure with nearly random crystallographic orientation and large fraction of high angle grain boundaries in nugget zone upon FSP, (ii) the regions (locations-2 and 4) within the nugget zone but close to the weld seam showed relatively finer grain structure (d~2 µm) as compared to the center region (location-1, d ~3 µm) of the nugget zone, and (iii) the

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microstructure between the weld seam and BM (locations-3 and 5) appeared to be similar to the as-received condition and it consisted of a wavy and elongated grain structure possessing slightly strengthened <111> texture on L-ST plane. Figure 10 presents the effect of tool traverse speed on microstructure and texture evolution during FSW of the alloy. The center region of nugget zone (loaction-1) of the sample friction stir welded at a tool traverse speed of 500 mm/minute was examined by EBSD. The OIM-IPF image (Fig. 10a) revealed that the alloy exhibited comparatively a reduced grain size (d ~2.1 µm) upon increasing the traverse speed. This observation is in contrast to that reported by Besel et al. [20] on a similar grade Al alloy. However, the (111) pole figure (Fig. 10b) and [001] inverse pole figure (Fig. 10c) indicated that the level of random texture development upon FSP is similar to that observed at lower tool traverse speed (Fig. 9d&h). Figure 11 shows the TEM results of the nugget zone (location-1) region of the sample friction stir welded at traverse speed of 250 mm/minute. The TEM images captured at three different magnifications are presented in Figures 11a-c. The bright-field micrographs (Fig. 11a & 11b) and high resolution TEM image (Fig. 11c) clearly revealed the presence of many fine precipitates (~10-50 nm) within fine grains and along grain boundaries as well. TEM-EDS line analysis (shown in Fig. 11d) performed on one of these precipitates confirmed that these precipitates are mostly Al-Sc based intermetallics. The presence of Zr atoms could not be detected and might be due to its very low concentration.

4. Discussion

The experimental observations reported in the above results section clearly indicate that the microstructure of AA5024 alloy is strongly influenced by the heat treatment temperature, the

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test temperature and strain rate under compression and friction stir welding conditions. The observed microstructural changes under different conditions and their effect on deformation behavior are discussed in-detail in the following sub-sections.

4.1. Effect of thermal treatment on microstructural evolution and mechanical properties The as-received 5024 alloy, which was in heavily deformed condition exhibited two exothermic peaks and then one endothermic peak in its DSC thermogram (Fig. 2) upon heating from room temperature to 903 K. These observations are somewhat similar to Kaiser et al. [27] investigations on thermal stability of 75% cold rolled Al-6Mg-(0-0.6) Sc alloy studied by hardness and DSC. They reported 643 K and 733 K as the recovery and recrystallization temperatures, respectively for Sc-containing Al alloy. Toropova et al. [28] reported that the fine precipitates in Al-Mg-Sc alloy start losing in their coherency and then the semi-/incoherent precipitates grow when heated at temperatures greater than 623 K. In addition, they also reported that Sc-containing Al alloys show static restoration process at temperatures above 648 K, which is nearly 100 K greater than other comparable non-Sc containing Al alloys. Jones and Humphreys [29] found that precipitates coarsening and coherency loss in Al-Sc alloys is associated with the movement of low angle boundaries during recovery. Therefore, the onset of slope change observed in DSC thermogram of the present alloy could be related to the starting of Oswald ripening of the particles and followed by initiation of static restoration process. A review article by Davydov et al. [30] indicated that the antirecrystallized effect (higher resistance to recrystallization) of the alloy is attained due to the pronounced effect of numerous fine precipitates of Al3(Sc1-xZrx). The two exothermic peaks observed in the present study can be ascribed to concurrent precipitate coarsening and static recovery (~713 K), and extended recovery (recrystallization, ~813 K) processes, since the alloy containing fine precipitates was in

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heavily deformed condition. According to Al-Sc alloy phase diagram, the Al3Sc particles start dissolving into Al matrix at temperatures above 573 K. However, the exact solvus temperature depends on Sc content as well as other solute elements (Mg and Zr) concentration in the matrix. For example, the pseudo binary phase diagrams between Al, Sc and Mg indicate that the solvus temperature of the precipitates increases but liquidus temperature of the alloy decreases with increased concentration of these solute elements [31, 32]. Based on this, the observed strong endothermic peak (~867 K) can be attributed to dissolution of the particles and partial melting of the alloy. The prior high angle grain boundaries become more planar due to the presence of secondphase particles which prevent local migration of boundaries. However, owing to coarsening and dissolution some precipitates would have lost their pinning ability and thereby limited control over grain growth. The observed microstructural stability of the alloy as a function of temperature corroborates the noted changes in DSC thermogram. Unlike discontinuous recrystallization where the initial strong deformation texture is replaced by a strong cube texture, the rolling texture is almost retained with a little change during continuous recrystallization of single phase alloys [33]. However, in quasi-single phase alloys like AA5024 alloy in which precipitates control grain boundary migration, the rolling texture can no longer exist under conditions of loss of pinning ability, especially at higher annealing temperatures (secondary recrystallization) [34]. The microhardness measurements and deformation behavior in compression at room temperature are in accordance with the microstructural changes. The as-received alloy showed highest hardness and yield strength owing to its initial work hardened temper condition, and its hardness and yield strength decreased upon annealing due to static restoration processes. The

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observed hardness values of the as-received alloy are in close agreement with that of earlier reported studies on a similar alloy [11]. Further, the decreasing trend in hardness upon annealing is also consistent with previous studies on similar alloy [11]. As anticipated, the fully recrystallized specimen (873 K 24 h) displayed significant strain hardening up to ε = 0.5. Whereas, the deformed (as-received) and partially recovered/recrystallized specimens showed initial rapid strain hardening followed by saturation in stress due to balancing of work hardening and softening mechanisms with further straining.

4.2. Effect of compression deformation on flow behavior and microstructural evolution at elevated temperature The work hardening could be associated with dislocations activity and their interaction with the precipitates present in the alloy. Typically, flow softening i.e. decrease in flow stress with increasing strain is due to dynamic restoration processes (recovery and recrystallization), texture softening, deformation heating and/or any other microstructural instabilities like cavitation/voids formation, coarsening of precipitates, etc. It is hard to expect any deformation heating at such lower strain rate of 1 x 10-3 s-1. Since the Al alloy is a high stacking fault energy material, there could be an easy cross-slip activity and thereby an increased dynamic recovery with further straining even at lower deformation temperatures. At higher temperatures, thermally activated processes such as enhanced cross-slip, dislocation climb, grain boundary sliding and migration, etc., can contribute to early flow softening in the alloy. At higher test temperatures (>523 K), the alloy attained a steady state flow behavior at small strains with very low flow stress values (<50 MPa) at a strain rate of 1 x 10-3 s-1 (Fig. 5a).

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The observed flow behavior at higher test temperatures (Fig. 5a & 5b) i.e. strain hardening, then softening, and again hardening is not unusual, especially in Al-Mg-Sc alloys when deformed in compression at high temperatures under quasi-static strain rates. For example, Smirnov et al. [35] reported similar flow curves for Al-Mg-Sc-Zr alloy at 603 K and 633 K in compression. On similar lines of their observations, it can be explained that the drop in true stress after attaining the peak stress due to work hardening is because of favorable conditions for softening by dynamic recrystallization. On the other hand, the secondary hardening could be a resultant of (i) decelerating dynamic recrystallization under the influence of precipitates, (ii) dynamic grain growth due to coarsening of precipitates, and (iii) increased frictional effects at larger strain under compression. In addition, it is to be noted that the compression tests were conducted at constant cross-head speeds. This means, there will be a gradual increase in true strain rate with strain in compression, which can lead to an increase in flow stress at higher strain. It is worth to mention that low flow stresses and higher strain rate sensitivity are characteristic features of superplastic phenomenon. The Al-Mg-Sc alloys are known to exhibit superplastic behavior at low strain rates but at temperatures above 623 K with typical characteristics of very low stress values (<20 MPa) and high strain rate sensitivity (m>0.3) [3, 4, 9, 12]. This is majorly because of the fine grained microstructure formation due to DRX which undergoes significant grain boundary sliding (GBS), one of the dominant strain contributing mechanisms. The as-received Al alloy contained a high density of dislocations due to prior cold working. It can be expected that these will transform into subgrain structure with low angle grain boundaries during heating (static recovery) if the test temperature is high enough. Similar to static recovery, dynamic recovery i.e. formation of subgrain structure occurs during hot

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deformation due to increased dislocation cross-slip, glide and climb activities. Further, it is important to note that the applied stress facilitates stress-assisted boundary migration due to additional driving pressure. This indicates that the subgrain size (ds) is a strong function of flow stress (σ) and it can be represented as [33, 36]: ݀௦   =  

ఈீ௕ ఙ

(1)

where, G is shear modulus, b is Burgers vector. Here, α is a constant and has a value of ~10 for fcc metals. This equation (1) elucidates that higher the flow stress, finer the subgrain. As noted earlier from flow behavior of the alloy as a function of strain rate (Fig. 5b), the alloy exhibited an increased flow stress requirement upon increasing strain rate. Hence, the sample deformed at higher strain rate can contain finer subgrains. The subgrains in metals are usually considered to be transient microstructural features because of progressive subgrain rotation and thereby formation of strain-induced high angle boundaries due to local variation in dislocation density and subgrain boundary tensions during subsequent hot deformation. The process of continuous reactions of formation of subgrain and followed by recrystallization during deformation is usually referred to as continuous dynamic recrystallization (cDRX). Owing to their high stacking fault energy, the aluminum alloys are known to undergo cDRX during hot deformation [33], which is the case in the present alloy as well. The role of numerous fine precipitates, which are appeared to be thermally stable, present in the as-received alloy also needs to be considered for explaining the microstructural changes taking place during hot deformation. The nature of deformation structure in the vicinity of particles/precipitates dictates the occurrence of the well-known particle-stimulated nucleation (PSN) of recrystallization process. Further, it has been reported that the particles are expected to affect the subgrain structure

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formation, if the interparticle spacing is less than the subgrain size [33]. It is unlikely to have PSN in the present alloy, as the size of precipitates is much finer than the critical size (typically greater than 1 µm) required for it under current set of deformation conditions. Moreover, Ridley et al. [37] reported that the subgrain structure and its misorientation development in Al alloys is not much affected by the presence of small stable Al3Zr particles. However, stable and high dense fine particles can prevent local migration of high angle boundaries and thereby suppress discontinuous recrystallization by Zener pinning of the subgrains under deformation conditions of slower strain rates. However, under the conditions of increased external stress (higher strain rates) and thermal activation, which provide additional driving pressures to overcome the particle pinning, there will be an enhanced stress-assisted and thermally activated boundary migration during hot deformation. In addition to this, the increased subgrain misorientation with strain and strain-induced boundary migration can result in randomization of crystallographic orientation during cDRX process in quasi-single phase Al alloys like in current AA5024 alloy.

4.3. Effect of friction stir welding process on microstructural evolution Malopheyev et al. [38] reported the preservation of fine nano dispersoids of Al3(Sc, Zr) in stir zone of FSW of cold rolled 1570C grade Al-Mg-Sc alloy. Kumar and Mishra [22] reported that Al-Mg-Sc alloy contains both primary and secondary Al3(Sc, Zr) dispersoids. While the primary dispersoids are in micron size range, the secondary dispersoids are in nanometer range. It was observed that the primary dispersoids get refined slightly and the secondary nanodispersoids control grain growth during FSW [22]. According to this, the observed precipitates in NZ of FSW sample of the current study seem to be secondary precipitates.

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The joint formation in FSW occurs by mechanical mixing of severely deformed (plasticised) material on both sides of the faying surfaces. Typically, friction stir welded joint constitutes three important zones, namely nugget zone (NZ), TMAZ and HAZ (heat affected zone). The central region within seam (NZ) undergoes severe plastic deformation at higher strain rates but also experiences peak temperatures and thereby anticipated that dynamic restoration processes can take place during FSW. As a result, the friction stir processed region exhibits fine recrystallized and equiaxed grain structure [21]. While the TMAZ region, which is expected to be present majorly on advancing side only, displays strong inhomogeneity in material flow and can contain a partially recovered/recrystallized microstructure. Whereas, the HAZ represents the transition region between the TMAZ and unaffected BM and it does not experience any mechanical deformation but there will be some thermal gradient. However, the existence of HAZ in FSW Al-Mg-Sc alloys is highly a debatable topic [18]. The heat generation during FSW which softens the base material around the tool, is majorly due to frictional and adiabatic heating [39]. Selection of proper welding parameters, geometry and dimensions of tool, work-piece thickness and thermal conductivities of the materials to be joined play significant role on heat evolution [39]. It is important to note that the amount of heat evolution affects the alloy microstructure including precipitation kinetics and their dissolution at different zones of FSW of Al alloy and subsequently the joint properties. Heat generation during FSW can be estimated indirectly from the heat input per unit length (P, J/mm) using the following simplified equation [40]: ܲ  =  

ଶగேఛ ௩

(2)

Where, N is the tool rotation rate in revolutions per minute, v is the tool traverse speed in mm/minute and τ is the torque in N-m.

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The above equation (2) indicates that for a given torque and rotation rate, the heat input to the material decreases with increasing the tool traverse speed. Besel et al. [20] reported that the maximum temperature achieved at a distance of 7 mm away from the FSW joint line of AlMg-Sc alloy was in the range of 612-650 K and was decreased with increasing the tool traverse speed from 480 to 720 mm/minute. Much higher temperatures (>673 K) can be expected at central region of NZ. Although the exact strain rates involved during FSW are not known, it is estimated to be in the range of 20-350 s-1 [41]. The Zener-Hollomon parameter (Z), which represents the combined effect of strain rate (ߝሶ) and temperature (T), is useful for assessing the deformation behavior of materials and it can be written as: ொ

ܼ  =   ߝሶ ݁‫ ݌ݔ‬ቀ ቁ ோ்

(3)

Where Q is the activation energy of deformation and R is gas constant. Dynamic restoration processes occur typically in a certain range of ߝሶ and T, and hence in a particular range of Z values. Further, under hot deformation conditions higher the Z, higher the flow stress, σ. According to equation (1), higher Z results in finer subgrain size. From hot compression deformation study conducted under quasi-static conditions involving strain rates of 10-3 – 10-1 s-1 (section 4.2), it is very clear that the alloy undergoes dynamic restoration at temperatures above 523 K. For the same Z value range for dynamic restoration of microstructures, the deformation processes involving higher strain rates like FSP require higher temperatures, which can easily be achieved in FSW due to friction and adiabatic heating. Higher heat generation at central region of NZ as well as at lower traverse speed can lead to significant grain coarsening. The observed wavy nature of microstructure (at locations 3 and 5, Fig. 8) could be the resultant of mild stir processing effect arising from the relative motion between tool and material flow due to shear

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deformation but less temperature raise at this location. Development of random texture (at locations 1, 2 & 4, Fig. 8) during FSW is on similar lines as that observed in compression testing (Fig. 7). Similar kind of texture development was reported by Malopheyev et al. [38] upon FSW of the cold rolled 1570C alloy. As explained earlier, higher strain rates and higher temperatures involved in FSW facilitate enhanced stress-assisted and thermally activated boundary migration by easy overcoming of the precipitate barriers. Increased subgrain misorientation due to severe straining in NZ along with enhanced boundary migration can produce highly randomized recrystalized grain structure during FSW. Kumar and Mishra [22] proposed geometric dynamic recrystallization (gDRX) as the possible mechanism for the grain refinement in FSP of similar alloy. It is understood that gDRX is a process of grain impingement during deformation to large strains, where subgrain size remains constant and the prior high angle grain boundaries serrated during deformation come closer. Further, it is important to note that the crystallographic texture remains almost unchanged as there will be limited boundary migration in gDRX process [33]. However, in the present study the observed texture change from strong rolling texture to nearly random orientation during compression deformation (section 4.2) as well as during FSW rules out the gDRX mechanism. As noted and explained in section 4.2, the cDRX process by progressive rotation of subgrains is believed to be the possible mechanism by which fine recrystallized and equiaxed grain structure is produced in NZ region during FSW of AA5024 alloy.

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5. Conclusions

The commercial AA5024 alloy in cold rolled condition was used to investigate the microstructural evolution during thermal treatment, quasi-static compression deformation and friction stir welding. The following conclusions can be drawn from the present study: 1. The as-received alloy in heavily deformed condition containing fine precipitates (<20 nm) of Al3(Sc1-xZrx) exhibited two exothermic peaks at ~713 K and ~813 K and one strong endothermic peak at ~867 K during continuous heating in DSC. The two exothermic peaks are attributed to the concurrent precipitate coarsening and static recovery and extended recovery (recrystallization) processes, respectively, while the endothermic peak to dissolution of the particles and partial melting of the alloy. 2. Upon static annealing, the as-received alloy showed initially continuous recrystallization process which includes recovery and extended recovery. However, at a higher temperature of 873 K, coarse equiaxed grain structure was formed with significant weakening of the initial strong deformation (rolling) texture because of coarsening and dissolution of precipitates. The alloy exhibited reduced hardness and compressive strength but enhanced strain hardening upon annealing. 3. The alloy displayed high positive strain rate sensitivity and low flow stresses at elevated temperatures (above 523 K) in compression, which are characteristics of superplasticity. This was associated with dynamic restoration processes of recovery and recrystallization (extended recovery) and significant change in crystallographic texture to nearly random orientation. 4. Upon friction stir welding of the as-received alloy, the nugget zone displayed fine recrystallized, equiaxed and random orientation grain structure with fine Al-Sc based

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intermetallic precipitates (~10-50 nm). With increasing the tool traverse speed from 250 to 500 mm/minute which leads to decrease in heat input, the alloy exhibited relatively finer grain size in nugget zone of FSW. 5. Based on detailed microstructural investigation of the alloy as a function of temperature and deformation, it is proposed that continuous dynamic recrystallization (cDRX) involving progressive rotation of subgrains is the prominent mechanism for the grain refinement during FSW of Al-Mg-Sc alloy.

Acknowledgements

Authors would like to acknowledge facilities such as National OIM-Texture laboratory, FISTDual beam FIB-SEM laboratory and SAIF at IIT Bombay for providing access to use electron microscopes and FIST-UTM at IIT Bombay for carrying out compression testing.

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Figure Captions Figure 1. Microstructures of the as-received AA5024 alloy: (a) Tri-planar OIM images, (b) (111) pole figure of L-LT, (c) (111) pole figure of L-ST, (d) bright-field and (e) dark-field TEM micrographs with an inset of SAED.

Figure 2. DSC thermograms of the as-received AA5024 alloy sample obtained in (a) first heating (H1), (b) second heating (H2) and (c) the subtracted (H1-H2) curve.

Figure 3. Triplanar OIM microstructures of the AA5024 alloy samples annealed for 24 hours at (a) 623 K, (b) 793 K, (c) 873 K, and (d) secondary electron micrograph in LT-ST plane of the sample annealed at 873 K.

Figure 4. (a) Variation in microhardness and (b) true stress-true strain curves at room temperature of the as-received and annealed AA5024 alloy samples.

Figure 5. True stress-true strain plots of the as-received AA5024 alloy as a function of (a) testing temperature at 1 × 10-3 s-1 and (b) strain rate at 623 K.

Figure 6. (a) Low magnification & (b) higher magnification IPF maps and (c) SEM micrograph the sample deformed in compression at 523 K; and (d) IPF map, (e) FIB (ion-channel contrasting) image and (f) SEM micrograph of the sample deformed at 623 K. The strain rate of compression deformation is 1 × 10-3 s-1. The compression axis and the plane of sample is shown in figure.

Figure 7. (a) Tri-planar IPF maps and (b) FIB image on L-LT plane of the AA5024 alloy sample compressed at 623 K and strain rate of 10-1 s-1.

Figure 8. The IPF maps from different regions of typical butt-joint configuration of FSWAA5024 alloy sample processed at a tool traverse speed of 250 mm/minute. Schematic showing the locations of the selected regions of the FSW sample is presented in figure.

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Figure 9. (111) pole figures and [001] IPFs from different regions of typical butt-joint configuration of FSW-AA 5024 alloy sample processed at a tool traverse speed of 250 mm/minute.

Figure 10. (a) IPF map, (b) (111) pole figure and (c) [001] IPF from center region of the buttjoint configuration of FSW-AA5024 alloy sample processed at tool traverse speed of 500 mm/minute.

Figure 11. Bright-field TEM micrographs at (a) low and (b) higher magnifications, (c) high resolution electron micrograph, and (d) EDS line profile elemental analysis profile obtained from one of the particles of a sample from location-1 (weld nugget portion) of FSW-AA5024 alloy processed at 250 mm/minute traverse speed.

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Declaration of interests ☒ 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. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:

Sumit Chhangani, Suresh Kumar Masa, Rohit T Mathew, M.J.N.V. Prasad, Sujata M