Skip to main content

Severe Plastic Deformation of Al–Mg–Si Alloys Processed Through Rolling Techniques: A Review

Abstract

In the present work, different combinations of rolling used by the researchers as a novel severe plastic deformation technique for the deformation of Al–Mg–Si alloys have been investigated. Reported research work is used for explaining the microstructures obtained after the processing with the help of Electron Back-Scattered Diffraction (EBSD), Transmission Electron Microscopy (TEM) and Differential Scanning Calorimetery (DSC). Based on the literatures investigated, it was found that cryorolling (CR) followed by warm rolling (WR) of Al–Mg–Si alloy provided the highest specific strength compared to all the investigated different combinations of rolling techniques available in the open literatures. The CR (70%) followed by WR (20%) at 145 °C followed by ageing at 125 °C for 60 h provided the highest hardness (130 HV) and ultimate tensile strength (400 MPa) in Al–Mg–Si alloy having a chemical composition of (Si:0.67 Fe:0.28 Cu:0.20 Mn:0.04 Mg:1.01 Cr:0.05 Zn:0.06 Ti:0.01 Al: balance). The CR followed by warm rolling helped in retaining the higher strength as well as higher ductility. It is because of dynamic recovery and precipitation evolution was dominated during the processing. The dynamic recovery was promoted the higher ductility, and precipitation evolution helped in the improvement in the strength due to precipitation strengthening. The peak ageing of warm rolled sample further helped in improvement in the mechanical properties by evolution of β′′-precipitates.

Introduction

Over the decades, researchers used various severe plastic deformation (SPD) techniques for altered the microstructure and to receive the desired mechanical properties [1]. The Al–Mg–Si alloys as a high stacking fault energy material have a greater deformability. The researchers deformed the Al–Mg–Si alloys at various working temperature by using various SPD techniques. They further used different types of ageing, even they combined the ageing phenomenon with the deformation temperature. By doing so, they significantly improved the mechanical properties of Al–Mg–Si alloys. In this work, author tried to combine the contribution of various researchers for deforming the Al–Mg–Si alloys by using various deformation techniques.

Vaseghi et al [2] performed the Equal Chanel Angular Pressing (ECAP) for deforming the Al–Mg–Si alloys. They studied the ageing phenomenon, without ECAP, pre-ECAP, post-ECAP and dynamic ageing (during the ECAP). They found that ageing on the without ECAP processed Al–Mg–Si alloys, and then, this aged Al–Mg–Si alloys followed by the ECAP provided the highest hardness (138 HV). Panigrahi et al. [3] performed the cryorolling on the 6063 Al Alloy up to the thickness reduction of 93%. They found that at the 93% thickness, ultra-fine grains (UFG) were formed in the 6063 Al alloy, and these UFG grains were acted as a high driving force for the precipitates evolution. Agena [4] performed the ECAP on the 6082 Al alloy, and then, tensile test was performed on the samples cut from the ND, TD and RD direction after the ECAP. He exhibits the anisotropy behaviour in flow properties of the ECAP-processed 6082 Al alloy with the help of experiment, computer software and the Ludwik equation. Vendra et al. [5] performed the Multi-Axial Forging (MAF) on 6063 Al alloy at the liquid nitrogen temperature and found that MAF at liquid nitrogen temperature helped in the formation of Nano-size, UFG and coarse grains, in the 6063 Al alloy. Panigrahi et al. [6] performed the cryorolling (CR) followed by ageing treatment on the Al–Mg–Si alloy. They reported that ageing followed by CR helped in increasing the strength and ductility of the Al–Mg–Si alloy, significantly, due to precipitation hardening and grain coarsening, respectively. Kumar et al. [7] performed the Warm Forging (WF) and Room temperature Forging on the 6082 Al alloy and found that after performing the WF at 2 cycles the hardness, tensile strength and J-integral value significantly improved. They reported that this was due to the formation of UFG, Dynamic recrystallized grains (DRX) and complete precipitation of β"-phase. Mirzakhani et al. [8] performed the ECAP and ageing on the Al–Mg–Si alloy individually and found that after ECAP the yield strength (YS) and ultimate tensile strength (UTS) of Al–Mg–Si alloy increased two times and three times, respectively, whereas after naturally ageing mechanical properties remain unchanged, and after artificially ageing at 180 °C, it improves the ductility and reduce the strength slightly. They reported it was due to the dislocations annihilation and grain coarsening. Osterreicher et al. [9] investigated the Al–Mg–Si alloy by performing the cyclic deformation combining with the artificial ageing. They found that at small cyclic deformation after each ten minutes combining with the artificial ageing at 160 ˚C can reduced the ageing time by factor of nine. Kolar et al. [10] studied the effect of pre-deformation before artificial ageing, and deformation accompanied with the artificial ageing on the AA6060 Al alloy. They suggested that deformation followed the artificial ageing provided the peak mechanical properties at the shorter ageing time. Murugesan et al. [11] deformed the AA 6082 alloy by using High Pressure Torque (HPT); they achieved highest hardness after the five turn i.e. 175 HV due to the formation of subgrain of size 200 nm. Fan et al. [12], performed the Hot Gas Forming with synchronous die quenching on the Al–Mg–Si alloy sheet for the formation of complex geometry shapes components. They suggest that by maintaining the cooling rate greater than 50 ˚C/sec by die quenching, the strength of the component matches the T6-Temper condition. Kwon et al. [13] found the optimum condition of the forging for the Al–Mg–Si alloy by using the power dissipation Map. Panigrahi et al. [14] investigated the Al–Mg–Si alloy by deforming it through the CR and then performed the short annealing and ageing on it. They reported that post-CR short annealing (155 °C for 5 min) and then ageing treatment (125 °C for 12 h) were the optimum condition for the formation of UFG-grains to improve in the tensile strength as well as ductility. Panigrahi et al. [15] performed the CR and room temperature rolling (RTR) on the 6063 Al alloy and deformed the sheet at different strain level such as 0.4, 2.3 and 3.8, respectively. They found that recrystallized UFG grains were formed in CR condition after the strain level of 3.8, whereas in RTR condition recrystallized UFG grains were not formed even at 3.8 strain level. Liu et al [16] investigated the deformation twin in the 6082 Al alloy processed through ECAP, which were never reported in the coarse grains. Roven et al. [17] investigated the precipitation behaviour in the Al–Mg–Si alloy produced through the ECAP at room temperature and at 175 °C. They observed the dynamic precipitation during ECAP at both temperatures. Weng et al. [18] investigated the effect of addition of Ag and Cu on the precipitation behaviour of Al–Mg–Si alloy. They reported that Ag/Cu addition enhances the precipitation kinetics during natural ageing and artificial ageing. This was due to the strong interaction between the Cu, Ag and Mg atoms. Das et al. [19] investigated the mechanical properties of 6063 Al alloy using new approach, i.e. instrumented ball indentation technique. They performed the ageing at two parameters, i.e. 220 °C/2 h and 220 °C/7 h. They found the improvement in mechanical properties after ageing at 220 °C/7 h due to formation of Mg2Si-precipitates of size 60 nm. Ye et al. [20] performed the compression test (CT) on the 6063 Al alloy at three conditions, i.e. as cast, solution treated and artificial aged (AA). They performed the CT at a constant strain of 0.2, and strain rate was varying from 1 × 10−3 to 3.5 × 103 s−1. They observed the positive strain rate sensitivity in AA condition during dynamic compression, and sensitivity under quasi-static compression was low. This was due to the presence of high dislocation density and homogenous distribution of β′′-precipitates inside the matrix of AA 6063-Al alloy. Odoh et al. [21], investigate the thermo-mechanical processing of three Al–Mg–Si alloys (effect of Cr-addition and increased Mg-Si content) over the temperature range from 400 to 550 ˚C and the strain rate varying from 0.01 to 10 s−1. They found that addition of 0.2wt.% Cr significantly improved the average steady flow stress between the strain rate of 0.01 and 1 s−1. Further, they developed the model with the help of predicted and experimental stress flow behaviour. Kumar et al. [22] performed the CR followed by annealing at different temperature ranges from 150 to 400 ˚C on the 6082 Al alloy. They reported that CR followed by annealing at 150 ˚C exhibits maximum hardness and tensile strength due to the precipitation of β”-precipitates and formation of the nano-size ultrafine grain. They also reported that high dislocation density inside the matrix of CR-6082 Al alloy helps in producing the dense oxide film on the surface and further helps in improving the corrosion resistance. Rao et al. [23] investigate the effect of CR and Warm rolling (WR) on the precipitation kinetics of the 6061 Al alloy. They observed that peak corresponding to the β’-precipitates was disappeared in the differential scanning (DSC) curve in pre-deformed 6061 Al alloy due to the simultaneously precipitation of the β’ and β’-precipitates. Panigrahi et al. [24] investigated the UFG-grain evolution produced through the CR of the Al–Mg–Si Al alloy and found the strain level of 3.6 during CR was enough to produce the UFG- Al–Mg–Si Al alloy. Yang et al. [25] used the differential speed rolling (DSR) to analyse the effect of deformation path on the texture and mechanical properties of the Al–Mg–Si alloy. They performed the DSR on the samples in three conditions, i.e. no rotation, rotation along the TD and rotation along the RD, with respect to the passage. They found that excellent strength about (~312 MPa) and ductility of (~7.4%) were achieved, due to formation of macro-shear bands. Kumar et al. [26] performed the CR and RTR, followed by annealing on the 6082 Al alloy. They observed that CR exhibited the best results as compared to RTR in terms of mechanical and corrosion properties. They reported that it was due to the suppression of dislocations annihilation and as well as suppression of precipitation during the CR. Kumar et al. [27] performed the CR followed by WR ranges over different temperatures (100–250 ˚C) on the 6082 Al alloy. They observed that CR+WR (100 ˚C) 6082 Al alloy exhibited the highest fatigue strength due to the presence of high dislocation density and small size of the AlFeMnSi-phase & Mg-Si-precipitates. Engler et al. [28] investigated the effect of the natural ageing (NA) and pre- artificial ageing (AA) on the precipitation kinetics of Al–Mg–Si alloy. They observed the volume fraction of the β”-precipitates were much larger in pre-AA sample, whereas the particle size of NA was much stronger than pre-AA. They finally concluded that despite the difference in the precipitation kinetics, both the samples have similar peak strength. Panigrahi et al. [29] compared the result of the CR and RTR, performed on the 6063 Al alloy. They observed that strength and ductility of the severely deformed 6063 Al alloy were increased after the ageing. They found that precipitation hardening was responsible for the strength enhancement, and grain coarsening was responsible for the increase in the ductility. Rao et al. [30], compared the result of the post-cryorolled treatment performed on the Al–Mg–Si alloy. They found that CR+WR sample has the better UTS and hardness value as compared to the CR+Short annealing (SA) sample. They observed it was due to dynamic ageing and formation of the dislocations free cell substructure during the CR+WR sample. Zi et al. [31] investigate the influence of the pre-ageing (PA) time and temperature on the secondary natural ageing (NSA) and paint baking (PB) hardening of the Al–Mg–Si alloy. They observed that PB hardness increased by the increasing PA temperature and NSA stability improved by the increasing in the PA time or decreasing the PA temperature. Panigrahi et al. [32] investigate the 6063 Al alloy produce through the CR and RTR. They found that CR sample has the strength of 257 MPa as compared to the RTR sample (232 MPa). The higher strength of the CR sample was due to the accumulation of the high amount of dislocation density inside the matrix of CR-6063Al alloy, as compared to RTR sample. Lu et al. [33] investigate the evolution of the secondary phase particles in the twin rolled casted (TRC) Al–Mg–Si alloy and relate it with the mechanical properties and recrystallized texture. They observed during the TRC and homogenization, different type of distribution (coarse and fine) of secondary phase particles was achieved. These distributions help in controlling the recrystallization texture intensity and significantly influenced the mechanical properties. Zhao et al. [34] investigated the effect of processing parameter of the friction stir processing (FSP), on the mechanical properties and microstructure evolution of the 6063 Al alloy. They found that density of the needle-shaped precipitates was reduced at the weld nugget (WN). Strength of the base metal is higher than the single or two pass FSP sample, but as the speed of FSP pass was increased, then strength of processed zone was increased as compared to the base metal. Wang et al. [35] investigated the effect of the none-recrystallized, partially recrystallized and fully recrystallized grains on the corrosion fatigue properties of the Al–Mg–Si alloy. They found that low-angle grain boundaries (LAGBs) along with the Mg–Si segregation have the significant role on the corrosion penetration depth, whereas the high-angle grain boundaries (HAGB) and precipitation free zone (PFZ) do not affect the corrosion resistance. Qian et al. [36] investigate the effect of the addition of Mn (0.05–1 wt.%) on the recrystallization resistance of the AA 6082 Al alloy under the condition of deformed as well as post-deformed annealing condition. They found that after the addition of 0.05 wt.% of Mn recovered grain structure remain stable up to 8 h of annealing, whereas in the Mn free base alloy static recrystallization occurs after 2 h of annealing and after 4 h of annealing grain growth started. As they increase the Mn content up to 1 wt.%, static recrystallization resistance was reduced. It was due to the increased Mn content that helps in reducing the dispersoid free zones (DFZ), which favour the recrystallization process. Rao et al. [37], performed the CR+Peak ageing (PA) and CR+WR+Peak ageing (PA) on the 6061 Al alloy. They found that CR+WR+PA sample exhibits the excellent strength (406 MPa) as compared to the CR+PA (365 MPa) sample. This was due the combined effect of the solid solution strengthening, dislocation strengthening, dynamic recovery, and dynamic ageing during CR+WR; further PA contributes in the strength by pinning of dislocations through nano-sized precipitates. Kim et al. [38] investigated the tensile deformation response of the Al–Mg–Si alloy under the condition of solution treatment (ST), natural ageing (NA) and artificial ageing (AA) when the pulse electric current is applied. They observed that as the pulse electric current is applied under the condition of ST, the flow curve and elongation both increased, and Portevin–Le Chatelier (PLC)-effect nearly disappeared, whereas in the case of NA, flow stress decreased and elongation significantly increased, and in the case of AA flow stress and elongation both decreased. They observed that thermal/electric current induced annealing occurred in all the investigated condition. This annealing, in the case of ST sample, led to early stage precipitation, which helped in increasing the strength as well as ductility, whereas in the case of AA sample, micro-voids start forming at the grain boundaries, around the precipitates, which led to the early failure. Kim et al. [39] observed the enhanced ductility in UFG Al–Mg–Si alloy produced through the high ratio differential speed rolling at warm temperature (493–553 K). They found the deformation mechanism for this was the pipe diffusion-controlled grain boundary sliding, and reason for increased in the ductility was the higher strain rate sensitivity, and increased strain hardening was supported by accelerated grain growth. Majchorwicz et al. [40] investigated the 6101 Al alloy produced through the hydrostatic extrusion combined with the artificial ageing. They observed the double fibre (<100> and <111>) texture after the hydrostatic extrusion. The <100> grain provided the precipitation strengthening by the β”-precipitates, and <111> grain provides the grain boundary strengthening. They observed the higher conductivity due to decomposition of the solid solution into the β”-precipitates at the grain interior and spherical β′& β-precipitates located at the grain boundaries. Winter et al. [41] investigated the 6082 Al alloy produced through the ECAP and compared the single pass ECAP and as extruded peak aged condition high cycle fatigue life. They observed the fatigue strength was influenced by the sensitivity to the certain loading direction and during the high cycle fatigue life investigation. Cuellar et al. [42] investigate the wear resistance behaviour of Al–Mg–Si alloy produce through ECAP. They found that wear resistance increased as the ECAP deformation passes increased due to the fragmentation and re-distribution of the secondary phase particles. Gbenebor et al. [43] investigated the effect of extrusion using different die geometry ranges from 15° to 75° on the 6063 Al alloy. They observed that extrusion using 75° die geometry of the 6063 Al alloy results into the outstanding compressive strength, plastic deformation, strain rate and energy absorbing capacity. This was due to the increased in the die angle which led to the decrease in the deformation zone and samples were deformed in the axial direction. Magalhaes et al. [44], performed the CR on the AA6061 Al alloy and then conducted the tensile testing at temperature of 298, 173 and 77 K with varying strain rate of 0.1 to 0.0001 s−1 on the deformed sample. They observed the suppression of PLC effect on the 77 K sample, and at 298 K sample significant PLC-effect was observed; this was due to the decrease in dynamic recovery rate at the low temperature. Yang et al. [45] performed the isothermal hot compression test on the 6A82 Al alloy with temperature ranges from 320 to 530 °C at varying strain rate from 0.001 to 10 s−1. They observed the continuous flow softening behaviour at the temperature range from 320 to 390 °C. This was due to the enhanced dynamic recrystallization and coarsening of the dynamic precipitation. Aruga et al. [46], investigated the cluster classification according to the composition based on the atomic density inside cluster through the high detection efficiency atom probe. It reveals that Si-rich cluster has value (Mg/Mg+Si) less than 0.4 and has smaller atomic density inside cluster. This was led to high vacancy concentration inside Si-rich cluster. Pan et al. [47] performed the grain refinement in the AA6061 Al alloy without the use of plastic deformation or matrix transformation. They preferred the cyclic electro-pulsing treatment, in which dislocations in the [111] slip plane was activated. This led to enhance in the [100]// ED texture as well as increases in the [111] planes as the electro-pulsing treatment repeated. Wang et al. [48] sequentially performed the natural ageing followed by the cold rolling, and then, they finally performed the post-annealing on the Al–Mg–Si alloys. They achieved the significant improvement in the strength and ductility was comparable to T6-treated alloy. During the post-annealing treatment, dislocation density decreases and formation of the subgrain as well nano-size grains was accelerated. This post-annealing accelerated the formation of the wall shape nano-sized precipitates, which was different with the traditional precipitation in the Al–Mg–Si alloy. Winter et al. [49] investigated the fatigue strength of the oxide coated 6082 Al alloy in which they used the ECAP processed substrate. They found that fatigue strength of the ECAP processed substrate was 63% of the none-ECAP uncoated substrate. The depreciation in the fatigue strength of the ECAP-processed substrate was due to the initiation of crack from the oxide coating towards the substrate. Li et al. [50] investigated the hot deformation behaviour of the Al–Mg–Si alloy between the temperature ranges of 400 and 550 °C with varying strain rate of 0.001−1s−1 at true strain of 1.2. They observed that deformed microstructure depends upon the Zener–Holloman value (Z). The value of Z decreased when subgrain size as well as fraction of the recrystallization increased. Lio et al. [51] investigate the hot deformation behaviour of the Al–Mg–Si alloys has different Si wt.% i.e. 0.7 wt.%, 7 wt.% and 12.3 wt.% between the temperature ranges of 573 and 773 K, at varying strain rate of 0.01, 0.1, 1 and 5 s−1 up to the true strain of 0.6. The higher amount of Si containing Al–Mg–Si alloy formed the large fraction of Si-particle; these particles hinder the dislocations motion, which helped in increasing the steady state flow stress as well as deformation activation energy. Hu et al. [52] investigated the Al–Mg–Si alloy produced by the repetitive continuous extrusion forming (R-Conform) followed by the T8 tempering treatment. After this, they observed the excellent strength and enhanced ductility in the Al–Mg–Si alloy. The refined subgrains and induced high dislocation densities during the R-Conform process further contribute in the formation of the large fraction of nano-size β′′-precipitates. Zhang et al. [53] investigated the Al–Mg–Si alloy used for the cable application. They deformed the sample at different strain, i.e. 0.98, 1.73 and 2.24, and after that, they performed the annealing at temperature ranges varying from 200 to 500 °C. They observed that recrystallized Al–Mg–Si alloy exhibits higher electrical conductivity as compared to the non-recrystallized Al–Mg–Si alloys. Further annealing helped in increasing the electrical conductivity due to decrease in the dislocation density as well as grain boundaries area. Gonzalez et al. [54], compared the influence of conventional shoot peening (CSP) and severe shoot peening (SSP) on the fatigue life of 6063 Al alloy. They observed that severe shoot peening was the severe plastic deformation surface modification technique, in which grain of nano-size was formed on the surface. They finally concluded that SSP has the better fatigue life compared to CSP; it was due to the formation of the nano-size thick surface layer and deep compressive residual field. Gao et al. [55] compared the Al–Mg–Si processed at different condition, i.e. solution treatment with hot air and solution treatment with electrical resistance. They observed that strength of the solution treatment with hot air Al–Mg–Si alloy was high due to the formation of the smaller size and more uniform recrystallized grain and has excellent bake hardenability due to the more fraction of clusters formation. Jiang et al. [56] investigated the 6063 Al alloy and 6082 Al alloy at two different heat-treated condition, i. e. underaged (UA) and peak aged (PA). They observed that fatigue crack in UA condition moved in first stage (Faceted form) in the 6063 Al alloy, and in the PA condition crack moved in the second stage (striation form) in the 6063 as well as in the 6082 Al alloy. Cerri et al. [57], performed the ECAP followed by ageing on the Al–Mg–Si alloy. They observed that alloy exhibited the softening after performing the ageing at higher temperature, i.e. 170 ˚C, whereas at the low temperature hardness remains constant due to the precipitation hardening. Hussain et al. [58] performed the CR followed by WR with different temperature ranges from 100 to 175 °C.They found the optimum ageing (12°C for 45 h) while performing the hardness measurement. The strength was significantly improved to 450 MPa at 125 ˚C/45 h due to evolution of the nano-size dense β”-precipitates. Kumar et al. [59] investigated the 6082 Al alloy processed through the CR and RTR at three deformation levels, i.e. 40, 70 and 90%. They observed that CR and RTR both improved the mechanical properties due to high dislocation density, sliding of grain boundary, and grain refinement. But CR exhibited the better mechanical properties as compared to RTR, for the same amount of deformation, because of the formation of the subgrain structure as well as high dislocation density. Bouquerel et al. [60] investigate the T6-6082 Al alloy processed through the artificial ageing before and after the forming process. They observed that intermetallic particles were removed from their locations due to the plastic deformation. Value of plastic deformation less than 0.1 helped in the formation of substructure inside the coarse grain, which further promoting the strength. Bochvar et al. [61] investigated the Al–Mg–Si alloys produced through the ECAP followed by the ageing with the addition of the Sc+Zr and Sc+Hf. They observed that the addition of metals does not affected the precipitation sequence, but precipitation kinetics shifted to lower temperature side. Rao et al. [62] performed the multiaxial forging at liquid nitrogen temperature at three different strain levels, i.e. 1.8, 3.6, and 5.4. They found that hardness and strength increased from 50 to115 Hv, and 180 to 388 MPa, respectively. This was due to the formation of the subgrain structure of the size of 250 nm surrounding by the high-angle grain boundaries. Wan et al. [63] investigated the influence of the welding speed on the microstructure as well as on the mechanical properties processed through the self-support friction stir welding (SSFSW). They observed that transverse tensile strength of the joint increased, whereas ductility decreased with increase in the welding speed (from 10–200 mm/min). Zhang et al. [64], investigated the Al–Si–Mg alloy processed through the high-pressure torsion (HPT) combined with the prior heat treatment. They observed the high fraction of nano-scale Si-particles when undergone through the HPT. These nano-size particles help in increasing the strength up to 550 MPa due to hindrance in the path of the dislocations. Song et al. [65] investigated the two-step ageing on the Al–Mg–Si alloy with the help of 3D-atom probe. They found that at two-step ageing at 170 °C for 0.6 ks, less fraction and smaller size of cluster were formed as compared to natural ageing at 2419.2 ks. It occurred due to the dissolution of the nanocluster occurred during the first stage of the two-step ageing. Wang et al. [66] studied the U-shaped profile of the Al–Mg–Si alloy processed through the hot extrusion, solid solution treatment and ageing. They found the coarse grain structure near to profile edge due to the dynamic recrystallization and excess grain growth during processing. The deformed subgrains were found between the two coarse grain layers due to the work hardening and dynamic recovery. Liu et al. [67] processed the Al–Mg–Si alloy through the subrapid solidification (SRS) technique instead of conventional solidification (CS) process. They observed through SRS technique fine grain structure promoted, high average partition coefficient of solute atoms and weaker micro-segregation, which finally accelerated the high cooling rate (160) as compared to the CS (30 ˚C/s). Fadhalah et al. [68] investigated the 6063 Al alloy processed through the multi-pass friction stir processing at three overlapping (OL) percentages, i.e. 25, 50 and 75. They observed the similar grain structure size (3-6 µm) for all OL percentages samples, but the strength and microhardness were reduced as the OL percentages increased due to the dissolution of the work hardening precipitates. Meng et al. [69] investigate the T6-6082 Al alloy processed through the metal inert gas welding (MIG). They observed the lowest hardness at the centre of the weld (84 HV), and hardness in the softest zone was 100 HV. They observed during the fatigue crack stability propagation stage, non-deformable secondary phase particles promote the crack under the action of dislocations and deformable secondary phase particles hinder the crack propagation. Liu et al. [70] investigated the hot deformation behaviour of the Al–Mg–Si alloy based on the diffusion mechanism. They observed the increases in the low-angle grain boundaries along with the increases in subgrains and recrystallized grains, as the induced strain was increased. The softening mechanism is more prominent due to the dynamic recrystallization as well as due to the dynamic recovery, in which dynamic recovery was the influential. Lu et al. [71] investigated the Al–Mg–Si alloy processed through the double-pass continuous expansion extrusion forming (CEEF) for increasing the rod diameter gradually. They observed the strength and diameter of the CEEF increased significantly due to the grain refinement by the continuous dynamic recrystallization (CDRX) as well as by the geometric dynamic recrystallization (GDRX). Liu et al. [72] investigated the UFG Al–Mg–Si alloy processed through the ECAP combined with the dynamic ageing. They observed that after the three ECAP pass strength (445 MPa) was significantly improved due to the formation of the grain of average size 239 nm. The ageing further contributes in the strength by interaction of the dislocations with the precipitates (nano-size β”-precipitates). They finally concluded that ECAP combined with the ageing provided the advantage to the antifriction performance of Al–Mg–Si alloy. Kim et al. [73] investigated the 6063 Al alloy produced through the mechanical milling followed by the hot extrusion. They observed in this process coarse grain combined with the ultra-fine grain (UFG) was formed. This bi-modal grain structure was help into attain the strength and ductility, due to the presence of UFG-grain and coarse grain, respectively. Lin et al. [74] investigated the welding joint (WJ) and base metal (BM) of the extruded Al–Mg–Si alloy, produced through the metal inert gas welding. They observed the strength of the BM was 76.5% of the BM, and fatigue strength (107- cycles) was for the BM (141.3 MPa) and for the WJ (61.7 MPa), respectively. They found the welding pores, incomplete fusion, and local stress concentration, in the WJ area, were the main responsible for fatigue crack initiation and propagation and further deteriorated the WJ fatigue life. Das et al. [75] investigated the 6063 Al alloy processed through the high-pressure torsion (HPT). They observed during the initial stage of straining, the subgrain was formed, and each subgrain has the dislocations inside them. As the strain induced increases, the size of the subgrain surrounding by the high-angle boundaries (HAGBs) decreased along with the increased in dislocations inside the subgrains. They observed that during straining Vickers hardness, yield strength and UTS increased simultaneously and reached to saturated level, and beyond this level mechanical properties remain unchanged in the steady state. Puplampu et al. [76] investigated the effect of the fire exposure on the AA 6061 Al alloy combined with the tensile, torsion, and combined loading. They observed the 60–70% strength reduction in the strength after the fire exposure. The combined loading helped in achieving the high strength as compared to the pure tensile loading under the exposure of fire. The microstructural investigation reveals that microstructure was relief in the fire exposed sample, and mechanical loading leads to the strain-induced grain refinement. Sunde et al. [77] examine the Al–Mg–Si alloy for improving the electrical conductivities. For this, they selected two conditions: (1) Conventional solution treatment, quenching, and artificial aging at 170°C, (2) Solution treatment, quenching, pre-ageing, 50% cold rolling reduction, and re-ageing at 170 °C. They found that in second condition sample, electrical conductivity was significantly improved. It was due to the solute depletion in the Al-matrix. Kumar et al. [78] investigated the plane stress fracture toughness of the 6082 Al alloy processed through CR, RTR followed by annealing at 200 °C. They observed the plane stress fracture toughness using load point displacement (LPD)-load and the J-Integral resistance (J-R) curve, and they found that J-R curve method was more accurate for evaluating the critical plane stress fracture toughness of the processed 6082 Al alloy. They found the high plane stress fracture toughness for the CR+AN sample was due to the presence of trans-granular facets, flat circular and elongated dimples, small-sized Si-rich precipitates and high kernel average misorientation. Yang et al. [79] investigated the reason of why the β′′-precipitates are a major strengthening phase, as compared to the β′-precipitates. For this, they prepared the two samples (1) underaged (only have β”-precipitates) and (2) overaged (approximately equal volume of β” and β’-precipitates). They found the reason for the major strengthening contribution from the β”-precipitates was due to the smaller lattice strain (0.99%) relative to the matrix for the β’-precipitates, as compared to the β”-precipitates (2.1%). Zheng et al. [80] investigated the 6082 Al alloy produced through the hot deformation and observed the microstructure during processing using Electron Back-Scattered Diffraction (EBSD) first time. They found that geometrically necessary dislocations (GND) were increased with the induced strain levels and distributed over the large area. Ma et al. [81] investigated the behaviour of T6-6082 Al alloy under the applied axial compression load at different five temperatures conditions. They found that flexure buckling occurs in all the samples under the applied load. They developed the model with the help of the experimental results and verified it with the test results, Chinese code (GB) and European Code (EC9). Dadbakhsh et al. [82] studied the ECAP-processed 6082 Al alloy combined with the ageing treatment performed before and after the ECAP. They found that maximum strengthening was given in the post-ECAP ageing condition, and the ageing temperature used in this condition was low compared to without ECAP ageing condition. Lai et al. [83] investigated the dislocations-induced precipitation behaviour of the Al–Mg–Si alloy. They found that dislocations induced precipitates consumed more Mg as compared to Si; then, it has more formation kinetics. They found that different kinds of the dislocations-induced precipitates were formed such as short-range ordered, long-range disordered polycrystalline precipitates and multi-phase composites precipitates. Wang et al. [84] investigate the Al–Mg–Si alloys processed through the electric pulse treatment (EPT). They found the EPT process triggered the precipitation of the Al5FeSi-phase. This phase pinned the dislocations effectively such that dislocations were remain after the aging, and these remained dislocations were promoted the deformation band, and these band convert the Mg2Si rod-like structure to globular. Shi et al. [85] investigated the cyclic deformation behaviour of the solution-treated AA6061 Al alloy. They found that after the cyclic deformation numerous clusters were evolved and further aging at 177 °C for 2 h, these clusters were transformed into the high number of precipitates. At this ageing condition, AA6061 Al alloy has the 10% more strength compared to the T6-treated condition. Teichmann et al. [86] investigated the AA6060 Al alloy processed through the deformation along with the heat treatment and compared it with the similar heat treatment condition without deformed and pre-deformed conditions. They observed that during the deformation along the heat treatment, dispersion of fine β″-precipitations takes place along the dislocations. In this condition, strength contribution from the β′′-precipitates was more compared to the dislocations strengthening. Saito et al. [87] investigated the effect of the quenching rate after the solution treatment combined with the 1% pre-deformation on the three Al–Mg–Si alloy. They found that Cu addition up to the 0.01% does not affect the precipitation behaviour, whereas when the Cu addition was 1%; then, in this condition fine dispersion of the short precipitates significantly improves the hardness. Kumar et al. [88] investigated the low cycle fatigue behaviour of the 6082 Al alloy processed through the CR, RTR, CR+Annealing and RTR+Annealing. They found that CR+Annealing exhibits the high low cycle fatigue life due to the presence of small sized Si-rich precipitates, secondary phase particles, subgrains formation (200–400 nm) and high stored energy (32.16 × 10−4 J/mol). Rayes et al. [89] investigated the 6082 Al alloy processed through the friction stir processing (FSP) from one to three passes at constant rotational speed accompanying with the three different transverse speed. As the number of passes was increased, it led to softening due to the increases in the grain size and precipitates dissolutions, whereas when the traverse speed increased, in that case strength was increased due to the reduction in the particle size as well in the particle area fraction. Kim et al. [90] investigated the Al–Mg–Si alloy processed through the step quenching (SQ) combined with the two-/multi-step ageing and through the bake hardening. They used the atom probe tomography (APT) and positron annihilation life time spectroscopy (PALS) to investigate the Nano cluster and vacancy formation. They found that SQ at the 100°C for 3.6 ks was more effective for suppressing the negative effects of natural ageing and promotes the precipitation behaviour. Biradar et al. [91] investigate the 6063 Al alloy processed through the accumulated roll bonding (ARB) at room temperature and from 423K to 773K. They found the superior mechanical properties (Hardness: 71 HV, UTS: 178 MPa) at 723 K, but the ductility reduces from 38.3 to 6.6%. They found the optimized mechanical properties (UTS: 165 MPa, Ductility: 15.5%) after the solution treatment at 523 K for 2 h of the ARB processed samples. Birol [92] performed the thixoforging on the wrought 6082 Al alloy. For this, 6082 billets have transformed into non-dendritic after the isothermal heating. After that, slug machined from the 6082 Al alloy was thixoforged after the isothermally heating at 645 ˚C for 5 min. Finally, after the T6 tempered condition, hardness of the thixoforged part increased to 95 \(\underset{\_}{+}\) 3 HB. Mao et al. [93] investigated the β”-precipitation behaviour of the Al–Mg–Si alloy using the experiment and through the modified multi-phase method (MPF) combined with the computer coupling of phase diagrams and thermochemistry. They concern about the interfacial energy anisotropy and elastic interaction, to describe the morphology of the β″-precipitates. Nikhil Kumar [94] investigated the microstructural as well as mechanical properties inhomogeneity in the 6082 Al alloy processed through the open die multi-axial forging (MAF). They found that hardness and tensile strength were varied from the surface to the centre from 100 HV to 120, 300 MPa to 330, respectively. He found the reason for this the strain gradient exit (high strain accumulation at the centre as compared to the surface) in the MAF-processed 6082 Al alloy. Kumar et al. [95] investigated the fracture toughness of the 6082 Al alloy processed through the CR, RTR, CR+AN and RTR+AN. They measure the relative improvement in the fracture toughness by calculating the equivalent energy fracture toughness (Kee) and J‐integral due to the inappropriate sample size available for the testing. They found that CR+AN has the highest fracture toughness as compared to all other investigated samples, because of the presence of Si-rich particles and small size AlFeMnSi- phases. Kumar et al. [96] investigated the mechanical properties of the 6082 Al alloy processed through the CR, RTR, followed by the annealing. They found that hardness and strength of the CR+AN sample were increased up to the 95 and 40%, respectively. The reason for this is the dislocation as well as precipitation strengthening. Kumar et al. [97] investigated the influence of the Mg/Si-precipitates on the hot compression behaviour of the 6082 Al alloy. They observed the influence of the strain rate was not tandem with the conventional results because of the dynamic strain ageing produced by the Mg/Si-precipitates during the compression test. Yuan et al. [98] investigated the precipitates morphology of the evolved precipitates and their correlation between the strength, ductility and toughness of the Al–Mg–Si alloy. They found during the investigation that as the morphology of the precipitates changes from the spherical to rod-/needle-shaped, the ductility of alloy was increased, but at the same time yield strength and fracture toughness were decreased. They found the strength–ductility relationship was conventional type, but the strength–toughness relationship was different from the conventional relationship. They found the reason for this precipitate’s interaction with the dislocations and simultaneously deformation disagreement of the precipitates and matrix. Yassar et al. [99] investigated the effect of the pre-deformation on the evolution of the β″ & β′ precipitates and Q′ phase in the Al–Mg–Si alloy. They found that β″-precipitates were replaced by the β′-precipitates and Q′-phase in the pre-deformed Al–Mg–Si alloy. They found that at the high temperature only Q′-phase exits. Lloyd et al. [100] investigated the surface plastic deformation developed over the different length scale due to deformation. They found that surface displacement occurs at grain boundaries due to bending, and it increased with the bending angle. This was due to the surface topography and plastic deformation near to grain boundaries. Yucel Birol [101] investigated the bake hardening response of the twin-roll cast Al–Mg–Si alloy. He observed that pre-straining improves the hardness up to 90 HV as compared to without pre-straining (72 HV). They observed that just after the solution treatment pre-straining helps in suppressing the natural aging and promotes the precipitation of the β″-precipitates. They observed this was due to the pre-straining induced dislocations, which fill the vacancies introduce due to quenching, and then, it avoids the cluster formation. These dislocations provide the site for the G-P zones, and these zones faster grow to β″-precipitates. Chapelle [102] investigated the cube recrystallization texture developed by the deformation at 400°C through the channel die compression test followed by the annealing at 510°C. They observed the low volume fraction of the cube texture before and after the deformation. Sabirov et al. [103] investigated the UFG Al alloy under the compression test at different strain rates. They found that mechanism of the plastic deformation strongly depends upon the base line strain rate. They further observed that decrease in the strain rate results into the activation of the micro-shear banding due to the grain boundary sliding. Huis et al. [104] investigated the large number of the early structure of the MgxSiyAlz-type precipitates, which have face-centred cubic-based structure. They found that if the Mg:Si ratio is less than one, then needle shape precipitation was favoured, in which Si-based pillars extend in the needle direction. If the Mg:Si ratio is greater than one, then platelet type precipitates were favoured, which have the Mg, Si and Al stacking layers. Pogatscher et al. [105] investigated the artificial aging behaviour of the 6061 Al alloy between the temperature ranges of 150 and 250°C using the atom probe tomography. They observed that hardening response of the alloy after the artificial ageing below the 210°C was reduced with prior natural ageing, and it improves above the 210. They concluded that artificial aging mechanism is controlled by the concentration of mobile vacancies. They used the temperature dependent dissolution of co-clusters and solute vacancy interactions to determine the mobile vacancy concentration. Ninive et al. [106] investigated the β′′-precipitates in the Al–Mg–Si alloy in detail. They found the composition of the β″-precipitate was Mg4Al3Si4 and has the possibility in changing the composition variation in single precipitates. They finally observed that particles size was larger than the directly observed size due to the coherency strain field. Starink et al. [107] investigated the co-cluster formation and their correlation with the strength using the 3D atom probe tomography in the Al–Mg–Si alloy. They found that short-range order strengthening due to the co-clusters was the main strengthening mechanism in this alloy, whereas Si-cluster little bit contributes to the strength due to their low enthalpy. Sha et al. [108] investigated the Al–Mg–Si alloy processed through the high-pressure torsion (HPT). They found the highest yield strength (475 MPa) after the tent revolution. They observed the small size grains after the processing at 100°C due to the strong solute segregation to grain boundaries and processing at 180°C led to the formation of coarser precipitates. They found the highest strength achieved after the ten-revolution was due to the high dislocation’s density, solute nanostructures and grain refinement. Du et al. [109] investigated the overageing which results into several stoichiometric particles in the Al–Mg–Si alloy using the CALPHAD-coupled Kampmann-Wagner Numerical (KWN) framework. Their model can be applied to the overageing condition, i.e. growth of post-β″-phases at the cost of the needle-shaped β′′-phases. Sauvage et al. [110] investigated the nano-scale feature such as grain size, solute content of matrix and precipitates size with the help of the atom probe tomography and transmission electron microscopy. They used the severe plastic deformation technique and post-processing heat treatment for obtaining the nano-scale features. They observed the deformation-induced dislocations and grain boundary segregation during the plastic deformation. Further, they observed the grain growth and precipitation during the post-processing heat treatment. Chorminski et al. [111] investigated the effect of the hydrostatic extrusion on the microstructural features of the Al–Mg–Si alloy. They found the <001> oriented grains with low-angle grain boundaries and <111> oriented grain with the high-angle grain boundaries after the processing. They found the precipitation sequence of coarse grain material was different for ultrafine grained material in the present case. They found the reason for this; the thermal shocks due to deformation led to the formation of the Mg–Si co-clusters, and then, Cu was moved into the Mg–Si co-clusters, which further hinder the growth of the Cu-free β″-phases. Zhu et al. [112] investigated the solute clustering in the Al–Mg–Si alloy. They remove the negative influence of the pre-aging on the mechanical properties, which generally occurred after the solution treatment. For this, they used the thermo-mechanical pre-aging processed by using coil cooling step for controlling the state of clustering. Mckenzie et al. [113] investigated the wrought 6016 Al alloy processed through the ECAP. They developed the 3 D two-phase composite model which was based on the dislocation density, and this model was accurate for predicting the cell size of the ultrafine grain material and also to provide the information regarding the dislocation formation. They found that the effect of the back-pressure results into the increase in the dislocation density inside the cell as well as at the cell boundaries. Jana et al. [114] investigated the fatigue properties of the Al–Si–Mg hypoeutectic alloy processed through the friction stir processing (FSP). They found that FSP-processed alloy has the five times the fatigue strength as compared to the cast alloy. They found the reason for this was the closure of cast porosities, which become crack initiation site in the cast alloy. Zhang et al. [115] investigated the Portevin–Le Chatelier effect in the Al–Mg–Si alloy by using the model based on the dynamic strain aging combined with the three-dimensional finite element analysis. They used the flat and round specimens for the investigation, and they observed that localized deformation zone in the round specimen has diffused double cone morphology, whereas for the flat specimen the propagative deformation bands have planar morphology and inclined tensile axis by 35°. Giofre et al. [116] investigated the meta-stable β″-phases formed in the 6xxx all alloy. They investigate the bulk, elastic strain and interface energies to understand the stability of a nucleating cluster. They used the mesoscale models to analyse the stability and interaction of precipitates based on the classical semi-quantitative nucleation theory. Mariani et al. [117] welded the foils of 6061 O Al alloy, to produce the monolithic and Si–C fibre-embedded samples through the ultrasonic consolidation (UC) process at the condition of contact pressure from 135 to 175 MPa, 20kHz frequency, 50% oscillation amplitude, 34.5 mm s−1 sonotrode velocity and 20°C. They found that deformation occurred by the non-steady-state dislocation glide and dynamic recovery occurred at upper and lower foils. They observed the friction at the welding interface led to temperature varied from 0.5 to 0.8 of melting temperature, high contact pressure and fast strain rates. This led to the microstructure variation by the continuous dynamic recrystallization (CDRX) with in the welding area. Shimizu et al. [118] investigated the 6061 Al alloy produced through the high-power ultrasonic additive manufacturing to understand the microstructure of the weld interface between the two metal layers. They found that microstructure of the weld interface consists the fine grain recrystallized structure results due to the shear deformation along the ultrasonically vibrating direction of the tape surface. Mckenzie et al. [119] investigated the effect of the ECAP on the microstructure evolution of the 6016 Al alloy and performed various heat treatments to understand the workability required to survive the severe plastic deformation. They found that F, W and T4 temper has the limited workability, whereas O and T7 can be pressed to high strain without failure. Woo et al. [120] investigated the 6061-T6 Al alloy processed through the friction stir processing (FSP). They prepared the two conditions for studying the texture evolution: (1) severe plastic deformation and friction heating; (2) alone friction heating. They compared the texture evolved in these two conditions and found that case 1 affects the texture evolution, while case 2 has little effect. Hannard et al. [121] investigated the three 6xxx Al alloy. They observed during the tensile testing that fracture strain was decreased as the tensile strength increased, but major difference in the investigated samples was the area of fracture. They used the 3D microtomography to explain the area of fracture by mechanism of the void nucleation, growth and coalescence process. Mckenzie et al. [122] investigated the simultaneously increase in the tensile strength and ductility of the O-temper 6016 Al alloy with or without back pressure, processed through the ECAP. They observed the 100% ductility after the heat treatment at 200 °C and strain rate of 10−4 s−1. Simar et al. [123] developed the microstructure-based strain hardening model for the precipitates hardening material for studied the supersaturated solid solution and relate the dynamic recovery with the yield strength, and presence & stability of orowan loops. Their model can study the (1) precipitates size, volume fraction and thermal cycles and (2) predict yield strength as function of microstructure. Sritharan et, al. [124] performed the interpreted tensile testing on the extruded 6061 Al alloy after the solution treatment, underaged, peak-aged and overaged conditions. They obtained the hysteresis loop after the interrupted tensile test. They found that width of the loop was increased as the pre-strain increased together with up to the peak aged condition, and at overaged condition it start decreasing. They found that back stresses were start decreasing at overaged condition due to the coarsening of precipitates, which led to the reduction in the dislocations pile-up. Shankar et al. [125] investigated the 6061 Al alloy processed through the ECAP at elevating temperature by the machining. The machining of the peak aged condition gives the chips of the fine microstructure as well as high hardness, as compared to the peak aged condition achieved through the ECAP at high operating temperature. They found the chips that cut from large strains were less stable during the annealing as compared to less-strains counterpart. They observed that initial rapid grain growth provides the stable microstructure as compared to the grain growth due to the prolonged annealing.

Based on the literatures, authors found that in all severe plastic deformation techniques rolling is widely used in the industries for producing the continuous long product such as sheets. In the present work, authors made an attempt to pile up the finding based on development of the ultra-fine grained (UFG) 6xxx-Al alloy by using rolling as a severe plastic deformation technique. Different types of rolling combinations used for the production of UFG-6xxx Al alloy were discussed. Various microstructural features of UFG-6xxx Al alloy were discussed with the help of High-Resolution (HR)-Transmission Electron Microscopy (TEM), TEM and Electron Back-Scattered Diffraction (EBSD). The mechanisms behind improvement in the mechanical properties of UFG-6xxx Al alloy were discussed in detail. The following are the outline of the present work:

  1. 1

    Different combination of rolling

  2. 2

    Microstructural Study

    1. 2.1

      Electron Back Scattered Diffraction (EBSD)

    2. 2.2.

      Transmission Electron Microscopy (TEM)

    3. 2.3.

      Differential Scanning Calorimeter (DSC)

  3. 3

    Mechanical Properties

Different Combination of Rolling

Panigrahi et al. [37] performed the rolling at liquid nitrogen temperature named as cryorolling. They maintained the liquid nitrogen temperature during the cryorolling (CR) by dipping the sample in liquid nitrogen temperature for 10 min initially and 5 min after each pass. Then, this CR sample was used for the ageing at 125 °C. They further performed the combination of CR followed by the warm rolling (WR) at 145 °C named as CR+WR. Then, this CR+WR sample was then used for the ageing at 100 –145 °C. The schematic of the experimental procedure is shown in Fig. 1. In our earlier work [22], they used the 45*30*10 mm3 sample size for the solution treatment (ST) at 540 °C for 24 h followed by the quenching at room temperature. We performed the cryorolling by dipping the sample in liquid nitrogen temperature for the 15 min before processing. Rao et al. [30] used the sample of size 10*30*40 mm3 for the solution treatment at 520 °C for 2 h in muffle furnace followed by quenching at room temperature for the cryorolling. They dipped the sample in cryocontainer for 10 min initially and then 3−5 min after the passes. They deformed the sample up to 4% in each pass and pour the liquid nitrogen on roller to suppress the heat generation. They further performed the cryorolling followed by warm rolling (WR) at 200 °C. They dipped the sample in oil bath furnace for the 5 min during the WR. Different sample conditions investigated by them are shown in Table 1 In one of our earlier study [26], we performed the CR and room temperature rolling (RTR) by dipping the sample at liquid nitrogen temperature and cold water, respectively, before as well as during the processing.

Fig. 1
figure 1

Process diagram for (a) post-CR aged sample, (b) post-CR+WR aged sample [37]. Reprinted with permission from Elsevier.

Table 1 Sample naming and their corresponding processing conditions [30]. Reprinted with permission from Elsevier.

Lu et al. [33] performed the twin-roll casting for the fabrication of 6-mm-thick strip using the alloy melt (Fig. 2). They further homogenized the sample at 550 °C for 0, 5 and 16 h. These homogenized samples were used for the cold rolling from 6 mm to 1 mm. These different conditions helped in studying the effect of secondary phase particles on microstructural evolution and mechanical anisotropy. After that, they performed the T4P treatment (solution treatment at 560 °C for 15 min in hot blast stove followed by pre-ageing at 100 °C for 5 h, and then, finally natural ageing (NA) was performed for 14 days). After the T4P treatment, bake hardening process was simulated on the sample by heating at 180 for 30 ˚C min.

Fig. 2
figure 2

Schematic diagram of the short processes used to prepare the TRC Al–Mg–Si alloy, including TRC, homogenization, cold rolling, and T4P+BH heat treatments [33]. Reprinted with permission from Elsevier.

Yang et al. [25] performed the differential speed rolling (DSR) by using the two same size rollers, but the roller speed ratio was 1:4 (lower roll/upper roll), where the speed of the lower rolls was fixed to ~ 5 m/min. They used the 4-pass during the DSR to achieve 75% thickness reduction. Bridar et al. [90] used 100*40*1 mm3 samples for the roll bonded after homogenization at 723 K for 1 h. Figure 3 shows different processes for the roll bonded. The surface of the samples was prepared by removing the contaminations from the surface by decreasing in acetone bath followed by scratching with the help of steel brush of 0.28 mm diameter (Fig. 4).

Fig. 3
figure 3

Schematic diagram showing the roll bonding process [90]. Reprinted with permission from Elsevier.

Fig. 4
figure 4

(a) Scratch brushing of Al-6063 strips. (b, c, d) Photographs showing bond failures. (e, f, g) Successful bonding of strips [90]. Reprinted with permission from Elsevier.

Microstructural Study

In this section, author investigated the microstructure obtained after different combinations of rolling with the help of HR-TEM/TEM, EBSD and DSC.

Electron Back-Scattered Diffraction

The EBSD micrograph of ST and WR 6061 Al alloy is shown in Fig. 1 [37]. The elongated grain of 45 µm size after solution treatment is shown in Fig. 5(a). Figure 5(b) depicts the orientation image micrograph of the CR+WR condition in which grey lines depict the low-angle grain boundaries, whereas black line depicts the high-angle grain boundaries. Figure 5(c) depicts the image quality (IQ)-map of CR+WR sample which was superimposed with the grain boundaries. In the IQ-map, red lines depict the low-angle grain boundaries (\({5}^{^\circ }\le \theta \le {15}^{^\circ }\)) and green lines depict the partially transformed boundaries (\({1}^{^\circ }\le \theta \le {5}^{^\circ }\)).

Fig. 5
figure 5

EBSD micrographs of (a) ST, (b) CR+WR, (c) Image Quality map with low-angle and high-angle grain boundaries of CR+WR sample (d) IPF colour code [37]. Reprinted with permission from Elsevier.

The misorientation distribution of different processed conditions such as ST, CR and CR+WR sample is shown in Fig. 6 [37]. They observed the high fraction of low-angle grain boundaries was found for CR and CR+WR sample. In the CR sample, high fraction of low-angle grain boundaries was formed due to the fragmentation of grain, whereas fraction of low-angle grain boundaries was decreased for CR + WR sample as compared to CR sample. The decrease in the low-angle grain boundaries in CR+ WR sample was due to formation of subgrains, which led to transformation of low-angle grain boundaries in to high-angle grain boundaries.

Fig. 6
figure 6

Grain boundary misorientation angle distribution of ST, CR, CR+WR samples [37]. Reprinted with permission from Elsevier.

The microstructure and corresponding texture of three different conditions, i.e. H0, H5 and H16, are shown in Fig. 7 [33]. The presence of large number of low-angle grain boundaries in H0 condition microstructure depicts the severe deformation in this condition. The texture component present mainly in the H0 condition was shear texture (R-Cube {001} <110> and γ fibre <111> // ND orientations) and rolling texture (Cu {112} <111> orientation) (Fig. 1(b)). As compared to H0 condition, when the homogenization time was increased, then recrystallization and grain growth take place. Due to this orientation of cube {001} <100> texture was developed as homogenization time increased, which led to development of abnormal grain growth (Fig. 7(e)).

Fig. 7
figure 7

Inverse pole figure (IPFs) and corresponding orientation distribution function (ODF) maps showing the microstructure and texture in (a, b) H0, (c, d) H5, and (e, f) H16. In IPFs, the white and black lines represent small angle grain boundaries (2 °–15 °) and large angle grain boundaries (>15°), respectively [33]. Reprinted with permission from Elsevier.

The (001) pole figure and inverse pole figure show the grain size distribution of the homogenized samples (Fig. 8) [33]. Figure 8 shows that as the homogenization temperature increased, the aspect ratio decreased from 1.99 to 1.81. This depicts the formation of fine equiaxed grain as the homogenization time was increased from 0 to 16 h. However, the solution-treated sample exhibits low intensity of texture, but as the homogenization time was increased, intensity of texture increased corresponding to them.

Fig. 8
figure 8

Inverse pole figure, statistical grain size distribution, and {001} pole figures in the solution treated sheets: (a, d, g) H0 and (b, e, h) H5, and (c, f, i) H16, respectively [33]. Reprinted with permission from Elsevier.

The ODF-map of three different conditions, i.e. H0, H5 and H16, is shown in Fig. 9 [33]. Figure 9(a) shows that in H0-sample, the main texture component is cube-ND {001} <130> orientation. It was observed that as the homogenization time was increased, intensity of the cube-ND texture was weakened. In the H5 condition, Cube-ND and Cube both textures are present (Fig. 9(b)). In the H16-sample, the Cube-ND texture was completely disappeared (due to suppression of the particle-stimulated nucleation (PSN) mechanism), and some fraction of R-Cube {001} <100> texture was appeared (Fig. 9(c)). The transformation from cube ND to cube texture was observed as the homogenization time was increased. This was due to the variance in the presence of the second phase particles as the homogenization time was increased [33].

Fig. 9
figure 9

ODF-maps of the solution treated sheets: (a) H0, (b) H5 and (c) H16 [33]. Reprinted with permission from Elsevier.

The deformed microstructure after CR and RTR after 92% thickness reduction are shown in Fig. 10. The black and white lines depict the high-angle grain boundaries (\(\ge 15^\circ \)) and low-angle grain boundaries (\(1.5^\circ \le \theta \le 15^\circ )\), respectively. It can be observed from the EBSD map that width of the subgrain of CR sample is smaller than the RTR sample.

Fig. 10
figure 10

Microstructures of CR and RTR 6063 Al alloy after 92% thickness reduction: (a) EBSD micrograph of CR material; (b) EBSD micrograph of RTR material [32]. Reprinted with permission from Elsevier.

The subgrain misorientation profile of the CR and RTR sample after the 30 and 92% thickness reduction is shown in Fig. 11. It can be observed from the figure that after the 30% thickness reduction, there is significant improvement in the fraction of low-angle grain boundaries in both rolling temperatures. As the thickness reduction was increased from 30 to 92%, fraction of high-angle grain boundaries was increased in both rolling conditions. Figure 11 shows that fraction of low-angle grain boundaries in CR conditions was higher as compared to RTR conditions after the both thickness reduction.

Fig. 11
figure 11

Frequency histograms of misorientation angles of CR and RTR 6063 Al alloy after different percentage of thickness reductions: (a) CR material at 30% reduction; (b) RTR material at 30% reduction; (c) CR material at 92% reduction; (d) RTR material at 92% reduction [32]. Reprinted with permission from Elsevier.

The EBSD map after the solution treatment of 6063 Al alloy depicts the average grain size of 35 µm as shown in Fig. 12(a) [29]. The EBSD micrograph of state 1 (T4 treatment+CR), state 2 (45 min at 520 ˚C+water quenched+CR) and sate 3 (2 h at 420 °C+Furnace cooled+CR) is shown in Fig. 12(b-d). In all three states, severely fragmented and elongated grains containing substructure surrounding by the low-angle grain boundaries along the rolling direction were observed.

Fig. 12
figure 12

EBSD micrographs of original and deformed Al 6063 alloy after different states (arrow indicates rolling direction) (a) starting solution treated material (b) state 1 (c) state 2 and (d) state 3 [29]. Reprinted with permission from Elsevier.

Figure 13 depicts the histogram of misorientation between the grain boundaries of three different states 1–3 [29]. This diagram depicts the medium high-angle grain boundaries for the state-1, and state-3 has the high fraction of low-angle grain boundaries.

Fig. 13
figure 13

Histograms showing the fraction of boundaries vs. misorientation angle (0) of the deformed Al 6063 alloy after state 1, state 2 and state 3, (A) fraction of low-angle grain boundaries (less than 10° misorientation), (B) fraction of medium angle grain boundaries (10–30° misorientation), (C) fraction of high-angle grain boundaries (more than 30˚ misorientation) [29]. Reprinted with permission from Elsevier.

Figure 14 represents the IQ-maps superimposed with the grain boundary misorientations of the CR+WR 6082 Al alloy at different warm rolling temperatures (100 °C–250). It shows that as the warm rolling temperature increased from 100 °C to 250, the volume fraction of the very low-angle grain boundaries (\(1.5^\circ \le \theta \le 3^\circ )\) was decreased. It was because the increase in rolling temperature led to annihilation of dislocations, and fraction of high-angle grain boundaries were increased due to recovery, recrystallization and grain growth, as the warm rolling temperature was increased.

Fig. 14
figure 14

Image quality map (IQ) superimposed with rotational angle (GR) after WR at (a) 100 °C, (b) 150, (c) 200 and (d) 250, respectively [27]. Reprinted with permission from Elsevier.

Figures 15and 16 represent the IPF-maps and partition of IPF-maps of CR+WR 6082 Al alloy rolled at 100 to 250°C. The partition of IPF-maps was done by using the criteria grain orientation spread (GOS) <1° for identified the recovered grain in different investigated conditions after the fatigue testing. The fraction of recovered grains was increased as the WR temperature increased from 100 to 250 °C. It was due to the decrease in fraction of very low-angle grain boundaries as the rolling temperature increased in the starting conditions.

Fig. 15
figure 15

IPF image of WR at (a) 100 °C, (b) 150, (c) 200 and (d) 250, respectively [27]. Reprinted with permission from Elsevier.

Fig. 16
figure 16

Partition of IPF image (PIPF) by giving condition GOS b 1° WR at (a) 100 (b) 150, (c) 200 and (d) 250°C, respectively [27]. Reprinted with permission from Elsevier.

Figures 17 and 18 represent the IPF-maps and partition of IPF-maps after the fatigue testing of CR+WR 6082 Al alloy. The partition of IPF maps was performed to identified the grain growth due to the cyclic loading by using the criteria (GOS<1˚ and grain size>1 µm). The excessive grain growth occurred during the cyclic loading of CR+WR sample after rolling at 250 °C, due to less stored energy in the starting condition.

Fig. 17
figure 17

IPF image of WR at (a) 100 °C, (b) 150, (c) 200 and (d) 250, respectively, after fatigue testing [27]. Reprinted with permission from Elsevier.

Fig. 18
figure 18

Partition of IPF image by giving condition GOS >1° & grain size >1 μm of WR at (a) 100, (b) 150, (c) 200 and (d) 250°C, respectively, after fatigue testing [27]. Reprinted with permission from Elsevier.

The IPF-maps of Al–Mg–Si alloy after the processing through the differential speed rolling (DSR) are shown in Fig. 19 [25]. The IPF-maps obtained at three different conditions depending upon the deformation path, i. e. no rotation (NR), rotation along the transverse axis (RT) and rotation along the rolling axis (RR). It is understood from the IPF-maps that lamellar-like structure with varying thickness from 3 to 10 µm was formed in the NR and RT sample. It was observed that most of the lamellar structures remain intact after the first DSR pass as well shear deformation localized on the same plane. The finer grain refinement is achieved after the RR as compared to RTR and RT. It was because first macroshear band generated in first pass of DSR was crossed by the second macroshear band in second pass of DSR.

Fig. 19
figure 19

IPF maps of the DSR-deformed samples via (a) NR (b) RT (c) RR methods [25]. Reprinted with permission from Elsevier.

The ODF-maps are drawn in the Euler angle range using the non-orthonormal symmetry at specific section of \(\varphi =0^\circ , 45^\circ , 65^\circ \) as shown in Fig. 20 [25]. The texture component is observed in FCC metals: (i) plain strain texture component such as brass, copper and S, (ii) shear texture component such as E, F and H, (iii) recrystallization component such as Cube, Goss and P. In all the investigated conditions, recrystallization texture component was found in more or less intensity. The volume fraction of recrystallization texture component was higher in RR condition as compared to the NR and RT conditions. The formation of new grains and their growth during deformation depend upon the strain gradient. In the RR condition, there was cross-interaction in lamellar bands inside the macro shear bands as the number of DSR passes increased. This will be led to accumulation of high strain gradient. This finally provides the preferential sites for the grain nucleation.

Fig. 20
figure 20

ODF images showing (a) the ideal texture components of FCC metal (b) the constituent texture components and their maximum values of DSR-deformed sample via NR, RT, and RR methods. (c) Relative fraction of texture components and (d) relative fractions of texture components [25]. Reprinted with permission from Elsevier.

The IPF-maps of Al–Mg–Si alloy after the solution treatment and then cryorolled at strain range of 0, 0.4, 0.8, 2.6 and 3.6 are shown in Fig. 21 [24]. The equiaxed grains with an average grain size of 70 µm are shown in Fig. 21 (a). Figure 21 shows that at low strain elongated cells were formed towards the rolling direction surrounding by the 5˚ to 15˚ orientated boundaries. As the imposed strain was increased to 2.3, lamellar structure surrounding by the high-angle grain boundaries with an average spacing of 3 µm parallel to rolling direction was observed. At the imposed strain of 3.6, these lamellar structures further fragmented and elongated subgrains of size within the range of 100–500 nm being formed.

Fig. 21
figure 21

EBSD micrographs of original and cryorolled 6063 Al alloy at different strains (a) Starting solution treated material (ST) (b) True strain (e) = 0.4 (c) e= 0.8 (d) e= 1.2 (e) e= 2.6 (f) e= 3.6 [24]. Reprinted with permission from Elsevier.

The boundary misorientation profile of ST and CR sample after imposed strain of 0.4, 0.8, 1.2, 2. 6 and 3.6 is shown in Fig. 22 [24]. At low strain of 0.4 and 0.8, fraction of low-angle grain boundaries were higher. But as the true strain increased to 2.6 and 3.6, volume fraction of low-angle grain boundaries was decreased, and volume fraction of high-angle grain boundaries was increased due to the dynamic recrystallization. The CR at high strain was accumulated the higher amount of dislocations as compared to CR at lower strain. These high dislocation densities act as a high driving force available for the dynamic recrystallization.

Fig. 22
figure 22

Frequency histograms of misorientation angles of original and cryorolled 6063 Al alloy at different strains (a) Starting solution treated material (ST) (b) True strain (e) = 0.4 (c) e= 0.8 (d) e= 1.2 (e) e= 2.6 (f) e= 3.6 [24]. Reprinted with permission from Elsevier.

To understand the recovery, recrystallization and grain growth behaviour of CR 6082 Al alloy, annealing was performed from 100 to 350°C, and their IQ-maps superimposed with the grain boundaries and IPF-maps are shown in Figs. 23 and 24 [22]. It was observed from the IQ-map of CR sample that CR sample has the high fraction of low-angle grain boundaries due to the suppression of dislocations annihilation due to the CR. As the annealing temperature increased up-to 250°C, volume fraction of low-angle grain boundaries was decreased and volume fraction of high-angle grain boundaries start increasing. This depicts about the recovery and recrystallization occurred up to annealing at 250°C. As the annealing was performed at 300 and 350°C, volume fraction of high-angle grain boundaries starts decreasing. This depicts the grain growth started after annealing at 300 °C.

Fig. 23
figure 23

Image quality map with superimposition of grain boundary angles of (a) CR and followed by annealing at (b) 150 (c) 200 (d) 250 (e) 300 (f) 350 ˚C, respectively, of 6082 Al alloy [22]. Reprinted with permission from Elsevier.

Fig. 24
figure 24

Inverse pole figure orientation map (IPF) of (a) CR and followed by annealing at (b) 150 (c) 200 (d) 250 e) 300 (f) 350°C, respectively, and Partition of IPF (Grain orientation spread <1˚ and grain size> 4 µm) satisfied by annealing at (g) 300 & (h) 350°C, respectively [22]. Reprinted with permission from Elsevier.

The partition of IPF maps of annealed sample at 300 and 350 ˚C by using the criteria (Grain orientation spread <1˚ and grain size> 4 µm) is shown in Fig. 24 (G & H) [22]. These divided IPF-maps supported the grain growth started after annealing at 300°C.

The IPF-maps of S. T, CR and RTR 6063 Al alloy at strain level of 0.4, 2.3 and 3.8 are shown in Fig. 25 [15]. The black and grey line depicts the high-angle grain boundaries (\(\ge 15^\circ )\) and low-angle grain boundaries (1.5–15°), respectively. The significant changes were found in the RTR and CR sample after imposing strain of 3.8. In the CR sample at strain of 3.8, most of the elongated grains were transformed into the equi-axed subgrains, whereas in the RTR sample at this strain equi-axed subgrain remains absent and elongated subgrains were found.

Fig. 25
figure 25

EBSD micrographs of the starting solution treated (ST), cryorolled (CR) and room temperature rolled of Al 6063 alloy at different strains (arrow indicates rolling direction) with colour coded orientation map along with the [0 0 1] inverse pole figure (a) ST material (b) CR material at e = 0.4 (c) RTR material at e = 0.4 (d) CR material at e = 2.3 (e) RTR material at e = 2.3 (f) CR material at e = 3.8 and (g) RTR material at e = 3.8. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article [15]. Reprinted with permission from Elsevier.

The grain boundary misorientation profile of the subgrain is formed during CR and RTR after imposing strain of 0.4, 2.3 and 3.8 as shown in Fig. 26 [15]. At low strain of 0.4, the low-angle grain boundaries were formed in both CR and RTR samples. After rolling at 2.3, volume fraction of low-angle grain boundaries was decreased and volume fraction of high-angle grain boundaries was increased, in both rolling conditions. But in CR condition volume fraction of high-angle grain boundaries formed was more as compared to high-angle grain boundaries formed in RTR condition. After the rolling at a strain of 3.8, in CR as well as RTR condition low-angle grain boundaries drastically decreased and volume fraction of high-angle grain boundaries further improved. As compared to RTR condition, in the CR condition, significant improvement in the fraction of high-angle grain boundaries was occurred.

Fig. 26
figure 26

Frequency histograms of the starting solution treated (ST), cryorolled (CR), and room temperature rolled (RTR) of Al 6063 alloy at different strains (a) starting ST material (b) CR material at e = 0.4 (c) RTR material at e = 0.4 (d) CR material at e = 0.8 (e) RTR material at e = 0.8 (f) CR material at e = 3.8 and (g) RTR material at e = 3.8 [15]. Reprinted with permission from Elsevier.

The IPF-map of S.T and CR 6063 Al alloy is shown in Fig. 27 [3]. The IPF map of S.T condition shows the equiaxed grain of average grain size of 40 µm, and IPF map of CR sample at the applied strain of 25 and 93% showed the elongated grain towards the rolling direction. It was found that as the deformation strain was increased, grains were severely fragmented and elongated towards rolling direction. The width of the elongated grains reduces to submicrometre range after the 93% deformation.

Fig. 27
figure 27

Inverse pole figures with high-angle grain boundaries indicated by thick lines (arrow indicates rolling direction). (a) Starting material. (b) Deformed 25% at cryogenic temperature. (c) Deformed 93% at cryogenic temperature. (d) Colour-coded orientation map along with the [001] inverse pole figure [3]. Reprinted with permission from Elsevier.

The IPF-maps of CR, RTR, CR followed by annealing at 200 and RTR followed by annealing at 200 °C of 6082 Al alloy are shown in Fig. 28 [87]. Figure 28 (a, b) shows that during CR homogeneous deformation occurred, and after RTR heterogenous deformation occurred. After annealing at 200 °C, it led to formation of subgrains in both CR and RTR conditions. The average grain size obtained in CR condition was 1 µm, and in the case of RTR condition average grain size was 3 µm. The suppression of dislocations annihilations and more dislocations tangling led to fine grain size in the case of CR.

Fig. 28
figure 28

IPF image of (a) CR, (b) RTR, (c) CR+AN (200 °C) and (d) RTR + AN (200 °C) 6082 Al alloy [87]. Reprinted with permission from Elsevier.

The grain boundary misorientation profile of CR, RTR, and annealing at 200 ˚C in both conditions is shown in Fig. 29 [87]. In the CR conditions, high fraction of low-angle grain boundaries were formed due to suppression of dislocations annihilation which led to more dislocations tangling, whereas in the case of RTR conditions due to dynamic recovery during processing itself it led to increase in high-angle grain boundaries and decrease in low-angle grain boundaries.

Fig. 29
figure 29

Misorientation angle of (a) CR, (b) RTR, (c) CR+AN (200 °C) and ( d) RTR+AN (200 ) 6082 Al alloy [87]. Reprinted with permission from Elsevier.

The IQ-map and KAM-map of the solution treated 6082 Al alloy are shown in Fig. 30 [78]. It can be observed from these maps that crack was moved along the grain boundaries in S.T condition. It was due to the absence of secondary phase particles, and other impurities led to grain boundaries become weaker region as compared to grain interior. In the case of CR, RTR and annealed condition, crack was moved trans-granularly. It was found that during crack propagation in the case of CR, RTR, and annealed conditions, at the crack tip there is precipitation of precipitates which can be seen IQ-map (Fig. 31) [78].

Fig. 30
figure 30

( a) Image quality (IQ) map and (b) kernel average misorientation (KAM) map of solution treated (ST) 6082 Al alloy [78]. Reprinted with permission from Elsevier.

Fig. 31
figure 31

Image quality (IQ) maps (a, b, c, d) and kernel average misorientation (KAM) maps (e, f, g, h) of: CR, RTR, CR+AN and RTR+AN sample, of 6082 Al alloy [78]. Reprinted with permission from Elsevier.

Transmission Electron Microscopy

The TEM micrograph of CR and CR+WR conditions is shown in Fig. 32 [37]. In the CR micrograph, high dislocations tangles and heavily deformed grain can be seen, whereas in the case of CR+WR condition due to dynamic recovery subgrain was formed which is shown in Fig. 32 (b). In the case of CR+WR condition, fine size of uniformly distributed Mg2Si-precipiates was formed which is shown in Fig. 32(c).

Fig. 32
figure 32

TEM micrographs of (a) CR, ( b )CR + WR, (c) Precipitates observed in CR+WR sample [37].

The TEM micrograph of CR+PA and CR+WR+PA samples is shown in Fig. 33 [37]. It was very clear from the TEM micrograph that ultra-fine grain of 300 nm average grain size was formed in both the conditions along with the fine crystallite of 100–150 nm size. These fine crystallites were formed along the second phase particles (SP). During the peak aging process, dislocation density was decreased due to the recovery, which further contributes to ductility during tensile testing. There is high density of β′′-precipitates of needle shaped found in CR+WR+PA condition.

Fig. 33
figure 33

TEM micrographs of Al 6061 alloy (a) CR+PA, (b) CR+WR+PA, (c) Needle-shaped precipitates of Mg2Si observed in CR+WR+PA [37]. Reprinted with permission from Elsevier.

The distribution of fine size particle or dispersoid in three different states, i.e. H0, H5 and H16, is shown in Fig. 34 [33]. Figure 34 (a & b) shows that in H0 and H5 rod-shaped Mg2Si and rounded α-Al(FeMn)Si particles were found in low number density (< 2 µm-2), whereas in the H16 condition nano-dispersoid (12.6 nm diameter) were found in the α-Al matrix in high number density (> 60 µm-2). These dispersoid belong to the combination of Mg and Si atoms, or in other words, these are meta-stable precipitates of Mg2Si.

Fig. 34
figure 34

TEM micrograph showing the distribution of fine particle or dispersoid in (a) H0, (b) H5, (c) H16 [33]. Reprinted with permission from Elsevier.

Figure 35 depicts that short homogenization time led to undissolved dispersoid of α−Al(Fe,Mn)Si as compared to longer time homogenization (Fig. 35(b)) [33]. It can be inferred that meta-stable Mg2Si formation occurred during the post-homogenization cooling.

Fig. 35
figure 35

TEM images showing the distribution of fine particles in the water-quenched homogenized samples: (a) H5-Quench, (b) H16-Quench [33]. Reprinted with permission from Elsevier.

Figure 36 depicts the micrograph of CR and RTR sample after the 92% thickness reduction [32]. It was clear from the micrograph of CR sample (Fig. 36(a)) that during CR ill-defined grain boundaries, dislocation tangles and heavily deformed grain were formed, whereas in the case of RTR subgrain surrounding by the well-defined grain boundaries was observed. When the CR sample underwent for annealing at 300 °C, recrystallized grain of 500 nm size was formed (Fig. 37) [32].

Fig. 36
figure 36

TEM microstructure of CR and RTR materials after 92% thickness reduction (a) CR material; (b) RTR material [32]. Reprinted with permission from Elsevier.

Fig. 37
figure 37

TEM micrograph of CR (92% reduction) 6063 Al alloy annealed at 300 °C for 5 min [32]. Reprinted with permission from Elsevier.

The TEM micrograph of CR, CR+SA and CR+WR sample is shown in Fig. 38 [30]. After the CR dislocation, cell structure were formed of average width of 200–300 nm surrounding by the low-angle grain boundaries. In the case of CR+SA, subgrain of 150–300 nm was formed. The dislocation density was decreased in the case of CR+SA as compared to CR due to the static recovery or recrystallization led by the annealing. In the case of CR+WR sample, partially recovered grain with full of dislocations tangles was observed.

Fig. 38
figure 38

TEM micrographs of Al 6061 alloy processed through various conditions: (a) CR, (b) CR+SA, (c) CR+WR [30]. Reprinted with permission from Elsevier.

Figure 39 depicts TEM micrograph of CR+SA and CR+WR sample after the peak ageing of 6061 Al alloy [30]. It was observed from the micrograph that after the peak ageing dislocation density was reduced in both conditions. The peak ageing led to static recovery in both conditions, and then, dislocation tangles were transformed into the more regular low-angle grain boundaries.

Fig. 39
figure 39

TEM micrographs of CR+SA and CR+WR Al 6061 alloy after peak ageing treatment: (a) SA+PA, (b) WR+PA [30]. Reprinted with permission from Elsevier.

Figure 40 depicts the TEM micrograph of SA+PA sample [30]. The dark field imaging depicts the precipitates in SA+PA sample. Figure 40 (b) depicts the micrograph from some other region. It can be concluded form the both figures that precipitates distributed heterogeneously throughout the sample. Figure 40(c) depicts the fine needle-shaped (average length of needle 11 nm) and spherical-shaped particles (size of 6 nm) at higher magnification. The spherical-shaped particles might be result of the end section of needle-shaped particles.

Fig. 40
figure 40

TEM image of precipitates in SA+PA sample: (a) dark field image of precipitates at lower magnification, (b) bright field image of precipitates at lower magnification, (c) bright field image of spherical- and needle-shaped precipitates at higher magnification, (d) SAED pattern of image shown in (a) [30]. Reprinted with permission from Elsevier.

Figure 41 represents the TEM micrograph of WR+PA sample. It was understood from the figure that after the WR+PA it led to the formation of very fine needle-shaped particles (2.5 nm average diameter, 8 nm average diameter) [30].

Fig. 41
figure 41

TEM image of precipitates in WR+PA sample: (a) bright field image of precipitates at lower magnification, (b) bright field image of needle-shaped precipitates at higher magnification, (c) SAED pattern of image shown in (a) [30]. Reprinted with permission from Elsevier.

The TEM micrograph of 6063 Al after the CR is shown in Fig. 42 [29]. This figure depicts that after the CR highly deformed grains containing the high dislocation densities were formed. The fragmentation of grains led to formation of subgrain of ultra-fine regime which were observed. The SAED pattern supports about the poly crystalline nature of image.

Fig. 42
figure 42

TEM micrograph of the solid solution treated, cryorolled Al 6063 alloy (state 2) (a) TEM micrograph (b) Electron diffraction pattern [29]. Reprinted with permission from Elsevier.

The TEM micrograph of CR+WR sample after rolling at 100 to 250 °C is shown in Fig. 43 [27]. It was observed that after rolling at 100 °C heavily deformed structure obtained. Rolling at this temperature further fragmented the spherical shape AlFeMnSi-phase to a size of 10–20 nm. Rolling at 150 ˚C led to formation of subgrain surrounding by the well-defined grain boundaries and size of AlFeMnSi-phase increased to 30–50 nm. After rolling at 200°C, fully recovered grains were formed and size of AlFeMnSi phase increased to 40–100 nm. After rolling at 250 °C, recrystallized grain of size more than 1 µm was formed and size of AlFeMnSi-phase increased to 80-200 nm.

Fig. 43
figure 43

Bright field TEM image of CR followed by WR at (a) 100 (b) 150 (c) 200 (d) 250°C of 6082 Al alloy [27]. Reprinted with permission from Elsevier.

The TEM micrograph of Al–Mg–Si alloy processed through the CR at a strain of 3.6 is shown in Fig. 44 [24]. It is evident from the figure that after the cryorolling heavily strained grain within the range of 100–400 nm was formed. The SAED pattern of this image depicts the polycrystalline nature of material.

Fig. 44
figure 44

TEM micrograph and SAD pattern of the cryorolled Al 6063 alloy after 97% reduction (e= 3.6) (a) TEM microstructure (b) SAD pattern [24]. Reprinted with permission from Elsevier.

The TEM micrograph of CR and CR+WR 6061 Al alloy is shown in Fig. 45 [23]. Figure 45(a) shows that after CR at a strain of 2.3 heavily deformed subgrain containing high dislocation density of average size of 500 nm was formed, whereas in the CR+WR sample, due to warm rolling at 145 °C partially recovered grains surrounding by the well-defined grain boundaries were formed.

Fig. 45
figure 45

Bright field TEM micrograph of Al 6061 alloy: (a) CR material and (b) WR material [23]. Reprinted with permission from Elsevier.

Figure 15 depicts the TEM micrograph of ST sample heated up to the formation of β″and β′- precipitates, respectively [23]. Figure 46 (a) shows needle-shaped β″-precipitates were formed homogenously through the matrix, whereas Fig. 46 (b) depicts the homogenously formation of rod-shaped β′- precipitates throughout the matrix.

Fig. 46
figure 46

Bright field TEM micrograph of ST material isochronally heated up to end of the major exothermic peaks (a) peak B (b) peak C [23]. Reprinted with permission from Elsevier.

Figures 47 and 48 depict the TEM micrograph of CR sample heated up to 238 at the heating rate of 20 °C/min [23]. It was cleared that two type of precipitates, i.e. needle-shaped and rod-shaped precipitates, were formed. The length of the needle-shaped precipitates was varied from 10 to 80 nm, and average diameter was 7 nm, whereas the average diameter for the rod-shaped precipitates was 35 nm.

Fig. 47
figure 47

Dark field TEM micrographs and (B = [001]A1 ) SAED for CR material heated to 238 (half reaction temperature) with 20 °C/min heating rate: (a) needle-shaped β″ precipitates at lower magnification, (b) (B=[001]A1) SAED for figure (a), and (c) β″ precipitates at higher magnification observed at different locations [23]. Reprinted with permission from Elsevier.

Fig. 48
figure 48

TEM micrograph of CR material heated to 238  °C (half reaction temperature) with 20 /min heating rate: (a) dark field TEM micrograph rod-shaped β′ precipitates, (b) bright field TEM micrograph of rod-shaped β′ precipitates observed at different location and (c) (B=[111]Al) SAED micrograph of figure (b) [23]. Reprinted with permission from Elsevier.

Figure 49 depicts the TEM micrograph of the CR sample after annealing up to the 275 at heating rate of 20°C/min up to the complete formation of β′-precipitates [23]. At this temperature, needle-shaped and rod-shaped precipitates of different sizes were found. It might be due to the (i) simultaneously precipitation of both precipitates at this temperature, (ii) transformation of β″ to β′-precipitates as the temperature increased.

Fig. 49
figure 49

Dark field TEM images of CR material heated to 275 °C (end reaction temperature) with 20 /min heating rate. Co-existence of very fine and coarse precipitates corresponding to β″ and β′ precipitates at (a) lower magnification and (b) higher magnification [23]. Reprinted with permission from Elsevier.

TEM micrograph of CR+WR 6061 Al alloy after heating up-to the 237°C is shown in Fig. 50. [23]. The major difference in the CR and CR+WR 6061 Al alloy heated up to this temperature was the size and distribution of precipitates. In the CR+WR 6061 Al alloy, β″ (in the range of 1-2.5 nm) & β′ precipitates (average diameter of 10 nm) were found, respectively. After heating up to the 295 ˚C, most of the β″-precipitates were transformed in to the β′ precipitates (Fig. 51).

Fig. 50
figure 50

Dark field TEM micrographs of WR material heated to 236  °C (half reaction temperature) with 20  °C/min heating rate: (a) needle-shaped β″ precipitates at lower magnification, (b) magnified view of figure (a), (c) existence of rod-shaped β′ precipitates in the same sample and (d) (B=[011]Al) SAED micrograph for figure (a) [23]. Reprinted with permission from Elsevier.

Fig. 51
figure 51

Dark field TEM micrographs of WR material heated to 295  °C (end reaction temperature) with 20 /min heating rate, showing fine β′-rod-shaped precipitates [23]. Reprinted with permission from Elsevier.

Figure 52 depicts the TEM micrograph of platelets shape stable β-precipitates formation in the CR as well as in CR+WR sample after the heating up to the 365 °C [23]. The precipitation strengthening contribution from this Mg2Si stable precipitates was small as compared to meta-stable precipitates.

Fig. 52
figure 52

Dark filed TEM micrograph along with SAED patterns of CR and WR material heated to 365  °C with 20 /min heating rate: (a) CR and (b) WR [23]. Reprinted with permission from Elsevier.

Figure 53 depicts the TEM micrograph of 6082 Al after the solution treatment, CR and CR followed by annealing from 150 to 300 °C for 1 h [22]. Figure 53 (a) shows the TEM image of the undissolved AlMn and AlMnSi-based dispersoids after the solution treatment. Figure 53 (b) represents the deformed grains containing dislocation tangles surrounding by the ill-defined grain boundaries. After annealing at 150 °C in the CR sample, subgrains surrounding by the well-defined grain boundaries of size 200–500 nm were formed. After annealing at 175°C, recovery takes place and needle-shaped β″-precipitates (Fig. 54(a)) of length 20–40 nm were found. The size of the subgrain remains within the range of 200–500 nm. After annealing at 225 ˚C, fully recovered and unrecovered areas were found. After annealing at 300 °C, recrystallized grains were found in the CR sample.

Fig . 53
figure 53

TEM Image of (a) ST, (b) CR followed by annealing at (c) 150 (d) 200 (e) 250 & (f) 300°C, respectively [22]. Reprinted with permission from Elsevier.

Fig. 54
figure 54

Needle (β″)-, rod (β′)- and plate-shaped (β) Mg2Si were shown in TEM micrograph after annealing at (a) 175 , (b) 225 & (c) 300°C, respectively [22]. Reprinted with permission from Elsevier.

Figure 54 depicts the TEM micrograph of CR sample after the annealing at 175 , 225 and 300 °C, respectively [22]. Figure 23 shows that after annealing at 175C needle-shaped β″-precipitates, after annealing at 225 ˚C needle (β″)- & rod-shaped (β′)-precipitates and after annealing at 300 ˚C stable Mg2Si-precipiates (β) were appeared.

Figure 55 represents the TEM image of CR and RTR Al 6063 alloy, after applying the true strain of 2.3 [15]. In the CR sample, heavily deformed grains containing high dislocation density surrounding by the ill-defined grain boundaries and dislocation tangles were formed. In the case of RTR, subgrains containing lesser dislocation density as compared to CR condition surrounding by the well-defined grain boundaries were formed. The dislocation tangles were remained absent in the case of RTR condition.

Fig. 55
figure 55

TEM micrographs of cryorolled (CR) and room temperature rolled (RTR) of Al 6063 alloy at a strain of 2.3; (a) CR material (b) RTR material [15]. Reprinted with permission from Elsevier.

Figure 56 represents the TEM micrograph of CR and RTR sample after applying the true strain of 3.8 [15]. At this, applied strain in the CR sample equiaxed subgrains surrounding by the well-defined grain boundaries was formed. These subgrains contain the lesser dislocation densities. In the case of RTR condition, at this applied strain elongated subgrain containing the dislocations was formed. At this applied strain, CR condition has the lesser dislocation density as compared to RTR condition.

Fig. 56
figure 56

TEM micrographs of cryorolled (CR) and room temperature rolled (RTR) of Al 6063 alloy at a strain of 3.8; (a) CR material (b) RTR material [15]. Reprinted with permission from Elsevier.

Differential Scanning Calorimetry (DSC) Analysis

The DSC peaks of Al 6061 alloy after processing at ST, CR and WR conditions at the scan rate of 20 /min over the temperature limit of 0-450 °C are shown in Fig. 57 [23]. The exothermic peak A is corresponding to the formation of nanoclusters. This peak was formed due to the formation of Mg/Si co-clusters. The Mg/Si co-clustering followed the sequence: Si, Mg atom clusters→ dissolution of Mg clusters→ formation of Mg/Si co clusters. In some cases, clustering reaction completed in three stages: first two reaction completed in with in one 1 h and third reaction continues to occur over the two weeks. As in the CR DSC peaks, two cluster exothermic reactions can be seen. Two large exothermic peaks (B and C) corresponded to the formation of β″ and β′-precipitates, respectively. The exothermic peak D was corresponding to the formation of stable Mg2Si phase. There is no endothermic peak found between the β″ and β′-precipitates. This ensures that β″-precipitates to β′-precipitates formation is a transformation reaction. In the case of CR conditions, cluster 1 could formed due to the interaction of Mg/Si atom with the dislocation or vacancies. The cluster 2 reaction shifted to high temperature side might be vacancies were annihilated due to interaction with the dislocations, and this delays the formation of cluster 2 reaction. The major difference between the DSC peaks of ST condition and rolled condition is the precipitation kinetics was accelerated in the rolled conditions.

Fig. 57
figure 57

Typical DSC thermograms of Al 6061 alloy at various heat-treated conditions: (a) ST, CR and WR and (b) close up view of image (a) showing only low temperature peaks [23]. Reprinted with permission from Elsevier.

Figure 58 represents the DSC peaks of ST, CR and CR+WR 6061 Al alloy at the varying scan rate from 10 to 25 °C/min [23]. The Kissinger’s analysis was used for calculating the activation energy of the clusters and other strengthening precipitates (Fig. 59). The activation energy obtained for the different state of Mg2Si precipitates is shown in Table 2. The activation energy for the formation of C2 in WR condition is lowest (75 kJ mol−1) as compared to the ST condition (88 kJ mol−1) as well as for the CR condition (87 kJ mol−1). In this study, it was found that dislocations can increase the activation energy by annihilated the vacancies and can decrease the activation energy by acting as a short circuit path which further led to enhancing the precipitation kinetics.

Fig. 58
figure 58

DSC thermograms of Al 6061 alloy at various heat-treated conditions with various heating rates: (a) ST, (b) CR, and (c) WR [23]. Reprinted with permission from Elsevier.

Fig. 59
figure 59

Kissinger plots for various exothermic peaks in (a) ST, (b) CR and (c) WR material [23]. Reprinted with permission from Elsevier.

Table 2 Peak temperatures and activation energies of various processed conditions [23]. Reprinted with permission from Elsevier.

Figure 60 represents the hardness measurement corresponds to the different precipitate’s evolution in the ST, CR and WR sample [23]. In the S.T condition maximum hardness (98 Hv) was achieved after the complete precipitation of β″-precipitates. In the CR condition, maximum hardness (105 HV) was achieved after the isochronal heating up to 200 ˚C. In the WR condition, maximum hardness was achieved after the isochronal heating up to the 150°C. It was observed that maximum hardness is obtained in CR and WR condition much before the complete precipitation of β″-precipitates. The maximum hardness obtained in CR and WR conditions was the combined effect of precipitation strengthening as well as grain boundary strengthening, whereas in the case of S.T condition maximum strength has the major contribution form the β″-precipitates strengthening.

Fig. 60
figure 60

DSC thermograms obtained with 20  °C/min heating rate along with the hardness plots of Al 6061 alloy: (a) ST, (b) CR, and (c) WR [23]. Reprinted with permission from Elsevier.

Figure 61 represents the DSC diagram of the ST, CR and CR followed by annealing from 150 to 400 °C of the 6082 Al alloy [22]. For finding the recovery and recrystallized peak, CR sample was scanned with respect to the S.T sample as shown in Fig. 61(a). These DSC curves were helped in finding the recovery, recrystallization and grain growth starting temperature in the CR 6082 Al alloy. In this study, recovery and recrystallization annealing temperature was find out by analysing different exothermic and endothermic peaks with the help of DSC measurement of CR 6082 Al alloy. The recovery process was completed by annealing up to 250 °C, and recrystallization process was completed within the annealing temperature of 300–350 °C.

Fig. 61
figure 61

DSC peaks of (A) ST, CR and CR Vs ST, and annealing at (B) 150 ,175, 200, 225, 250, 275; & (C) 300, 325, 350, 375, 400°C, respectively [22]. Reprinted with permission from Elsevier.

Figure 62 represents the DSC curve of ST and CR (25 and 93% reduction) of 6063 Al alloy [3]. In the DSC analysis, five exothermic peaks and two endothermic peaks were found in all investigated cases. In this study, they claimed that β″-precipitates to β′-precipitates is not a transformation reaction. First β″-precipitates were dissolved, and then, finally β′-precipitates were formed. The DSC curve of CR up to 93% deformation shifted to left side as compared to other investigated samples. It was due to the availability of more dislocations, accelerated the precipitation kinetics.

Fig. 62
figure 62

DSC plots for the heat flow of precipitate reactions of the starting solution treated and cryorolled (CR) Al 6063 alloy samples: (a) the starting material; (b) CR with 25% reduction; (c) CR with 93% reduction [3]. Reprinted with permission from Elsevier.

The activation energy of maximum strengthening precipitates (β″) was calculated with the help of Kissinger’s analysis for the ST and CR sample as shown in Table 3 and Fig. 63 [3]. The activation energy of CR up to the 93% reduction was lowest as compared to other investigated samples. The structural defect such as dislocations and low-angle grain boundaries enhances the precipitation kinetics in two ways: (i) precipitates nucleation increased due to the decrease in the activation energy, (ii) acts as a short circuit path for diffusion of solute atoms.

Table 3 Activation energy values of β″ precipitation in starting solution treated and cryorolled (25% and 93% thickness reduction) Al 6063 alloy samples using Kissinger analysis [3]. Reprinted with permission from Elsevier.
Fig. 63
figure 63

Kissinger plot for the precipitation kinetics of starting solution treated and cryorolled Al 6063 alloy samples with 25 and 93% thickness reduction [3]. Reprinted with permission from Elsevier.

Figure 64 represents the DSC curves for the ST, CR, CR+SA and CR+WR sample of 6061 Al alloy [30]. The DSC curve of ST sample shows the conventional precipitation sequence of Mg2Si precipitates, i. e. super saturated solid solution (SSSS) → GP (Guinier–Preston)-I zones (peak 1) → GP-II zones (β″) (peak 2) → β′ (peak 3) → β (peak 4). The DSC curve of CR condition has only three exothermic peaks. In this condition, peak corresponded to the β″-precipitates shifted to left side and peak corresponded to the β′-precipitates was suppressed. In the case of CR+SA and CR+WR sample, exothermic peaks corresponded to the cluster’s formation were disappeared, and other peaks position remain same as found in the CR sample, except intensity of exothermic β″-peaks (peak 3) reduced in these two conditions as compared to CR sample.

Fig. 64
figure 64

DSC plots of Al 6061 alloy at a heating rate of 20 ˚C/min, subjected to different processing conditions: (a) ST, (b) CR, (c) CR + SA, (d) CR+WR [30]. Reprinted with permission from Elsevier.

Figure 65 depicts the DSC curve of CR followed by warm rolling at different temperature from 100 to 250 °C of the 6082 Al alloy [27]. As the rolling temperature increased, it shifts the precipitation kinetic to left side. In this study, they claimed that rolling at higher temperature provided the higher energy for the diffusion of solute, which further accelerated the precipitation kinetics.

Fig. 65
figure 65

DSC exothermic and endothermic peaks of WR at different temperature [27]. Reprinted with permission from Elsevier.

Mechanical properties

The mechanical properties obtained after the processing through different combinations of rolling techniques of the Al–Mg–Si alloys are summarized in Table 4. After carefully observation from the literature, it was found that cryorolling (CR) followed warm rolling (WR) at optimized temperature, and then, again ageing of the Al–Mg–Si alloy significantly improves the mechanical properties. In all different combinations of rolling techniques investigated in the present work from the literatures, the CR (70)+WR (20%) at 145 °C followed by peak aging at 125°C for 60 h has highest hardness (130 HV), Yield strength (YS, 375 MPa) and tensile strength (TS, 400 MPa), respectively. This was due the combined effect of the solid solution strengthening, dislocation strengthening, dynamic recovery and dynamic ageing during CR+WR; further, PA contributes in the strength by pinning of the dislocation by the nano-sized precipitates.

Table 4 Mechanical properties of Al–Mg–Si alloy processed through different combinations of rolling techniques.

Discussion

It is evident from the reported literatures that rolling at liquid nitrogen temperatures is more advantageous as compared to rolling at room temperature. It is because of rolling at liquid nitrogen temperature helped in suppression of dislocations annihilation and promoted the dislocations strengthening. The reported literatures support this argument of higher strength of CR sample as compared to RTR sample as listed in Table 4 [22, 23, 26, 30, 32, 37]. Further, cryo-rolling followed by annealing exhibited the superior mechanical properties as compared to CR and RTR sample as listed in Table 4 [14, 15, 22, 23, 26, 29, 32, 78, 87]. In the reported literatures, it was found that CR helped in the suppression of precipitates evolution during the processing and hence provided the more precipitation strengthening during the annealing. It was observed that β′′-precipitates provided the excellent precipitation strengthening as compared to other precipitates (β′ & β). It is because of needle-shaped morphology of β′′-precipitates and their nano-size, provided the excellent pinning effect due to their non-shearble nature. Further researcher performed the warm rolling on the cryorolled sample and observed the superior mechanical properties (hardness, tensile strength, ductility, fatigue strength (low and high cycle fatigue life), fracture strength) as compared to RTR, CR and CR+AN sample as listed in Table 4 [23, 26, 37]. The CR followed by WR helped in precipitation of precipitates as well as dynamic recovery during processing which helped in retaining the higher strength as well as ductility. This dual effect helped in maintaining the higher strength as well as higher ductility. This dual combination further helped in improving the low cycle fatigue life (because of higher ductility), high cycle fatigue life (because of improvement in yield strength) and fracture toughness (due to higher strength and ductility). The mechanical properties can be further improved by performing the peak ageing on the CR+WR sample. In this approach, WR was performed at a temperature lower than the precipitation temperature of β″-precipitates. After this, the ageing on WR sample helped in significant improvement in the mechanical properties of Al–Mg–Si alloys. The slight change in the chemical composition of Al–Mg–Si alloys further affected the precipitation kinetics. It is advised that the optimization of warm rolling temperatures is needed, and for this the use of DSC study will be more effective. In all different rolling technique grain refinement is a primary mechanism, but in the CR + WR technique grain refinement together with the precipitates evolution is the major advantage.

Conclusions

Authors investigated the literatures reported by the researchers in the field of the severe plastic deformation of Al–Mg–Si alloy. It was found during the investigation that rolling technique was widely used in industries to produce the continuous long product such as sheet. Author more focused towards different combinations of rolling techniques used by the researchers as a severe plastic deformation technique, to produce the Ultra-fine/Nano size Al–Mg–Si alloys. Author gathered the results based on the Transmission Electron Microscopy (TEM), Electron Back-scattered diffraction (EBSD), Differential Scanning Calorimetry (DSC) and Mechanical properties, processed through different combinations of rolling techniques of the Al–Mg–Si alloy. The following key points point are concluded based on the investigated literatures:

  1. (1)

    Different combinations of rolling were used by the researchers to produce the ultra-fine-grained Al–Mg–Si alloy. In out of all the combinations of rolling, the CR+WR technique was found more suitable for the production of high specific strength Al–Mg–Si alloy. This technique provides the dual strengthening mechanism at the time of processing itself. The CR helps in accumulating the higher amount of dislocations inside the grains, and WR further contributes in the formation of subgrains surrounding by the well-defined grain boundaries due to dynamic recovery during the processing. These subgrains were formed with in the ultra-fine regime and provide the grain boundary strengthening according to the hall-patch relation. The WR temperature (By optimizing meta stable precipitations) helped in the precipitations of meta-stable precipitates, and strain induced during warm rolling helped in further fragmentation of these precipitates. The CR+WR provides the dual strengthening such as grain boundary strengthening and precipitation strengthening, to the Al–Mg–Si alloy during the processing itself. Further ageing will contribute more in increasing the strength of the CR+WR Al–Mg–Si alloy.

  2. (2)

    Two different conflicted results were presented after the CR processing of Al–Mg–Si alloy. In one approach, it was presented that during the CR as the severe strain was increased the greater number of heavily deformed grain were formed. These grains contained the high dislocation tangles and surrounding by the ill-defined grain boundaries. In the second approach, as the imposed strain increased during cryorolling, subgrain of well-defined grain boundaries were formed. In this approach, high dislocation density available will act as driven force for the formation of subgrain contained lesser amount of dislocation.

  3. (3)

    The precipitation sequence of Mg-/Si-based precipitates, i.e. super saturated solid solution (SSSS) → GP (Guinier-Preston)-I zones (peak 1) → GP-II zones (β″) (peak 2) → β′ (peak 3) → β (peak 4), were remain same in most of the literatures. But there is conflict in the reaction of the formation of the precipitates. In the first approach, the β′-precipitates were formed after the complete dissolution of β″-precipitates. In the second approach, β″-precipitates were completely transformed into the β′ -precipitates during the heat treatment.

The CR+WR technique based on the investigated literatures is recommended for the formation of sheet-based product, where high specific strength of the Al–Mg–Si alloys is required. The Al–Mg–Si alloys processed through the CR + WR have the optimum hardness, yield strength, tensile strength, fatigue properties (both low cycle and high cycle fatigue strength), and fracture toughness as compared to the processed through the other rolling techniques.

Data availability

The raw/processed data required to reproduce these findings cannot be shared at this time as the data also form part of an ongoing study.

References

  1. A. Azushima, R. Kopp, A. Korhonen, D.Y. Yang, F. Micari, G.D. Lahoti, A. Yanagida, Severe plastic deformation (SPD) processes for metals. CIRP Ann.57(2), 716–735 (2008)

    Article  Google Scholar 

  2. M. Vaseghi, H.S. Kim, A combination of severe plastic deformation and ageing phenomena in Al–Mg–Si Alloys. Mater Des. 1980–2015(36), 735–740 (2012)

    Article  CAS  Google Scholar 

  3. S.K. Panigrahi, R. Jayaganthana, V. Pancholi, M. Gupta, A DSC study on the precipitation kinetics of cryorolled Al 6063 alloy. Mater Chem Phys. 122, 188–193 (2010)

    CAS  Article  Google Scholar 

  4. A.S.M. Agena, A study of flow characteristics of nanostructured Al-6082 alloy produced by ECAP under upsetting test. J Mater Process Technol. 209(2), 856–863 (2009)

    CAS  Article  Google Scholar 

  5. S. Lokesh Vendra, S. Goel, R. Nikhil Kumar, Jayaganthan., A study on fracture toughness and strain rate sensitivity of severely deformed Al 6063 alloys processed by multiaxial forging and rolling at cryogenic temperature. Mater Sci Eng A. 686, 82–92 (2017)

    Article  CAS  Google Scholar 

  6. S.K. Panigrahi, R. Jayaganthan, A study on the mechanical properties of cryorolled Al–Mg–Si alloy. Mater Sci Eng A. 480, 299–305 (2008)

    Article  CAS  Google Scholar 

  7. N. Kumar, G.M. Owolabi, R. Jayaganthan, Al 6082 alloy strengthening through low strain multi-axial forging. Mater. Charact. 155, 109761 (2019)

    CAS  Article  Google Scholar 

  8. B. Mirzakhani, Y. Payandeh, Combination of sever plastic deformation and precipitation hardening processes affecting the mechanical properties in Al–Mg–Si alloy. Mater Des. 68, 127–133 (2015)

    CAS  Article  Google Scholar 

  9. A. Johannes, osterreicher., Combined cyclic deformation and artificial ageing of an Al–Mg–Si alloy. Mater Lett: X. 10, 100072 (2021)

    Google Scholar 

  10. M. Kolar, K.O. Pedersen, S. Gulbrandsen-Dahl, K. Marthinsen, Combined effect of deformation and artificial aging on mechanical properties of Al–Mg–Si Alloy. Trans Nonferrous Met Soc China. 22(8), 1824–1830 (2012)

    CAS  Article  Google Scholar 

  11. A.P. Murugesan, V. Rajinikanth, B. Mahato, M. Wegner, M. Witte, G. Wilde, S.G. Chowdhury, Concurrent precipitation and associated texture evolution in AA 6082 alloy during high pressure torsion (HPT) processing. Mater Sci Eng: A. 700, 487–494 (2017)

    CAS  Article  Google Scholar 

  12. X. Fan, Z. He, X. Knag, S. Yuan, Deformation and strengthening analysis of Al–Mg–Si alloy sheet during hot gas forming with synchronous die quenching. J Manuf Process. 57, 452–461 (2020)

    Article  Google Scholar 

  13. Y.N. Kwon, Y.S. Lee, J.H. Lee, Deformation behaviour of Al–Mg–Si alloy at the elevated temperature. J. Mater. Process. Technol. 187–188, 533–536 (2007)

    Article  CAS  Google Scholar 

  14. S.K. Panigrahi, R. Jayaganthan, Development of ultrafine grained Al–Mg–Si alloy with enhanced strength and ductility. J Alloys Compd. 470, 285–288 (2009)

    CAS  Article  Google Scholar 

  15. S.K. Panigrahi, R. Jayaganthan, Development of ultrafine-grained Al 6063 alloy by cryorolling with the optimized initial heat treatment conditions. Mater Des. 32, 2172–2180 (2011)

    CAS  Article  Google Scholar 

  16. M. Liu, H.J. Roven, Yu. Yingda, J.C. Werenskiold, Deformation structures in 6082 aluminium alloy after severe plastic deformation by equal-channel angular pressing. Mater. Sci. Eng., A. 483–484, 59–63 (2008)

    Article  CAS  Google Scholar 

  17. H.J. Roven, M. Liu, J.C. Werenskiold, Dynamic precipitation during severe plastic deformation of an Al–Mg–Si Al alloy. Mater. Sci. Eng., A. 483–484, 54–58 (2008)

    Article  CAS  Google Scholar 

  18. Y. Weng, Z. Jia, L. Ding, Y. Pan, Y. Liu, Q. Liu, Effect of Ag and Cu additions on natural aging and precipitation hardening behavior in Al–Mg–Si alloys. J Alloys Compd. 695, 2444–2452 (2017)

    CAS  Article  Google Scholar 

  19. G. Das, M. Das, S. Ghosh, A.K. Paritosh Dubey, Ray., Effect of aging on mechanical properties of 6063 Al-alloy using instrumented ball indentation technique. Mater Sci Eng A. 527, 1590–1594 (2010)

    Article  CAS  Google Scholar 

  20. T. Ye, L. Li, P. Guo, G. Xiao, Z. Chen, Effect of aging treatment on the microstructure and flow behavior of 6063 aluminum alloy compressed over a wide range of strain rate. Int. J. Impact Eng. 90, 72–80 (2016)

    Article  Google Scholar 

  21. D. Odoh, Y. Mahmoodkhani, M. Wells, Effect of alloy composition on hot deformation behavior of some Al–Mg–Si alloys. Vacuum. 149, 248–255 (2018)

    CAS  Article  Google Scholar 

  22. N. Kumar, P.N. Rao, R. Jayaganthan, H.G. Brokmeier, Effect of cryorolling and annealing on recovery, recrystallisation, grain growth and their influence on mechanical and corrosion behaviour of 6082 Al alloy. Mater chem phys. 165, 177–187 (2015)

    CAS  Article  Google Scholar 

  23. P. Nageswera Rao, B. Viswanadh, R. Jayaganthan, Effect of cryorolling and warm rolling on precipitation evolution in Al 6061 alloy. Mater Sci Eng A. 606, 1–10 (2014)

    Article  CAS  Google Scholar 

  24. S.K. Panigrahi, R. Jayaganthan, V. Chawla, Effect of cryorolling on microstructure of Al–Mg–Si alloy. Mater Lett. 62, 2626–2629 (2008)

    CAS  Article  Google Scholar 

  25. H.W. Yang, I.P. Widiantara, Y.G. Ko, Effect of deformation path on texture and tension properties of submicrocrystalline Al–Mg–Si alloy fabricated by differential speed rolling. Materials Letter. 213, 54–57 (2018)

    CAS  Article  Google Scholar 

  26. R. Nikhil Kumar, H.-G. Jayaganthan, Effect of deformation temperature on precipitation, microstructural evolution, mechanical and corrosion behavior of 6082 Al alloy. Trans. Nonferrous Met. Soc. China. 27, 475–492 (2017)

    Article  Google Scholar 

  27. N. Kumar, R. Sunkulp Goel, H.B. Jayaganthan, Effect of grain boundary misorientaton, deformation temperature and AlFeMnSi-phase on fatigue life of 6082 Al alloy. Mater Charact. 124, 229–240 (2017)

    CAS  Article  Google Scholar 

  28. O. Engler, C.D. Marioara, Y. Aruga, M. Kozuka, O.R. Myhr, Effect of natural ageing or pre-ageing on the evolution of precipitate structure and strength during age hardening of Al–Mg–Si alloy AA 6016. Mater. Sci. Eng., A. 759, 520–529 (2019)

    CAS  Article  Google Scholar 

  29. S.K. Panigrahi, R. Jayaganthan, V. Pancholi, Effect of plastic deformation conditions on microstructural characteristics and mechanical properties of Al 6063 alloy. Mater Des. 30, 1894–1901 (2009)

    CAS  Article  Google Scholar 

  30. D. Singh, R. Jayaganthan, Effect of post cryorolling treatments on microstructural and mechanical behaviour of ultrafine grained Al–Mg–Si alloy. J Mater Sci Technol. 30(10), 998–1005 (2014)

    Article  CAS  Google Scholar 

  31. Y. Zi, L. Zeqin, D. Leyvraz, J. Banhart, Effect of pre-ageing on natural secondary ageing and paint bake hardening in Al–Mg–Si alloys. Materialia. 7, 100413 (2019)

    CAS  Article  Google Scholar 

  32. S.K. Panigrahi, R. Jayaganthan, Effect of rolling temperature on microstructure and mechanical properties of 6063 Al alloy. Mater. Sci. Eng., A. 492, 300–305 (2008)

    Article  CAS  Google Scholar 

  33. Lu et al., Effect of plastic deformation conditions on microstructural characteristics and mechanical properties of Al 6063 alloy-phase particle evolution in a twin-roll-casted Al-Mg-Si alloy on recrystallization texture and mechanical anisotropy. Mater. Charact. 176, 111038 (2021)

    Article  CAS  Google Scholar 

  34. H. Zhao, Q. Pan, Q. Qin, Wu. Yujiao, Su. Xiangdong, Effect of the processing parameters of friction stir processing on the microstructure and mechanical properties of 6063 aluminum alloy. Mater. Sci. Eng., A. 751, 70–79 (2019)

    CAS  Article  Google Scholar 

  35. Yu. Wang, Y. Deng, J. Chen, Q. Dai, X. Guo, Effects of grain structure related precipitation on corrosion behavior and corrosion fatigue property of Al–Mg–Si alloy. J. mater. Res. technol. 9(3), 5391–5540 (2020)

    CAS  Article  Google Scholar 

  36. X. Qian, N. Parson, X. Grant Chen, Effects of Mn content on recrystallization resistance of AA6082 aluminum alloys during post-deformation annealing. J Mater Sci Technol. 52, 189–197 (2020)

    Article  Google Scholar 

  37. P.R. Nageswara raoJayaganthan, Effects of warm rolling and ageing after cryogenic rolling on mechanical properties and microstructure of Al 6061 alloy. Mater Des. 39, 226–233 (2012)

    Article  CAS  Google Scholar 

  38. M.J. Kim, M.G. Lee, K. Hariharan, S.T. Hong, I.S. Choi, D. Kim, H.N. Han, Electric current–assisted deformation behavior of Al–Mg–Si alloy under uniaxial tension. Int J Plast. 94, 148–170 (2017)

    CAS  Article  Google Scholar 

  39. W.J. Kim, S.J. Yoo, Enhanced ductility and deformation mechanisms of ultrafine-grained Al–Mg–Si alloy in sheet form at warm temperatures. Scripta Mater. 61, 125–128 (2009)

    CAS  Article  Google Scholar 

  40. K. Majchrowicz, Z. Pakiela, W. Chrominski, M. Kulczyk, Enhanced strength and electrical conductivity of ultrafine-grained Al–Mg–Si alloy processed by hydrostatic extrusion. Mater. Charact. 135, 104–114 (2018)

    CAS  Article  Google Scholar 

  41. L. Winter, K. Hockauf, S. Winter, T. Lampke, Equal-channel angular pressing influencing the mean stress sensitivity in the high cycle fatigue regime of the 6082 aluminum alloy. Mater. Sci. Eng., A. 795, 140014 (2020)

    CAS  Article  Google Scholar 

  42. E. Ortiz-cuellar, M.A.L. Hernandez-Rodriguez, E. Garcia-Sanchez, Evaluation of the tribological properties of an Al–Mg–Si alloy processed by severe plastic deformation. Wear. 271, 1828–1832 (2011)

    CAS  Article  Google Scholar 

  43. O.P. Gbenebor, O.S.I. Fayomi, A.P.I. Popoola, A.O. Inegbenebor, F. Oyawale, Extrusion die geometry effects on the energy absorbing properties and deformation response of 6063-type Al–Mg–Si aluminum alloy. Results in Physics. 3, 1–6 (2013)

    Article  Google Scholar 

  44. D.C.C. Magalhaes, A.M. Kliauga, V.L. Sordi, Flow behavior and fracture of Al–Mg–Si alloy at cryogenic temperatures. Trans. Nonferrous Met. Soc. China. 31, 595–608 (2021)

    CAS  Article  Google Scholar 

  45. Q.Y. Yang, Y.A.N.G. Dong, Z.Q. Zhang, L.F. Cao, X.D. Wu, G.J. Huang, L.I.U. Qing, Flow behavior and microstructure evolution of 6A82 aluminium alloy with high copper content during hot compression deformation at elevated temperatures. Transact Nonferrous Met Soc China. 26(3), 649–657 (2016)

    CAS  Article  Google Scholar 

  46. Y. Aruga, M. Kozuka, T. Sato, Formulation of initial artificial age-hardening response in an Al–Mg–Si alloy based on the cluster classification using a high-detection-efficiency atom probe. J. Alloy. Compd. 739, 1115–1123 (2018)

    CAS  Article  Google Scholar 

  47. D. Pan, Y. Wang, Q. Guo, D. Zhang, X. Xu, Y. Zhao, Grain refinement of Al–Mg–Si alloy without any mechanical deformation and matrix phase transformation via cyclic electro-pulsing treatment. Mater Sci Eng: A. 807, 140916 (2021)

    CAS  Article  Google Scholar 

  48. S.H. Wang, C.H. Liu, J.H. Chen, X.L. Li, D.H. Zhu, G.H. Tao, Hierarchical nanostructures strengthen Al–Mg–Si alloys processed by deformation and aging. Mater Sci Eng: A. 585, 233–242 (2013)

    CAS  Article  Google Scholar 

  49. L. Winter, K. Hockau, T. Lampke, High cycle fatigue behavior of the severely plastically deformed 6082 aluminum alloy with an anodic and plasma electrolytic oxide coating. Surf. Coat. Technol. 349, 576–658 (2018)

    CAS  Article  Google Scholar 

  50. J. Li, X. Wu, L. Cao, B. Liao, Y. Wang, Q. Liu, Hot deformation and dynamic recrystallization in Al–Mg–Si alloy. Mater Charact. 173, 110976 (2021)

    CAS  Article  Google Scholar 

  51. H. Liao, Wu. Yuna, K. Zhou, J. Yang, Hot deformation behavior and processing map of Al-Si-Mg alloys containing different amount of silicon based on Gleebe-3500 hot compression simulation. Mater. Des. 65, 1091–1099 (2015)

    CAS  Article  Google Scholar 

  52. Hu. Jiamin, W. Zhang, Fu. Dingfa, J. Teng, H. Zhang, Improvement of the mechanical properties of Al–Mg–Si alloys with nano-scale precipitates after repetitive continuous extrusion forming and T8 tempering. J Mater Res Technol. 8(6), 5950–5960 (2019)

    Article  CAS  Google Scholar 

  53. J. Zhang, M. Ma, F. Shen, D. Yi, B. Wang, Influence of deformation and annealing on electrical conductivity, mechanical properties and texture of Al–Mg–Si alloy cables. Mater Sci Eng: A. 710, 27–37 (2018)

    CAS  Article  Google Scholar 

  54. J. González, S. Bagherifard, M. Guagliano, I.F. Pariente, Influence of different shot peening treatments on surface state and fatigue behaviour of Al 6063 alloy. Eng Fract Mech. 185, 72–81 (2017)

    Article  Google Scholar 

  55. G.J. Gao, H.E. Chen, L.I. Yong, J.D. Li, Z.D. Wang, R.D.K. Misra, Influence of different solution methods on microstructure, precipitation behavior and mechanical properties of Al–Mg–Si alloy. Trans Nonferrous Met Soci China. 28(5), 839–847 (2018)

    CAS  Article  Google Scholar 

  56. D. Jiang, C. Wang, Influence of microstructure on deformation behaviour and fracture mode of Al–Mg–Si alloy. Mater. Sci. Eng., A. 352, 29–33 (2003)

    Article  CAS  Google Scholar 

  57. E. Cerri, P. Leo, Influence of severe plastic deformation on aging of Al–Mg–Si alloys. Mater. Sci. Eng., A. 410–411, 226–229 (2005)

    Article  CAS  Google Scholar 

  58. M. Hussain, P.N. Rao, D. Singh, R. Jayaganthan, S. Goel, K.K. Saxena, Insight to the evolution of nano precipitates by cryo rolling plus warm rolling and their effect on mechanical properties in Al 6061 alloy. Mater Sci Eng: A. 811, 141072 (2021)

    CAS  Article  Google Scholar 

  59. V. Kumar, D. Kumar, Investigation of tensile behaviour of cryorolled and room temperature rolled 6082 Al alloy. Mater. Sci. Eng., A. 691, 211–217 (2017)

    CAS  Article  Google Scholar 

  60. J. Bouquerel, B. Diawara, A. Dubois, M. Dubar, J.-B. Vogt, D. Najjar, Investigations of the microstructural response to a cold forging process of the 6082–T6 alloy. Mater. Des. 68, 245–258 (2015)

    CAS  Article  Google Scholar 

  61. N.R. Bochvar, O.V. Rybalchenko, N.Y. Tabachkova, G.V. Rybalchenko, N.P. Leonova, L.L. Rokhlin, Kinetics of phase precipitation in Al–Mg–Si alloys subjected to equal-channel angular pressing during subsequent heating. J Alloys Compd. 881, 160583 (2021)

    CAS  Article  Google Scholar 

  62. P.N. Rao, D. Singh, R. Jayaganthan, Mechanical properties and microstructural evolution of Al 6061 alloy processed by multidirectional forging at liquid nitrogen temperature. Mater Des. 56, 97–104 (2014)

    CAS  Article  Google Scholar 

  63. L. Wan, Y. Huang, W. Guo, S. Lv, J. Feng, Mechanical Properties and Microstructure of 6082–T6 Aluminum Alloy Joints by Self-support Friction Stir Welding. J. Mater. Sci. Technol. 30(12), 1243–1250 (2014)

    CAS  Article  Google Scholar 

  64. X. Zhang, L.K. Huang, B. Zhang, Y.Z. Chen, F. Liu, Microstructural evolution and strengthening mechanism of an Al–Si–Mg alloy processed by high-pressure torsion with different heat treatments. Mater Sci Eng: A. 794, 139932 (2020)

    CAS  Article  Google Scholar 

  65. M. Song, J. Kim, Microstructural evolution at the initial stage of two-step aging in an Al–Mg–Si alloy characterized by a three dimensional atom probe. Mater Sci Eng: A. 815, 141301 (2021)

    CAS  Article  Google Scholar 

  66. X.D. Wang, L.I.U. Xiong, D.I.N.G. Hao, S.R. Yan, Z.H. Xie, B.Q. Pan, W.Y. Wang, Microstructure and mechanical properties of Al− Mg− Si alloy U-shaped profile. Trans Nonferrous Met Soc China. 30(11), 2915–2926 (2020)

    CAS  Article  Google Scholar 

  67. Liu et al., Microstructure and mechanical properties of Al–Mg–Si alloy fabricated by a short process based on sub-rapid solidification. J Mater Sci Technol. 41, 178–186 (2020)

    Article  Google Scholar 

  68. K.J. Al-Fadhalah, A.I. Almazrouee, A.S. Aloraier, Microstructure and mechanical properties of multi-pass friction stir processed aluminum alloy 6063. Mater. Des. 53, 550–560 (2014)

    CAS  Article  Google Scholar 

  69. X. Meng, S. Yang, Yu. Yubao Huang, J.G. Fang, Qi. Xiong, C. Duan, Microstructure characterization and mechanism of fatigue crack propagation of 6082 aluminum alloy joints. Mater. Chem. Phys. 257, 123734 (2021)

    CAS  Article  Google Scholar 

  70. S. Liu, Q. Pan, M. Li, X. Wang, X. He, X. Li, J. Lai, Microstructure evolution and physical-based diffusion constitutive analysis of Al–Mg–Si alloy during hot deformation. Mater Des. 184, 108181 (2019)

    CAS  Article  Google Scholar 

  71. R. Lu, S. Zheng, J. Teng, J. Hu, D. Fu, J. Chen, H. Zhang, Microstructure, mechanical properties and deformation characteristics of Al–Mg–Si alloys processed by a continuous expansion extrusion approach. J Mater Sci Technol. 80, 150–162 (2021)

    CAS  Article  Google Scholar 

  72. M. Liu, J. Chen, Y. Lin, Z. Xue, H.J. Roven, P.C. Skaret, Microstructure, mechanical properties and wear resistance of an Al–Mg–Si alloy produced by equal channel angular pressing. Prog Nat Sci: Mater Int. 30(4), 485–493 (2020)

    CAS  Article  Google Scholar 

  73. M. Shakoori Oskooie, H. Asgharzadeh, H.S. Kim, Microstructure, plastic deformation and strengthening mechanisms of an Al–Mg–Si alloy with a bimodal grain structure. J alloys an compound. 632, 540–5548 (2015)

    CAS  Article  Google Scholar 

  74. S. Lin, Y.L. Deng, J.G. Tang, S.H. Deng, H.Q. Lin, L.Y. Ye, X.M. Zhang, Microstructures and fatigue behavior of metal-inert-gas-welded joints for extruded Al–Mg–Si alloy. Mater Sci Eng: A. 745, 63–73 (2019)

    CAS  Article  Google Scholar 

  75. M. Das, G. Das, M. Ghosh, V. Matthias Wegner, S.G. Rajnikant, T.K.P. Chowdhury, Microstructures and mechanical properties of HPT processed 6063 Al alloy. Mater Sci Eng A. 558, 525–532 (2012)

    CAS  Article  Google Scholar 

  76. S.B. Puplampu, A. Siriruk, A. Sharma, D. Penumadu, Multiaxial deformation behavior of aluminum alloy 6061 subjected to fire damage. Mech. Mater. 159, 103885 (2021)

    Article  Google Scholar 

  77. J.K. Sunde, C.D. Marioara, S. Wenner, R. Holmestad, On the microstructural origins of improvements in conductivity by heavy deformation and ageing of Al–Mg–Si alloy 6101. Mater Charact. 176, 111073 (2021)

    CAS  Article  Google Scholar 

  78. N. Kumar, G.M. Owolabi, R. Jayaganthan, O.O. Ajide, Plane stress fracture toughness of cryorolled 6082 Al alloy. Theoret. Appl. Fract. Mech. 95, 28–41 (2018)

    CAS  Article  Google Scholar 

  79. M. Yang, H. Chen, A. Orekhov, Q. Lu, X. Lan, K. Li, Y. Du, Quantified contribution of β ″and β′ precipitates to the strengthening of an aged Al–Mg–Si alloy. Mater Sci Eng: A. 774, 138776 (2020)

    CAS  Article  Google Scholar 

  80. J.-H. Zheng, C. Pruncu, K. Zhang, K. Zheng, J. Jiang, Quantifying geometrically necessary dislocation density during hot deformation in AA6082 Al alloy. Mater. Sci. Eng., A. 814, 141158 (2021)

    CAS  Article  Google Scholar 

  81. H. Ma, Q. Hou, Yu. Zhiwei, P. Ni, Stability of 6082–T6 aluminum alloy columns under axial forces at high temperatures. Thin-Walled Struct. 157, 107083 (2020)

    Article  Google Scholar 

  82. S. Dadbakhsh, A. Karimi Taheri, C.W. Smith, Strengthening study on 6082 Al alloy after combination of aging treatment and ECAP process. Mater Sci Eng A. 527, 4758–4766 (2010)

    Article  CAS  Google Scholar 

  83. Y.X. Lai, W. Fan, M.J. Yin, C.L. Wu, J.H. Chena, Structures and formation mechanisms of dislocation-induced precipitates in relation to the age-hardening responses of Al–Mg–Si alloys. J. Mater. Sci. Technol. 41, 127–138 (2020)

    Article  Google Scholar 

  84. Wang et al., Superior mechanical properties induced by the interaction between dislocations and precipitates in the electro-pulsing treated Al–Mg–Si alloys. Mater. Sci. Eng., A. 735, 154–161 (2018)

    CAS  Article  Google Scholar 

  85. L. Shi, K. Baker, R. Young, J. Kang, J. Liang, B. Shalchi-Amirkhiz, H. Zurob, The effect of chemical patterning induced by cyclic plasticity on the formation of precipitates during aging of an Al–Mg–Si alloy. Mater Sci Eng: A. 815, 141265 (2021)

    CAS  Article  Google Scholar 

  86. Teichmann et. al., The effect of simultaneous deformation and annealing on the precipitation behaviour and mechanical properties of an Al–Mg–Si alloy Mater. Sci. Eng., A 565, 228–235 (2013)

    CAS  Article  Google Scholar 

  87. T. Saito, C.D. Marioara, J. Røyset, K. Marthinsen, R. Holmestad, The effects of quench rate and pre-deformation on precipitation hardening in Al–Mg–Si alloys with different Cu amounts. Mater Sci Eng: A. 609, 72–79 (2014)

    CAS  Article  Google Scholar 

  88. N. Kumar, S. Goel, R. Jayaganthan, G.M. Owolabi, The influence of metallurgical factors on low cycle fatigue behavior of ultra-fine grained 6082 Al alloy. Int. J. Fatigue. 110, 130–143 (2018)

    CAS  Article  Google Scholar 

  89. M.M. El-Rayes, E.A. El-Danaf, The influence of multi-pass friction stir processing on the microstructural and mechanical properties of Aluminum Alloy 6082. J Mater Process Technol. 212(5), 1157–1168 (2012)

    CAS  Article  Google Scholar 

  90. Kim et al., The nanocluster formation and vacancy behavior of step-quenched Al–Mg–Si alloy and its effect on transition to β′′-phase via advanced methods. Mater. Sci. Eng., A. 811, 141032 (2021)

    CAS  Article  Google Scholar 

  91. A. Biradar, R. Rasiwasia, J. Soni, M. Orłowska, M. Rijesh, Thermomechanical roll bonding of Al-6063 strips. J Alloys Compd. 855, 157401 (2021)

    CAS  Article  Google Scholar 

  92. Y. Birol, Thixoforging experiments with 6082 extrusion feedstock. J alloy Compd. 455, 178–185 (2008)

    CAS  Article  Google Scholar 

  93. H. Mao, Y. Kong, D. Cai, M. Yang, Y. Peng, Y. Zeng, Y. Du, β’’needle-shape precipitate formation in Al–Mg–Si alloy: Phase field simulation and experimental verification. Comput Mater Sci. 184, 109878 (2020)

    CAS  Article  Google Scholar 

  94. N. Kumar, An exploration of microstructural in-homogeneity in the 6082 Al alloy processed through room temperature multi-axial forging. Mater. Charact. 176, 111134 (2021)

    CAS  Article  Google Scholar 

  95. N. Kumar, G.M. Owolabi, R. Jayaganthan, S. Goel, Correlation of fracture toughness with microstructural features for ultrafine-grained 6082 Al alloy. Fatigue Fract Eng Mater Struct. 41(9), 1884–1899 (2018)

    CAS  Article  Google Scholar 

  96. N. Kumar, G.M. Owolabi, R. Jayaganthan, G. Warner, Effect of Annealing on Mechanical Properties and Metallurgical Factors of Ultrafine-Grained 6082 Al Alloy. Trans Indian Inst Met. 72(9), 2523–2531 (2019)

    CAS  Article  Google Scholar 

  97. N. Kumar, G.M. Owolabi, R. Jayaganthan, O.O. Ajide, S. Sonker, G. Warner, Hot-Compression Response of Solution-Treated Al–Mg–Si Alloy. J. of Materi Eng and Perform. 28, 7602–7615 (2019)

    CAS  Article  Google Scholar 

  98. S.P. Yuan, G. Liu, R.H. Wang, G.-J. Zhang, X. Pu, J. Sun, K.-H. Chen, Effect of precipitate morphology evolution on the strength–toughness relationship in Al–Mg–Si alloys. Scripta Mater. 60, 1109–1112 (2009)

    CAS  Article  Google Scholar 

  99. R.S. Yassar, D.P. Field, H. Weiland, The effect of pre-deformation on the β’’ and β’ precipitates and the role of Q’ phase in an Al–Mg–Si alloy; AA6022. Scripta Mater. 53, 299–303 (2005)

    CAS  Article  Google Scholar 

  100. D.J. Lloyd, D. Steele, J.H. Huang, Plasticity associated with grain boundaries during the bending of an Al–Mg–Si-based alloy. Scripta Mater. 63, 426–429 (2010)

    CAS  Article  Google Scholar 

  101. Y. Birol, Pre-straining to improve the bake hardening response of a twin-roll cast Al–Mg–Si alloy. Scripta Mater. 52, 169–173 (2005)

    CAS  Article  Google Scholar 

  102. S. de La Chapelle, Cube recrystallization textures in a hot deformed Al–Mg–Si alloy. Scripta Mater. 45, 1387–1391 (2001)

    Article  Google Scholar 

  103. I. Sabirov, M.R. Barnett, Y. Estrin, P.D. Hodgson, The effect of strain rate on the deformation mechanisms and the strain rate sensitivity of an ultra-fine-grained Al alloy. Scripta Mater. 61, 181–184 (2009)

    CAS  Article  Google Scholar 

  104. M.A. van Huis, J.H. Chen, M.H.F. Sluiter, H.W. Zandbergen, Phase stability and structural features of matrix-embedded hardening precipitates in Al–Mg–Si alloys in the early stages of evolution. Acta Mater. 55, 2183–2199 (2007)

    Article  CAS  Google Scholar 

  105. S. Pogatscher, H. Antrekowitsch, H. Leitner, T. Ebner, P.J. Uggowitzer, Mechanisms controlling the artificial aging of Al–Mg–Si Alloys. Acta Mater. 59, 3352–3363 (2011)

    CAS  Article  Google Scholar 

  106. P.H. Ninive, A. Strandlie, S. Gulbrandsen-Dahl, W. Lefebvre, C.D. Marioara, S.J. Andersen, O.M. Løvvik, Detailed atomistic insight into the β ″phase in Al–Mg–Si alloys. Acta mater. 69, 126–134 (2014)

    CAS  Article  Google Scholar 

  107. M.J. Starink, L.F. Cao, P.A. Rometsch, A model for the thermodynamics of and strengthening due toco-clusters in Al–Mg–Si-based alloys. Acta Mater. 60, 4194–4207 (2012)

    CAS  Article  Google Scholar 

  108. Sha et al., Strength, grain refinement and solute nanostructures of an Al–Mg–Si alloy (AA6060) processed by high-pressure torsion. Acta Mater. 63, 169–179 (2014)

    CAS  Article  Google Scholar 

  109. Q. Du, K. Tang, C.D. Marioara, S.J. Andersen, B. Holmedal, R. Holmestad, Modeling over-ageing in Al–Mg–Si alloys by a multi-phase CALPHAD-coupled Kampmann-Wagner Numerical model. Acta Mater. 122, 178–186 (2017)

    CAS  Article  Google Scholar 

  110. X. Sauvage, E.V. Bobruk, MYu. Murashkin, Y. Nasedkina, N.A. Enikeev, R.Z. Valiev, Optimization of electrical conductivity and strength combination by structure design at the nanoscale in Al–Mg–Si alloys. Acta Mater. 98, 355–366 (2015)

    CAS  Article  Google Scholar 

  111. W. Chrominski, M. Lewandowska, Precipitation phenomena in ultrafine grained Al–Mg–Si alloy with heterogeneous microstructure. Acta Mater. 103, 547–557 (2016)

    CAS  Article  Google Scholar 

  112. S. Zhu, H.-C. Shih, X. Cui, Yu. Chung-Yi, S.P. Ringer, Design of solute clustering during thermomechanical processing of AA6016 Al–Mg–Si alloy. Acta Mater. 203, 116455 (2021)

    CAS  Article  Google Scholar 

  113. P.W.J. Mckenzie, R. Lapovok, Y. Estrin, The influence of back pressure on ECAP processed AA 6016: Modeling and experiment. Acta Mater. 55, 2985–2993 (2007)

    CAS  Article  Google Scholar 

  114. S. Jana, R.S. Mishra, J.B. Baumann, G. Grant, Effect of friction stir processing on fatigue behavior of an investment cast Al–7Si–0.6 Mg alloy. Acta Mater. 58, 989–1003 (2010)

    CAS  Article  Google Scholar 

  115. S. Zhang, P.G. Mccormick, Y. Estrin, The morphology of portevin–le chatelier bands: finite element simulation for Al–Mg–Si. Acta mater. 49, 1087–1094 (2001)

    CAS  Article  Google Scholar 

  116. D. Giofre, W.A. Till Junge, M.C. Curtin, Ab initio modelling of the early stages of precipitation in Al-6000 alloys. Acta Mater. 140, 240–249 (2017)

    CAS  Article  Google Scholar 

  117. E. Mariani, E. Ghassemieh, Microstructure evolution of 6061 O Al alloy during ultrasonic consolidation: An insight from electron backscatter diffraction. Acta Mater. 58, 2492–2503 (2010)

    CAS  Article  Google Scholar 

  118. S. Shimizu, H.T. Fujii, Y.S. Sato, H. Kokawa, M.R. Sriraman, S.S. Babu, Mechanism of weld formation during very-high-power ultrasonic additive manufacturing of Al alloy 6061. Acta Mater. 74, 234–243 (2014)

    CAS  Article  Google Scholar 

  119. P.W.J. Mckenzie, R. Lapovok, ECAP with back pressure for optimum strength and ductility in aluminium alloy 6016. Part 1: Microstructure. Acta Mater. 58, 3198–3211 (2010)

    CAS  Article  Google Scholar 

  120. W. Woo, H. Choo, D.W. Brown, S.C. Vogel, P.K. Liaw, Z. Feng, Texture analysis of a friction stir processed 6061–T6 aluminum alloy using neutron diffraction. Acta Mater. 54, 3871–3882 (2006)

    CAS  Article  Google Scholar 

  121. F. Hannard, T. Pardoen, E. Maire, C. Le Bourlot, R. Mokso, A. Simar, Characterization and micromechanical modelling of microstructural heterogeneity effects on ductile fracture of 6xxx aluminium alloys. Acta Mater. 103, 558–572 (2016)

    CAS  Article  Google Scholar 

  122. P.W.J. Mckenzie, R. Lapovok, ECAP with back pressure for optimum strength and ductility in aluminium alloy 6016. Part 2: Mechanical properties and texture. Acta Mater. 58, 3212–3222 (2010)

    CAS  Article  Google Scholar 

  123. A. Simar, Y.B.A.T. Bre´chetde MeesterDenquinPardoen, Sequential modeling of local precipitation, strength and strain hardening in friction stir welds of an aluminum alloy 6005A–T6. Acta Mater. 55, 6133–6143 (2007)

    CAS  Article  Google Scholar 

  124. T. Sritharan, R.S. Chandel, Phenomena in interrupted tensile tests of heat-treated aluminium alloy 6061. Acta Mater. 45(8), 3155–3161 (1997)

    CAS  Article  Google Scholar 

  125. M. Ravi Shankar, S. Chandrasekar, A.H. King, W. Dale Compton, Microstructure and stability of nanocrystalline aluminum 6061 created by large strain machining. Acta Mater. 53, 4781–4793 (2005)

    Article  CAS  Google Scholar 

Download references

Acknowledgement

The author would like to acknowledge the financial support received under the scheme of seed grant provided to Dr. Nikhil Kumar by the IIT(BHU), Varanasi.

Author information

Authors and Affiliations

Authors

Corresponding author

Correspondence to Nikhil Kumar.

Ethics declarations

Conflict of Interest

I have no conflict of interest.

Additional information

Publisher's Note

Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.

Rights and permissions

Reprints and Permissions

About this article

Verify currency and authenticity via CrossMark

Cite this article

Kumar, N. Severe Plastic Deformation of Al–Mg–Si Alloys Processed Through Rolling Techniques: A Review. Metallogr. Microstruct. Anal. 11, 353–404 (2022). https://doi.org/10.1007/s13632-022-00859-6

Download citation

  • Received:

  • Accepted:

  • Published:

  • Issue Date:

  • DOI: https://doi.org/10.1007/s13632-022-00859-6

Keywords

  • UFG
  • Cryorolling
  • Warm rolling
  • Al–Mg–Si alloy
  • SPD