Hot Deformation Behavior and Microstructural Evolution of AA7050 Aluminum Alloy under Plane Strain Compression

Abstract

The deformation behavior, processing maps and microstructure evolution of AA7050 aluminum alloy were investigated using plane strain compression at temperatures of 300∼450°C and various strain rates ranging from 0.01 to 10 s⁻¹. Based on the flow stress and processing maps, the optimum hot working domains were established in the temperatures range of 420∼450°C and strain rates from 0.01 to 0.06 s⁻¹. The microstructure characterization of the deformed sample with the maximal power dissipation efficiency indicated that the deformation mechanism is the combined effect of continuous dynamic recrystallization (CDRX) and discontinuous dynamic recrystallization (DDRX). At low strain, dynamic recrystallization (DRX) nucleation start preferentially at triple junctions. With increasing strain, continuous dynamic recrystallization grains are developed by the progressive rotation of sub-grains and the transformation from low angle boundaries to high angle boundaries within deformed grains. At the same time, discontinuous dynamic recrystallization grains spread along grain boundaries, which is related to the local grain boundary bulging.

1. Introduction

In order to meet the rapid growth in the aerospace industry, the study of various aerospace materials is receiving a wide interest due to the cost reduction by reducing weight and extending service life [1]. Al-Zn-Mg-Cu alloys is considered as one of the most practical solutions for aerospace industries due to its excellent strength and low density [2]. However, the poor formability at room temperature and valuable processing cost become an obstacle for further application in aircraft materials for these high alloyed alloys. Therefore, different thermomechanical processing are applied to improve the workability of high strength aluminum alloy.

In recent years, many studies have been focused on the flow behavior and deformation mechanism in Al-Zn-Mg-Cu alloys by uniaxial isothermal compression tests [3–6]. During hot working, the flow stress was controlled by the result of competition between work hardening and dynamic softening. The dynamic softening of Al alloy was considered to be originated by thermal softening and microstructural softening. The thermal softening occurred mainly due to deformation heating [7]. The microstructural softening was attributed to dynamic recovery, dynamic recrystallization and dynamic precipitation [8–10]. Moreover, Recent research in 7000 alloy have attributed the deformation mechanism to deformation parameters, initial microstructure, heat treatment and hot working process [11–13]. Shen [14] and Yang [15] showed that continuous dynamic recrystallization (CDRX) is directly related to the change from low angle boundaries to high angle boundaries and the progressive rotation of sub-grains in Al-Cu-Li alloy and AA7085 aluminum alloy. Zhao [3] investigated the deformed samples of Al-Zn-Mg-Cu alloy with strain rates from 10⁻⁵ to 10⁻³ s⁻¹ and showed that discontinuous dynamic recrystallization (DDRX) occurs at grain boundaries. Rokni [16] found that geometric dynamic recrystallization (GDRX) could be obtained below 520°C. However, only a few investigations addressed the deformation mechanism in the sheet forming process of hot rolling. There has not been any attempt to understand the deformation mechanism of the alloy under different deformation conditions from a grain orientation point of view. This because that uniaxial isothermal compression cannot accurately simulate the rolling conditions for industrial production. Hot rolling is the essential step to fabricate thin and thick sheets of Al alloy. Thus, plane strain compression (PSC) is adopted to obtain available data applied frequently due to the variations in deformation conditions generally observed in industrial hot rolling schedules.

As the PSC process closely simulates rolling behavior, the production process of thin and thick sheets will be reproduced during the laboratory testing. It is helpful to better understanding the dynamic response of the metals and subsequent microstructure evolution from an industrial perspective. In the present work, the flow behavior, deformation parameter and microstructure evolution will be investigated. The optimum deformation parameter and restoration mechanism of AA7050 aluminum alloy were predicted.

2. Experimental

The alloy used in this study was received as-ingoted sheet with mainly chemical composition of 6.0Zn-2.2Mg-2.24Cu-0.10Zr-0.25Fe-0.15Si-0.10Ti-0.20Ni-0.15Cr-Al (mass%). The ingoted sheet was subjected to homogenization at 475°C for 48 h and water quenched before compression. The specimens with nominal dimensions of 20 × 15 × 10 mm³ were machined from the homogenized plate. Plane strain compression (PSC) was carried out on Gleeble-3500 simulator. Schematic diagrams experimental of PSC was presented in Fig. 1. The PSC samples were first heated to 300, 350, 400 and 450°C at a heating rate of 5°C/s, held for 180 s and subsequently compressed to a total strain of 1.2 at strain rates of 0.01, 0.1, 1 and 10 s⁻¹. After deformation, the samples were quenched to preserve the deformation microstructure. The samples were sectioned by wire-electrode cutting machine, as shown in Fig. 1. Microstructural characterization was performed on the central zone of the RD × ND plane for the PSC samples. Deformation substructure and DRX microstructure were investigated using an optical microscope (OM) complemented by an electron backscattering diffraction (EBSD) pattern analysis equipped with an HKL Channel 5. The preparation of the OM and the electron back-scattered diffraction (EBSD) samples was mechanically polished, followed by electrolytic polishing with a solution of 10% HClO₄+90% C₂H₅OH. EBSD examination was conducted by FEG-SEM with a step size of 1.0∼1.8 µm.

Fig. 1

Schematic illustration of the compressive deformation process.

3. Results and Discussion

3.1 Flow stress

Based on the isothermal compression data, effect of the deformation parameters on the flow behavior is displayed in Fig. 2. It could be noted that the flow stress increases firstly to the peak stress and then decreases to the steady-state at strain rates of 0.01∼1 s⁻¹ (Fig. 2(a)–(c)), indicating that the dynamic softening is sufficient to counteract the work-hardening during hot deformation process. At strain rate of 10 s⁻¹, all the flow stress increases rapidly to the peak stress and then decreases slowly regime up to the end of straining. The shape of true strain-true stress curves under high strain rate conditions exhibits a prominent flow softening effect (Fig. 2(d)).

Fig. 2

True stress-true strain curves of AA7050 aluminum alloy deformed under different conditions: (a) 0.01 s⁻¹, (b) 0.1 s⁻¹, (c) 1 s⁻¹ and (d) 10 s⁻¹. (online color)

The associated work hardening rate ($\theta = \partial\sigma/\partial\varepsilon$) and the relative softening (S = [(σ_p − σ_i)/σ_p] × 100%) are considered as stronger indications to estimate work hardening effect and the extent of the dynamic flow softening, respectively. S is the relative softening, σ_p the peak stress and σ_i the stress at different strain. At the initial stage of deformation (ε < 0.1), the work hardening rate is extremely high due to dislocation multi-plication and dislocation pile up for each strain rate (Fig. 3). When deformed at strain rates of 0.01∼1 s⁻¹, the work hardening rate decreases rapidly and then fluctuates around zero with increasing strain, indicating that the dynamic softening and the work-hardening reach a dynamic balance (Fig. 3(a)–(c)). Under these deformation conditions, the relative softening increase firstly and then decrease slightly after reaching maximum values, which is considered to be originated by the extent of the contribution of dynamic recovery, dynamic recrystallization and dynamic precipitation (Fig. 4(a)–(c)). Moreover, the softening extent of the 350°C samples is higher than that of the other samples at strain rate of 0.01 and 0.1 s⁻¹. This because that the thermal softening can be negligible under elevated temperature at low strain rate. At temperatures below 400°C, dynamic recovery and dynamic precipitation dominate the flow softening [7]. At 350°C, the amount of MgZn₂ phases precipitated from the Al matrix are prone to coarsening gradually. The pinning effect of dislocation become less and less effective, leading to a sudden drop in flow stress and an increase in the relative softening (Fig. 4(a)–(b)). By further rising temperature to 400°C, the amount of MgZn₂ phases dissolution. The contribution from dynamic recovery and dynamic recrystallization play a dominant role in the flow softening. Thus, the maximum relative softening could be observed at 400°C (Fig. 4(c)). With further increasing strain rate, the work hardening rate gradually decreases and eventually drops below zero, indicating that the dynamic softening is sufficient to counteract the work-hardening at medium and high strains (ε $\geqslant$ 0.25) (Fig. 3(d)). Under this condition, thermal softening may be lead to the occurrence of adiabatic shear band. In combination with the results of dynamic recovery (DRV) and dynamic recrystallization (DRX), the relative softening increase with increasing strain and reach the maximum value at 450°C (Fig. 4(d)).

Fig. 3

Work hardening rate of AA7050 aluminum alloy deformed under different conditions: (a) 0.01 s⁻¹, (b) 0.1 s⁻¹, (c) 1 s⁻¹ and (d) 10 s⁻¹. (online color)

Fig. 4

Relative softening of AA7050 aluminum alloy deformed under different conditions: (a) 0.01 s⁻¹, (b) 0.1 s⁻¹, (c) 1 s⁻¹ and (d) 10 s⁻¹. (online color)

3.2 Processing maps

According to the processing map based on the dynamic materials model (DDM), the hot deform work could be considered as a total power dissipater (P), which consists of the dissipater power content (G) through plastic deformation and dissipater power co-content (J) through microstructure transition [17, 18]. The phenomenological model can be estimated by expressed as:

  

\begin{equation} p = \sigma\dot{\varepsilon} = G + J = \int_{0}^{\dot{\varepsilon}}{\sigma d\dot{\varepsilon}} + \int_{0}^{\sigma}\dot{\varepsilon}d\sigma \end{equation} (1)

For a work-piece deformed under hot working condition, the distribution between the system power dissipation is decided by the strain rate sensitivity (m), which can be given by the following form:

  

\begin{equation} \left(\frac{\partial J}{\partial G}\right)_{\varepsilon,T}{} = \frac{\partial P}{\partial G} = \frac{\sigma\text{d}\dot{\varepsilon}}{\dot{\varepsilon}\text{d}\sigma} = \left[\frac{\partial\left(\ln\sigma \right)}{\partial\left(\ln\dot{\varepsilon} \right)}\right]_{\varepsilon,T}{} = m \end{equation} (2)

For ideally plastic flow, m = 1. The maximum value of J is obtained, J_max = P/2. The power dissipation efficiency due to structural changes in the deformation process is denoted by η. It can be defined as eq. (3).

  

\begin{equation} \eta = \frac{J}{J_{\max}} = \frac{2m}{m + 1} \end{equation} (3)

The excellent workability could not be determined by only the power dissipation efficiency maps since some flow instabilities such as flow localization, cracking, shear bands and voids, may be appear in the area with the high η value [19–21]. Therefore, the instability maps based on the extreme principle of irreversible thermodynamics paly important role to identifying the domains of flow instability. The equation below can be used to describe the unstable flow behavior:

  

\begin{equation} \xi = \left\{\frac{\partial \ln [m/(m + 1)]}{\partial\ln\dot{\varepsilon}} \right\} + m < 0 \end{equation} (4)

The corresponding three-dimensional distribution diagrams of the η changing with deformation temperature, strain rate and strain are shown in Fig. 5(a). The colorful grids represent the percentage value of power dissipation efficiency. The material with higher efficiency of power dissipation usually shows higher workability. In addition, dynamic recovery and dynamic recrystallization are more likely to occur in regions with high efficiency. The minimum values of the power dissipation for DRV and DRX are about 20% and 35%, respectively [22, 23]. From the processing maps in Fig. 5, it can be seen that the hot working domains with the high efficiency of power dissipation (η $\geqslant$ 35%) have no obvious change with the development of deformation, which occurred at temperature range 425∼450°C and strain rate range 0.01∼0.04 s⁻¹. The peak values of efficiency of power dissipations increase with increasing strain at 450°C and 0.01 s⁻¹. The maximum values are 37.5%, 38.4%, 42.7% and 43.6% with strains of 0.1, 0.4, 0.8 and 1.2, respectively. This may be that DRV and DRX consume the majority of deformation energy at an early stage of deformation. More the power input is needed to trigger microstructure transition at high strain. On the other hand, the domains with the power dissipation efficiency of 0.2∼0.35 expand from high temperature and low strain rate to low temperature and low strain rate with increasing strain (Fig. 5(a)). As the strain increased from ε = 0.1 to ε = 1.2, the deformation domains with the efficiency of about 0.2∼0.35 transform from 380∼450°C and 0.01∼0.4 s⁻¹ to 300∼425°C and 0.01∼0.88 s⁻¹. The above results can be explained by microstructure transition. At low strain, the stored energy might not be high enough to trigger microstructure transition at low temperature. As the strain increases, further deformation accelerates the motion of dislocation and the well-developed substructure. DRV can be conducted completely due to sufficient time for the dislocation climbing and crossing at low temperature and medium-low strain rate. Therefore, the efficiency of power dissipation represented dynamic recovery can transform from the domains with high temperature and low strain rate to the domains with low temperature and medium-low strain rate. In addition, it is interesting to note that DRV and DRX may be occur not only in the domains with high η value, but also inside the domains with low η value. This is because that microstructure transition is very complex. Variations of instability, such as adiabatic shear band, flow localization, dynamic strain aging, kink band, mechanical twinning and flow rotation, may be are formed in the domains with high strain rate, resulting in low η value. The complex evolution process gives rise to the various value of η. Thus, the value of η is only a reference factor to indicate microstructure transition and cannot be used to judge softening mechanism under different condition.

Fig. 5

The 3D processing maps of AA7050 aluminum alloy at different strains: (a) power dissipation efficiency map and (b) instability map. (online color)

The variations of the 3D instability maps based on the ξ of AA7050 aluminum alloy under different deformation conditions are shown in Fig. 5(b). In these figures, the wine domains and the green domains represent safe region and flow instability region, respectively. As shown in the figures, the instability region decrease initially and then slightly increase with the increase of strain. At a strain of ε = 0.1, the instability domains spread from low temperature and low strain rate to high temperature and high strain rate. The map exhibits two safe domains. The first safe domain occurs in the temperature range of 370∼450°C and strain rate range of 0.01∼0.26 s⁻¹ with a peak efficiency of 35.75%. The second safe domain occurs in the temperature range of 300∼317°C and strain rate range of 1.05∼10 s⁻¹ with a peak efficiency of 10%. The other areas are the instability regions with peak efficiency of 24%, which occurs at about 350°C and 0.01 s⁻¹. By further straining to ε = 0.4, the instability region occurred in the temperature range of 300∼340°C and strain rate range of 0.01∼0.36 s⁻¹ transforms the safe domain. The new instability region is observed at temperatures of 300∼450°C and strain rates of 0.17∼10 s⁻¹. The result is consistent with study on 7005 aluminum alloy processing map by isothermal hot compression. They attributed the cause of unstable plastic deformation to the change of the deformation mechanism [24]. With increasing strain, the instability regions deformed at 420∼450°C slightly extends toward low strain rates (0.06∼0.82 s⁻¹) initially and then reaches a relatively stable state at the strain of 1.2. This instability map characteristics suggest that the flow instability may be occur at any deformation condition.

Due to the limitation and instability of individual power dissipation diagram for optimizing process parameter, the processing map is established by superimposing the instability map on the power dissipation map considering the compensation of strain. It is confirmed that two stability domains were obtained, where the efficiency of power dissipation reach a local maximum values. Domain I occurs in the temperature range of 370∼420°C and strain rate range of 0.01∼0.32 s⁻¹ with a peak efficiency of 35% occurring at about 420°C and 0.01 s⁻¹. Domain II occurs in the temperature range of 420∼450°C and strain rate range of 0.01∼0.06 s⁻¹ with a peak efficiency of 35% occurring at about 450°C and 0.01 s⁻¹. The well-known that the higher efficiency of power dissipation, the better workability of deformed material. The optimum hot-working condition for 7050 aluminum alloy is 420∼450°C and strain rate range of 0.01∼0.06 s⁻¹.

3.3 Microstructural evolution

The typical optical macrographs under different deformation conditions are shown in Fig. 6. It can be seen that a qualitatively similar elongated grains accompanied by a great number of dispersion particles can be observed in the deformed samples from the unsafe domain (Fig. 6(a)). The maximum efficiency value is about 7.3%, which is lower than that of the efficiency value associated with DRV of about 30%. This indicates that deformed grains undergo work hardening and dynamic precipitation in these domains. Moreover, the typical microstructure flaws of flow instability, such as adiabatic shear band, severe deformation zone, flow localization, cracking and kinking, are not observed at high strain rate and low temperature. This result is not coincident with the cases observed in other study using uniaxial isothermal compression, where the typical instability microstructures were observed at high strain rate in the unsafe regions [25, 26]. This is because that the shear band propagation under plane strain compression is significantly different from those under uniaxial compression. Domanti [27] pointed out the failure of materials is directly related to additional tensile stress caused by non-uniform deformation during hot-working. During plane strain compression, the non-uniform deformation of AA7050 aluminum alloy is related to the friction and deformation heating. Previous researches showed that the contribution from frictional effect is slight under all deformation conditions, whereas deformation heating has a significant influence on the flow stress at low temperature and high strain rate [28]. High strain rate cannot offer sufficient time for homogeneous deformation. Unstable microstructures lead to the change of the flow instability factor. In this study, the temperature increase due to deformation heating during plane strain compression is lower than that of uniaxial compression [7]. The maximum values of the temperature increment are 17, 18, 19 and 10°C in the temperature range of 300∼450°C and strain rate of 10 s⁻¹, respectively (Fig. 7). Therefore, the adiabatic heating at high strain rate could not provide enough driving force for the shear band propagation. In addition, the grain shape of the 400°C/1 s⁻¹ specimen has no obvious change compared to the 350°C/10 s⁻¹ one, while the whole image seems to be more evident and clear (Fig. 6(b)). This indicates that dissolution of second-phase particles can occur at 400°C. The mechanism seems to be related to dynamic recovery and dissolution of second-phase particles. With decreasing strain rate, the microstructure exhibits more uniform than that of high strain rate (Fig. 6(c)–(d)). These specimens are located in the safe domains, where the peak efficiency of power dissipation are about 24.9% and 43.6%. This indicates that dynamic recovery and recrystallization may occur in these specimens. The microstructure deformed at 350°C and 0.01 s⁻¹ exhibits elongated grains accompanied by the majority of second-phase particles, indicating occurrence of DRV and dynamic precipitation (Fig. 6(c)). With increasing temperature, some fine grains are observed at 450°C and strain rate of 0.01 s⁻¹, which is attributed to formation of substructure and occurrence of DRX (Fig. 6(d)).

Fig. 6

Microstructures of AA7050 aluminum alloy under different deformation conditions: (a) 350°C, 10 s⁻¹, (b) 400°C, 1 s⁻¹, (c) 350°C, 0.01 s⁻¹, (d) 450°C, 0.01 s⁻¹. (online color)

Fig. 7

The actual temperatures of AA7050 aluminum alloy during hot compression: (a) 300°C, (b) 350°C, (c) 400°C and (d) 450°C. (online color)

In order to further investigate the deformation mechanism under the optimum hot working conditions, the microstructure evolution of the 450°C/0.01 s⁻¹ specimen obtained from EBSD data is shown in Fig. 8. In these figures, green, red and blue lines represent boundaries with misorientations of 2∼5°, 5∼10° and 10∼15°, respectively, and black lines represent misorientations >15°. When the specimen deformed to ε = 0.05, fine DRX grain is observed triple junctions (TJs) as marked with arrows in Fig. 8(a). The low angle grain boundaries (LAGBs) across the whole deformed grains, resulting in the orientation gradient within the deformed grains. The LAGBs fraction is 12.2% and the average misorientation is 35.7° (Fig. 9(a)). Such a microstructure evolution exhibits TJs is the most preferential DRX nucleation site at an early stage of deformation. By further straining to 0.3, more LAGBs spread along the grain boundary and invade progressively into deformed grains, leading to the formation of the complicated cell substructure and the obvious orientation gradient. Some LAGBs are transformed into high angle grain boundaries (HABs), indicating the formation of the new grains boundaries due to CDRX (Fig. 8(b)). It is seen that the fraction of LAGBs increases and reaches its maximum value, while the average misorientation decrease from 35.7° to 12.6° (Fig. 9(b)). This observation suggests that although the transformation of LABs to HABs has a profound effect on the recrystallization, the presence of sub-grains and the orientation of the deformed grains have a significant impact on recrystallization progress. At strain of ε = 0.6, the grain boundary has migrated, as evident from the features displayed by arrows in Fig. 8(c). Discontinuous recrystallization appear on the grain boundary along grain-boundary serration and bulging. The frequency of the LABs decreases to 56.7%, and the average misorientation increases to 15.8° (Fig. 9(c)). As the strain increases to 0.9, the majority of the newly recrystallized grains are formed and non-uniformly distributed. These new recrystallized grains are formed not only along the original elongated grains, but also inside recovered grains (Fig. 8(d)). The values of the LABs and the average misorientation are lower 13.5% and higher 5.9° than that of the alloy deformed at strain of ε = 0.6 (Fig. 9(d)). This indicates that the rate of the transfer of LABs into HABs is higher than the rate of generation of the LABs. As the strain increased from ε = 0.9 to ε = 1.2, more fine DRX grains are found. Some fine crystallites marked with white arrows are formed within deformed grains (Fig. 8(e)). This observation suggests that discontinuous growth of recrystallized grains at the expense of the recovered areas. In addition, the frequency of the LABs decreases to 37.5%, and the average misorientation inecreases to 24.8° (Fig. 9(e)). This information suggests that the extent of recrystallization is higher than the extent of recovery at medium-high strain. Such a microstructure evolution suggests that deformation mechanism is associated with dynamic recovery, continuous dynamic recrystallization and discontinuous dynamic recrystallization.

Fig. 8

Microstructure evolution of AA7050 aluminum alloy deformed to various strains at 450°C and 0.01 s⁻¹: (a) ε = 0.05, (b) ε = 0.3, (c) ε = 0.6, (d) ε = 0.9 and (e) ε = 1.2. (online color)

Fig. 9

The corresponding boundary misorientation distributions of AA7050 aluminum alloy deformed to various strains at 450°C and 0.01 s⁻¹: (a) ε = 0.05, (b) ε = 0.3, (c) ε = 0.6, (d) ε = 0.9 and (e) ε = 1.2. (online color)

4. Conclusions

Hot deformation behavior of AA7050 aluminum alloy was studied by plane strain compression. The processing map and microstructure evolution for the alloy were successfully developed. The main results of the current study are as follows:

1. (1)    The flow behavior exhibit different characteristics under different deformation conditions. At low strain rates (≤1 s⁻¹), the curves of the work hardening rate fluctuate around zero after the strain of 0.1, while they drop below zero in the case of high strain rate (10 s⁻¹). However, the relative softening increase firstly and then decrease slightly at low strain rates (≤1 s⁻¹), whereas they increases monotonically with increasing strain in the case of high strain rate (10 s⁻¹).
2. (2)    The processing maps of AA7050 aluminum alloy considering the compensation of the strain exhibit one safe deformation domain and one instability domain. The optimum processing parameters for hot working of AA7050 aluminum alloy are 420∼450°C and strain rate range of 0.01∼0.06 s⁻¹, where the peak efficiency of power dissipation is about 43.6%.
3. (3)    Microstructure observation in unsafe domain has no obvious microstructure flaws of flow instability. The main softening mechanism in these domains is DRV and dynamic precipitation. Microstructure observation in safe domain present dynamic recrystallization. The microstructure evolution with the peak efficiency value exhibits complex substructure and dynamic recrystallization grain, which is related to dynamic recovery, continuous dynamic recrystallization and discontinuous dynamic recrystallization.

Acknowledgments

This study is supported by the General Project of Chongqing Natural Science Foundation (cstc2021jcyj-msxmX0580).

REFERENCES

• 1)   X.S.  Zhang,  Y.J.  Chen and  J.L.  Hu: Recent advances in the development of aerospace materials, Prog. Aerosp. Sci. 97 (2018) 22–34. doi:10.1016/j.paerosci.2018.01.001
• 2)   T.  Rahmaan,  J.  Noder,  A.  Abedini,  P.  Zhou,  C.  Butcher and  M.J.  Worswick: Anisotropic plasticity characterization of 6000- and 7000-series aluminum sheet alloys at various strain rates, Int. J. Impact Eng. 135 (2020) 103390. doi:10.1016/j.ijimpeng.2019.103390
• 3)   J.H.  Zhao,  Y.L.  Deng,  J.  Tan and  J.  Zhang: Effect of strain rate on the recrystallization mechanism during isothermal compression in 7050 aluminum alloy, Mater. Sci. Eng. A 734 (2018) 120–128. doi:10.1016/j.msea.2018.07.068
• 4)   B.  Ke,  L.Y.  Ye,  J.G.  Tang,  Y.  Zhang,  S.D.  Liu,  H.Q.  Lin,  Y.  Dong and  X.D.  Liu: Hot deformation behavior and 3D processing maps of AA7020 aluminum alloy, J. Alloy. Compd. 845 (2020) 156113. doi:10.1016/j.jallcom.2020.156113
• 5)   H.R.  Rezaei Ashtiani and  P.  Shahsavari: Constitutive modeling of flow behavior of precipitation-hardened AA7022-T6 aluminum alloy at elevated temperature, Trans. Nonferrous Met. Soc. China 30 (2020) 2927–2940. doi:10.1016/S1003-6326(20)65432-2
• 6)   L.L.  Liu,  G.Q.  Zhao,  G.C.  Wang,  X.W.  Ma,  Z.H.  Yan and  S.P.  Cao: Hot deformation behavior and microstructure evolution model of 7055 aluminum alloy, J. Mater. Res. Technol. 27 (2023) 3191–3213. doi:10.1016/j.jmrt.2023.10.165
• 7)   Q.Y.  Yang,  X.Y.  Liu,  Y.X.  Liu,  X.Z.  Fan and  M.  Shu: The flow softening behavior and deformation mechanism of AA7050 aluminum alloy, Mater. Trans. 60 (2019) 2041–2047. doi:10.2320/matertrans.MT-M2019114
• 8)   H.B.  Shao,  Y.C.  Huang,  H.  Li,  Y.  Liu and  Z.B.  Xiao: Workability and dynamic restoration mechanisms of Al-6.2Zn-2.1Mg-0.15Zr alloy during medium temperature compression, J. Mater. Res. Technol. 24 (2023) 143–158. doi:10.1016/j.jmrt.2023.03.001
• 9)   Z.  Ma,  X.D.  Shu,  H.J.  Xu,  G.B.  Wang,  F.  Tong,  Z.X.  Li,  K.  Chen,  G.P.  Li and  B.  Li: Hot compression deformation behavior and microstructure evolution of Al-0.5Mg-0.4Si alloy, J. Alloy. Compd. 1002 (2024) 175282. doi:10.1016/j.jallcom.2024.175282
• 10)   S.D.  Liu,  J.H.  You,  X.M.  Zhang,  Y.L.  Deng and  Y.B.  Yuan: Influence of cooling rate after homogenization on the flow behavior of aluminum alloy 7050 under hot compression, Mater. Sci. Eng. A 527 (2010) 1200–1205. doi:10.1016/j.msea.2009.10.055
• 11)   D.G.  He,  H.  Xie,  Y.C.  Lin,  X.T.  Yan,  Z.B.  Xu and  G.  Xiao: Microstructure evolution mechanisms and a physically-based constitutive model for an Al-Zn-Mg-Cu-Zr aluminum alloy during hot Deformation, J. Mater. Res. Technol. 26 (2023) 4739–4754. doi:10.1016/j.jmrt.2023.08.203
• 12)   R.  Luo,  Y.  Cao,  H.K.  Bian,  L.L.  Chen,  C.T.  Peng,  F.Y.  Cao,  L.X.  Ouyang,  Y.  Qiu,  Y.J.  Xu,  A.  Chiba and  X.N.  Cheng: Hot workability and dynamic recrystallization behavior of a spray formed 7055 aluminum alloy, Mater. Charact. 178 (2021) 111203. doi:10.1016/j.matchar.2021.111203
• 13)   Z.  Zhang,  H.L.  Cong,  Z.J.  Yin,  B.  Qi,  Y.  Dong,  L.J.  Kong,  H.C.  Li and  K.G.  Prashanth: Thermal deformation behavior and microstructural evolution of the rapidly-solidified Al-Zn-Mg-Cu alloy in hot isostatic pressing state, J. Mater. Res. Technol. 30 (2024) 5927–5939. doi:10.1016/j.jmrt.2024.04.235
• 14)   B.  Shen,  L.  Deng and  X.Y.  Wang: A new dynamic recrystallisation model of an extruded Al-Cu-Li alloy during high-temperature deformation, Mater. Sci. Eng. A 625 (2015) 288–295. doi:10.1016/j.msea.2014.11.095
• 15)   Q.Y.  Yang,  Z.H.  Deng,  Z.Q.  Zhang,  Q.  Liu,  Z.H.  Jia and  G.J.  Huang: Effects of strain rate on flow stress behavior and dynamic recrystallization mechanism of Al-Zn-Mg-Cu aluminum alloy during hot deformation, Mater. Sci. Eng. A 662 (2016) 204–213. doi:10.1016/j.msea.2016.03.027
• 16)   M.R.  Rokni,  A.  Zarei-Hanzaki,  A.A.  Roostaei and  H.R.  Abedi: An investigation into the hot deformation characteristics of 7075 aluminum alloy, Mater. Des. 32 (2011) 2339–2344. doi:10.1016/j.matdes.2010.12.047
• 17)   A.  Ghosh,  A.  Elasheri,  N.  Parson and  X.G.  Chen: Hot deformation behavior and processing maps for an Al-Mg-Si-Zr-Mn alloy, J. Alloy. Metall. Syst. 6 (2024) 100077. doi:10.1016/j.jalmes.2024.100077
• 18)   Y.  Sun,  Z.H.  Cao,  Z.P.  Wan,  L.X.  Hu,  W.H.  Ye,  N.K.  Li and  C.L.  Fan: 3D processing map and hot deformation behavior of 6A02 aluminum alloy, J. Alloy. Compd. 742 (2018) 356–368. doi:10.1016/j.jallcom.2018.01.299
• 19)   Y.B.  Yang,  Z.M.  Zhang,  X.B.  Li,  Q.  Wang and  Y.H.  Zhang: The effects of grain size on the hot deformation and processing map for 7075 aluminum alloy, Mater. Des. 51 (2013) 592–597. doi:10.1016/j.matdes.2013.04.034
• 20)   H.  Wang,  D.F.  Yin,  M.C.  Zhao,  Y.  Tian and  A.  Atrens: The hot deformation characteristics and the associated microstructural evolution of an Al-Cu-Li-Zn alloy studied by constitutive equations and processing maps, J. Mater. Res. Technol. 27 (2023) 401–418. doi:10.1016/j.jmrt.2023.09.222
• 21)   N.  Bayat-Tork,  R.  Mahmudi and  M.M.  Hoseini-Athar: Hot deformation constitutive analysis and processing maps of extruded Mg-Gd binary alloys, J. Mater. Res. Technol. 9 (2020) 15346–15359. doi:10.1016/j.jmrt.2020.10.101
• 22)   L.  Ou,  Z.Q.  Zheng,  Y.F.  Nie and  H.G.  Jian: Hot deformation behavior of 2060 alloy, J. Alloy. Compd. 648 (2015) 681–689. doi:10.1016/j.jallcom.2015.07.027
• 23)   M.  Seyed-Salehi,  B.H.  Kim,  S.Y.  Yang,  S.K.  Kim and  G.E.  Bozchaloei: Effect of Cr elimination on flow behavior and processing map of newly developed ECO-7175 aluminum alloy during hot compression, Trans. Nonferrous Met. Soc. China 32 (2022) 1442–1459. doi:10.1016/S1003-6326(22)65885-0
• 24)   C.N.  Lin,  Y.C.  Tzeng,  S.L.  Lee,  Y.K.  Fuh,  A.  Łukaszek-Sołek,  C.Y.  Lin,  M.C.  Chen and  T.A.  Pan: Optimization of hot deformation processing parameters for as-extruded 7005 alloys through the integration of 3D processing maps and FEM numerical simulation, J. Alloy. Compd. 948 (2023) 169804. doi:10.1016/j.jallcom.2023.169804
• 25)   H.  Yi,  J.  Ding,  C.G.  Ni,  J.H.  Dai,  Y.  Tang,  X.G.  Chen,  K.H.  Song and  X.C.  Xia: Hot compression deformation behavior and processing maps of Al-0.5Mg-0.4Si-0.1Cu alloy, J. Mater. Res. Technol. 19 (2022) 4890–4904. doi:10.1016/j.jmrt.2022.07.027
• 26)   C.G.  Qin,  B.  Jiang,  M.L.  Hu,  Y.  Wang,  H.Y.  Xu,  Y.  Guo and  Z.S.  Ji: Hot compression deformation behavior and processing maps of Al-0.5Mg-0.4Si-0.1Cu alloy, J. Mater. Res. Technol. 33 (2024) 4725–4738. doi:10.1016/j.jmrt.2024.10.163
• 27)   A.T.J.  Domanti,  D.J.  Horrobin and  J.  Bridgwater: An investigation of fracture criteria for predicting surface fracture in paste extrusion, Int. J. Mech. Sci. 44 (2002) 1381–1410. doi:10.1016/S0020-7403(02)00047-4
• 28)   Q.Y.  Yang,  X.Y.  Liu and  Y.S.  Cheng: Hot Deformation Behavior of 7085 Aluminum Alloy Based on Deformation Heating and Friction, Mater. Trans. 61 (2020) 318–325. doi:10.2320/matertrans.MT-M2019279