Impact of Aging Treatment on Microstructure and Performance of Al-Zn-Mg-Cu Alloy Thin Sheets

: Al-Zn-Mg-Cu alloy, recognized for its heat-treatable and super-high strength, is a pivotal material for critical structural components in aviation. Consequently, grasping the intricate relationship between heat treatment processes and the microstructural characteristics and properties of the Al-Zn-Mg-Cu alloy is of paramount practical signiﬁcance. This study encompasses a spectrum of heat treatment methodologies, delving into the diverse effects of distinct aging treatments on the microstructure and performance of thin Al-Zn-Mg-Cu alloy sheets. The ﬁndings illuminate that short natural aging treatments bestow heightened strength and hardness upon the alloy, albeit at the expense of ductility. Nevertheless, the alloy retains commendable malleability, enabling signiﬁcant cold deformation. In contrast, artiﬁcial aging markedly elevates the alloy’s strength and hardness, albeit concurrently inducing a precipitous decline in ductility, thereby attenuating the alloy’s cold formability.


Introduction
The utilization of Al-Zn-Mg-Cu alloy in aviation for critical structural components is attributed to its remarkable characteristics, including high strength, low density, and excellent fracture toughness.The alloy's elevated alloying element content and the abundance of nanoscale precipitates resulting from heat treatment contribute to its impressive mechanical performance, albeit at the expense of reduced processability [1,2].Consequently, a pivotal area of research revolves around enhancing the alloy's cold formability while concurrently maintaining its strength levels [3].
The strengthening effect of Al-Zn-Mg-Cu alloy is mainly achieved via aging treatment [4].During the aging treatment process, a large number of fine precipitates will form inside the alloy, which can interact with dislocations, hinder their movement, and thus produce a strengthening effect.Previous studies have shown that the precipitation sequence of the main strengthening phases in the Al-Zn-Mg-Cu alloy series is: supersaturated solid solution (SSSS)-Guinier-Preston zone (GP zone)-η'-η [5][6][7][8].Correspondingly, there has been developed various distinct categories of artificial aging, such as peak aging (T6) [9,10], two-stage overaging (T7X) [11,12], and retrogression and reaging (RRA) [13], which aim to meet diverse performance requirements.After T6 treatment, a large number of nanoscale precipitates η' are dispersed inside the Al-Zn-Mg-Cu alloy, causing the alloy with the highest strength [14,15], but correspondingly the plasticity and anti-corrosion performance are poor [9].In order to solve the problem of poor fracture toughness and corrosion resistance of T6 state alloy, a two-stage overaging process was proposed.In the first stage of two-stage aging, a large number of fine and uniformly distributed precipitates are formed in the matrix.Then, in the second aging stage, the precipitates rapidly coarsen in a high-temperature environment and exhibit a discontinuous distribution state [16], which effectively reduces the sensitivity to stress corrosion cracking of the alloy and improves fracture toughness at the cost of losing a certain amount of strength [11,12].The RRA process was proposed to compensate for the shortcomings of the T6 and T7X processes, which can balance high strength and good corrosion resistance, but the process is more complex and has a longer cycle [17][18][19].In recent years, people have proposed thermomechanical treatment processes to reduce cycle time and achieve better comprehensive performance.This process can effectively combine deformation strengthening and heat treatment strengthening.For example, deformation accelerates the aging process, while precipitates delay the recovery of dislocations, leaving more work hardening inside the alloy [20], enabling the alloy to achieve good comprehensive properties.
Previous studies have shown that most strengthening methods for aluminum alloys come at the cost of sacrificing a certain degree of ductility to achieve high strength.For structures requiring heightened ductility, people often employ approaches such as microalloying and refining processing techniques to finely adjust the microstructure, striking a balance between strength and ductility.Research has demonstrated that incorporating Sc and Zr elements into Al-Zn-Mg-Cu alloys can effectively refine grain size and maintain consistent deformation and recrystallization behavior during forming and heat treatment [21,22].Jiang et al. [23] revealed that judiciously increasing Zn content and introducing Sc and Zr elements into the 7000 series aluminum alloys can transcend the strength-ductility trade-off.However, microalloying can elevate production costs, especially with rare earth elements.Hence, a pivotal research focus for Al-Zn-Mg-Cu alloys lies in achieving high strength while preserving excellent formability using modifications in thermal and heat treatment processes.Zuo et al. [24] put forth an enhanced two-step hot rolling process, termed Dual Heat Rolling (DHR), which encompasses pre-deformation, brief intermediate annealing, and final hot rolling to yield refined grains, consequently enhancing the alloy's ductility and corrosion resistance.Nonetheless, literature addressing the influence of heat treatment processes on Al-Zn-Mg-Cu alloy's formability remains relatively scarce.
The main material of the LNG Open Rack Vaporizer (ORV) is aluminum alloy, with a large number of aluminum alloy support components inside, which have certain requirements for the strength and plasticity of the aluminum alloy.This study employs a cold-rolled thin sheet of Al-Zn-Mg-Cu alloy designed for aviation applications as the experimental material.It investigates the impacts of different aging treatments on the alloy's mechanical properties and ductility.The aim is to comprehend the mechanisms by which heating treatments influence both the strength and ductility of the alloy, thus providing technical support for the production of structural components for the LNG Open Rack Vaporizers.

Experimental Methods
This experiment employs a 2.0 mm-thick Al-Zn-Mg-Cu alloy sheet provided using a specific enterprise.The initial state of the sheet is O state (annealed state).The primary chemical composition is as follows: 6.21% Zn, 1.82% Cu, 2.16% Mg, 0.39% Mn, 0.11% Fe, and 0.08% Si (wt%).Initially, the sheet undergoes a solution treatment at 470 • C for 38 min, and the solution-treated samples were then water quenched to room temperature.Subsequently, ten groups of samples undergo single-stage aging treatment, with the process conditions depicted in Figure 1.Furthermore, 16 groups of samples following solution treatment experienced a two-stage aging treatment, with the experimental plan outlined in Table 1.Additionally, four groups of samples after solution treatment are subjected to natural aging for durations of 0.5 h, 1 h, 2 h, and 4 h, respectively.bending angle of 80.8°.Each sample group underwent three repeated tests, and the average value was used as the experimental result.Hardness testing was conducted on the cross-section of the specimens using an HV-10Z micro-Vickers hardness tester (SFMIT, Changzhou, China), with a testing load of 29.42 N and a load-holding time of 10 s.Additionally, samples from the deformed region after bending were observed for their microstructures using the electron backscatter diffraction (EBSD) technique.

Microstructure and Mechanical Performance of Al-Zn-Mg-Cu Alloy Annealed Thin Sheets
Figure 2 illustrates the microstructure along the longitudinal section of the annealed thin sheet.As depicted, the internal grain structure of the aluminum alloy thin sheet exhibits a relatively flattened equiaxed morphology with consistent grain sizes.The mechanical properties of the annealed samples are listed in Table 2.The sheet's yield strength is 102.5 MPa, tensile strength is 219.5 MPa, elongation is 21.28%, and the hardness is 65.46 ± 3.5 HV.

Solution Treatment
The First Stage of Aging Treatment The Second Stage of Aging Treatment Temperature Time Optical microscopy (OM) was employed to observe the microstructures of samples in the annealed state after anodization and samples following solution treatment.The anodization solution consisted of 5 mL BF 3 + 100 mL H 2 O, with a voltage of 20 V and a current of 0.5-0.7 A. Several ϕ3 mm small discs were extracted from the samples of T6 and T74 state, thinned to 70 µm using sandpaper, and subsequently subjected to electrolytic twin-jet polishing before being observed under transmission electron microscopy (TEM, FEI, Hillsboro, OR, America) to analyze the morphology of the precipitated phases.Standard tensile samples were wire cut along the rolling direction line on the heat-treated thin plate.Additionally, the tensile tests were performed on the various experimental sample groups using a CMT5105 universal testing machine (MTS, Jinan, China), with each group subjected to three repeated tests and the average value used as the experimental result.The heattreated samples underwent bending tests on the same machine with a controlled bending angle of 80.8 • .Each sample group underwent three repeated tests, and the average value was used as the experimental result.Hardness testing was conducted on the cross-section of the specimens using an HV-10Z micro-Vickers hardness tester (SFMIT, Changzhou, China), with a testing load of 29.42 N and a load-holding time of 10 s.Additionally, samples from the deformed region after bending were observed for their microstructures using the electron backscatter diffraction (EBSD) technique.

Microstructure and Mechanical Performance of Al-Zn-Mg-Cu Alloy Annealed Thin Sheets
Figure 2 illustrates the microstructure along the longitudinal section of the annealed thin sheet.As depicted, the internal grain structure of the aluminum alloy thin sheet exhibits a relatively flattened equiaxed morphology with consistent grain sizes.The mechanical properties of the annealed samples are listed in Table 2.The sheet's yield   As the natural aging time extends, the strength and hardness of the samples gradually increase while the elongation rate diminishes rapidly.At a natural aging time of 3 h, the tensile strength rises to 367 MPa, the yield strength reaches 179 MPa, and the elongation drops to 18.31%.The images in Figure 4 depict the macroscopic appearance of samples after bending deformation following various durations of natural aging, along with the corresponding springback angles.Notably, it is observed that when the natural aging time remains within 3 h, there are no noticeable cracks on the sample surfaces after the bending deformation.Furthermore, measurements indicate that as the natural aging time is extended, the springback angle of the samples initially decreases before subsequently increasing.
For instance, at a natural aging time of 1 h, the springback angle of the samples measures

Effects of Natural Aging on the Performance of Al-Zn-Mg-Cu Alloy Thin Sheets
Figure 3 illustrates the changing trends in tensile performance, hardness, and bending behavior of the thin sheet samples as a function of natural aging time.Within the first 0.5 h following solution treatment, the samples exhibit a tensile strength of 341 MPa, a yield strength of 153 MPa, an elongation of 20.90%, and a Vickers hardness of 84.52 ± 3.2.As the natural aging time extends, the strength and hardness of the samples gradually increase while the elongation rate diminishes rapidly.At a natural aging time of 3 h, the tensile strength rises to 367 MPa, the yield strength reaches 179 MPa, and the elongation drops to 18.31%.The images in Figure 4 depict the macroscopic appearance of samples after bending deformation following various durations of natural aging, along with the corresponding springback angles.Notably, it is observed that when the natural aging time remains within 3 h, there are no noticeable cracks on the sample surfaces after the bending deformation.Furthermore, measurements indicate that as the natural aging time is extended, the springback angle of the samples initially decreases before subsequently increasing.
For instance, at a natural aging time of 1 h, the springback angle of the samples measures only 1.6 ± 0.1°.The images in Figure 4 depict the macroscopic appearance of samples after bending deformation following various durations of natural aging, along with the corresponding springback angles.Notably, it is observed that when the natural aging time remains within 3 h, there are no noticeable cracks on the sample surfaces after the bending deformation.Furthermore, measurements indicate that as the natural aging time is extended, the springback angle of the samples initially decreases before subsequently increasing.For instance, at a natural aging time of 1 h, the springback angle of the samples measures only 1.6 ± 0.1  Figure 5 presents the microstructure of the deformation zone in the sample (designated as T4N1) after 1 h of natural aging and subsequent bending deformation.As observed in the image, the grains on the RD-ND plane of the sample exhibit an overall fusiform shape, with certain grains experiencing fragmentation during the bending process.Figure 5b reveals the presence of numerous subgrain boundaries (<2°) and high-angle grain boundaries within the sample's interior, while the occurrence of low-angle grain boundaries (2°~15°) is quite limited.This signifies that the grains in the deformation zone undergo distortion during the bending process, leading to the formation of a substantial quantity of subgrain boundaries due to the relatively moderate extent of deformation.

Impact of Single-Stage Aging on Microstructure and Properties of Al-Zn-Mg-Cu Alloy Thin Sheets
Figure 6 illustrates the patterns of change in tensile performance and hardness during the single-stage aging of Al-Zn-Mg-Cu alloy thin sheets across different aging temperatures.As depicted, the variation in tensile strength remains relatively minimal within the temperature range of 100 °C to 130 °C.However, a noticeable decrease in tensile strength occurs after aging treatment at 140 °C (a reduction from 556 MPa at 120 °C to 537 MP).Conversely, the changes in yield strength and elongation are more pronounced with increasing aging temperature.With the rising temperature, the yield strength initially increases before showing a subsequent decline, while the elongation rate initially drops before increasing again.At an aging temperature of 120 °C, the yield strength (481 MPa) reaches the peak value, coinciding with the lowest elongation rate (12.55%).From Figure 6b, it can be seen that the variation in hardness of the thin sheets during aging treatment Figure 5 presents the microstructure of the deformation zone in the sample (designated as T4N1) after 1 h of natural aging and subsequent bending deformation.As observed in the image, the grains on the RD-ND plane of the sample exhibit an overall fusiform shape, with certain grains experiencing fragmentation during the bending process.Figure 5b reveals the presence of numerous subgrain boundaries (<2 • ) and high-angle grain boundaries within the sample's interior, while the occurrence of low-angle grain boundaries (2 • ~15 • ) is quite limited.This signifies that the grains in the deformation zone undergo distortion during the bending process, leading to the formation of a substantial quantity of subgrain boundaries due to the relatively moderate extent of deformation.Figure 5 presents the microstructure of the deformation zone in the sample (designated as T4N1) after 1 h of natural aging and subsequent bending deformation.As observed in the image, the grains on the RD-ND plane of the sample exhibit an overall fusiform shape, with certain grains experiencing fragmentation during the bending process.Figure 5b reveals the presence of numerous subgrain boundaries (<2°) and high-angle grain boundaries within the sample's interior, while the occurrence of low-angle grain boundaries (2°~15°) is quite limited.This signifies that the grains in the deformation zone undergo distortion during the bending process, leading to the formation of a substantial quantity of subgrain boundaries due to the relatively moderate extent of deformation.

Impact of Single-Stage Aging on Microstructure and Properties of Al-Zn-Mg-Cu Alloy Thin Sheets
Figure 6 illustrates the patterns of change in tensile performance and hardness during the single-stage aging of Al-Zn-Mg-Cu alloy thin sheets across different aging temperatures.As depicted, the variation in tensile strength remains relatively minimal within the temperature range of 100 °C to 130 °C.However, a noticeable decrease in tensile strength occurs after aging treatment at 140 °C (a reduction from 556 MPa at 120 °C to 537 MP).Conversely, the changes in yield strength and elongation are more pronounced with increasing aging temperature.With the rising temperature, the yield strength initially increases before showing a subsequent decline, while the elongation rate initially drops before increasing again.At an aging temperature of 120 °C, the yield strength (481 MPa) reaches the peak value, coinciding with the lowest elongation rate (12.55%).From Figure 6b, it can be seen that the variation in hardness of the thin sheets during aging treatment between 100 °C and 140 °C is relatively small, fluctuating within the range of 193.7 HV to     7a, it can be seen that with the extension of aging time, both the yield strength and tensile strength of the alloy remain basically unchanged, while the elongation shows a trend of gradually decreasing and then increasing.From Figure 7b, it can be seen that the hardness of the thin plate shows a trend of first increasing and then gradually decreasing with the extension of aging time, but the change amplitude is relatively small.Based on these experimental results, it can be concluded that, compared to aging time, aging temperature has a more significant influence on the alloy's overall properties.Figure 8 displays the macroscopic appearance and microstructure at the fracture surface of the sample (T6N2) following the process of solution treatment at 470 °C for 38 min, followed by aging treatment at 120 °C for 24 h.In Figure 8a, it is evident that pronounced cracks have emerged in the region of bending deformation in the sample.These cracks are situated precisely along the ridge line of the bending angle, exhibiting a fairly straight configuration that traverses the entire width of the sample.Figure 8b reveals distinct grain deformation in the bending deformation zone.The grains on the outer perimeter of the bending deformation zone (left section in Figure 8b) have elongated, as shown in Figure 8c.On the contrary, the grains inside the bending deformation zone have been com-  From Figure 7a, it can be seen that with the extension of aging time, both the yield strength and tensile strength of the alloy remain basically unchanged, while the elongation shows a trend of gradually decreasing and then increasing.From Figure 7b, it can be seen that the hardness of the thin plate shows a trend of first increasing and then gradually decreasing with the extension of aging time, but the change amplitude is relatively small.Based on these experimental results, it can be concluded that, compared to aging time, aging temperature has a more significant influence on the alloy's overall properties.7a, it can be seen that with the extension of aging time, both the yield strength and tensile strength of the alloy remain basically unchanged, while the elongation shows a trend of gradually decreasing and then increasing.From Figure 7b, it can be seen that the hardness of the thin plate shows a trend of first increasing and then gradually decreasing with the extension of aging time, but the change amplitude is relatively small.Based on these experimental results, it can be concluded that, compared to aging time, aging temperature has a more significant influence on the alloy's overall properties.Figure 8 displays the macroscopic appearance and microstructure at the fracture surface of the sample (T6N2) following the process of solution treatment at 470 °C for 38 min, followed by aging treatment at 120 °C for 24 h.In Figure 8a, it is evident that pronounced cracks have emerged in the region of bending deformation in the sample.These cracks are situated precisely along the ridge line of the bending angle, exhibiting a fairly straight configuration that traverses the entire width of the sample.Figure 8b reveals distinct grain deformation in the bending deformation zone.The grains on the outer perimeter of the bending deformation zone (left section in Figure 8b) have elongated, as shown in Figure 8c.On the contrary, the grains inside the bending deformation zone have been compressed, presenting an equiaxed shape, but the degree of grain deformation is relatively Figure 8 displays the macroscopic appearance and microstructure at the fracture surface of the sample (T6N2) following the process of solution treatment at 470 • C for 38 min, followed by aging treatment at 120 • C for 24 h.In Figure 8a, it is evident that pronounced cracks have emerged in the region of bending deformation in the sample.These cracks are situated precisely along the ridge line of the bending angle, exhibiting a fairly straight configuration that traverses the entire width of the sample.Figure 8b reveals distinct grain deformation in the bending deformation zone.The grains on the outer perimeter of the bending deformation zone (left section in Figure 8b) have elongated, as shown in Figure 8c.On the contrary, the grains inside the bending deformation zone have been compressed, presenting an equiaxed shape, but the degree of grain deformation is relatively small, as shown in Figure 8d.This observation underscores the subpar bending performance of the sample treated with the T6 treatment.

Impact of Dual-Stage Aging Treatment on Microstructure and Properties of Al-Zn-Mg-Cu Alloy Thin Sheets
Figure 9 illustrates the curves depicting the changes in tensile performance and hardness during the dual-stage aging treatment of Al-Zn-Mg-Cu alloy thin sheets concerning the temperature and duration of the second-stage aging.As observed in the figure, as the temperature of the second-stage aging rises, there is a noticeable reduction in the strength and hardness of the samples, accompanied by a significant increase in the elongation rate.Furthermore, extending the duration of the second-stage aging results in a gradual decline in both strength and hardness while the elongation rate remains relatively stable.This phenomenon can be attributed to the different morphologies of precipitates formed during the distinct stages of the two-stage aging treatment for Al-Zn-Mg-Cu alloy.During the initial phase of low-temperature aging in the first stage, the alloy's grain interiors experience the development of numerous fine and dispersed GP zones or coherent precipitate phases.Subsequently, in the second stage, characterized by higher temperature and prolonged aging, the precipitate particles formed during the first stage transform into semicoherent η' precipitates or equilibrium η precipitates [25,26].Consequently, in the secondstage aging process, the size and quantity of precipitates within the alloy undergo changes influenced by aging temperature and duration, leading to variations in the alloy's strength and hardness with changing aging conditions.Moreover, a comparative analysis highlights that aging temperature exerts a more pronounced influence on the alloy's mechanical properties.As observed in the figure, as the temperature of the second-stage aging rises, there is a noticeable reduction in the strength and hardness of the samples, accompanied by a significant increase in the elongation rate.Furthermore, extending the duration of the second-stage aging results in a gradual decline in both strength and hardness while the elongation rate remains relatively stable.This phenomenon can be attributed to the different morphologies of precipitates formed during the distinct stages of the two-stage aging treatment for Al-Zn-Mg-Cu alloy.During the initial phase of low-temperature aging in the first stage, the alloy's grain interiors experience the development of numerous fine and dispersed GP zones or coherent precipitate phases.Subsequently, in the second stage, characterized by higher temperature and prolonged aging, the precipitate particles formed during the first stage transform into semi-coherent η' precipitates or equilibrium η precipitates [25,26].Consequently, in the second-stage aging process, the size and quantity of precipitates within the alloy undergo changes influenced by aging temperature and duration, leading to variations in the alloy's strength and hardness with changing aging conditions.Moreover, a comparative analysis highlights that aging temperature exerts a more pronounced influence on the alloy's mechanical properties.Figure 10 illustrates the macroscopic appearance and microstructure at the fracture surface of the sample (T74N3) after undergoing the process of solution treatment at 470 °C for 38 min, followed by two-stage aging treatments at 115 °C for 8 h and 165 °C for 16 h.In Figure 10a, it is evident that prominent cracks have developed in the region of bending deformation within the sample.These cracks align precisely with the ridge line of the bending angle, displaying a relatively straight configuration that spans the entire width of the sample.Additionally, as depicted in Figure 10b, the microstructure in the bending deformation zone showcases marked grain deformation.The grains on the outer side of the bending deformation zone (left side in Figure 10b) are noticeably elongated, resembling a fibrous pattern, while those on the inner side are compressed, adopting a polygonal shape.Noteworthy is the comparison with the morphology of the T6N2 sample after bending, as shown in Figure 8.This reveals that the T74N3 sample displays a more pronounced level of grain deformation.This observation suggests that, in comparison to the forming performance of the thin sheet after the T6 treatment, the forming performance of the thin sheet following the T74 treatment has been moderately enhanced, although it still falls short of achieving a significant degree of bending deformation.In Figure 10a, it is evident that prominent cracks have developed in the region of bending deformation within the sample.These cracks align precisely with the ridge line of the bending angle, displaying a relatively straight configuration that spans the entire width of the sample.Additionally, as depicted in Figure 10b, the microstructure in the bending deformation zone showcases marked grain deformation.The grains on the outer side of the bending deformation zone (left side in Figure 10b) are noticeably elongated, resembling a fibrous pattern, while those on the inner side are compressed, adopting a polygonal shape.Noteworthy is the comparison with the morphology of the T6N2 sample after bending, as shown in Figure 8.This reveals that the T74N3 sample displays a more pronounced level of grain deformation.This observation suggests that, in comparison to the forming performance of the thin sheet after the T6 treatment, the forming performance of the thin sheet following the T74 treatment has been moderately enhanced, although it still falls short of achieving a significant degree of bending deformation.

Discussion
As commonly recognized, solid solution strengthening, grain boundary strengthening, precipitation strengthening, and dislocation strengthening stand as the major mechanisms for enhancing the strength of Al-Zn-Mg-Cu alloys [27].It is generally understood that the yield strength of aluminum alloys can be expressed as follows [28]: where  is the yield strength of pure aluminum, ∆ , ∆ , ∆ , and ∆ represent the strength increments attributed to grain boundaries, dislocations, solid solution, and precipitation, respectively.Precipitation strengthening, in particular, serves as the predominant factor in elevating the strength of Al-Zn-Mg-Cu alloys [27,28].The reinforcement by precipitation phases primarily revolves around hindering the movement of dislocations through various interaction mechanisms.The central mechanism of precipitation strengthening involves the interference of precipitation phases with dislocation mobility, accomplished using a range of interaction mechanisms.Depending on the properties of the precipitate phase and its crystallographic orientation relationship with the matrix, diverse interactions may take place between the precipitate particles and dislocationsthese interactions could encompass either dislocation cutting through or bypassing the precipitate phase.For small-sized precipitate phases that exhibit complete coherence with the matrix, the dislocation shearing mechanism comes into effect.However, when the dimensions of coherent precipitate phases surpass a critical threshold (2 nm), or if the precipitate phase is incoherent, the strengthening mechanism transitions toward dislocation bypassing [29].
Studies have shown that in Al-Zn-Mg-Cu alloys treated with relatively short natural aging durations, the predominant form of precipitation phases is in the shape of GP zones.These structural entities have small dimensions, typically only a few nanometers, and maintain a fully coherent relationship with the Al matrix.As a result, during deformation, the mobile dislocations tend to shear through them.The strengthening effect of the alloy is influenced by the interplay between the size and volume fraction of the strengthening phases [30], as outlined below: where c represents the alloy constant, φ and r represent the volume fraction and size of the strengthening phase, respectively.The strengthening effect of the alloy is enhanced as both the volume fraction and size increase.
During the initial stages of natural aging, with the passage of time, solute atoms continuously precipitate from the matrix.This leads to an increasing number of GP zones within the matrix, which gradually grow in size, thereby amplifying the alloy's strengthening effect.Consequently, in this study, samples subjected to solution treatment exhib-

Discussion
As commonly recognized, solid solution strengthening, grain boundary strengthening, precipitation strengthening, and dislocation strengthening stand as the major mechanisms for enhancing the strength of Al-Zn-Mg-Cu alloys [27].It is generally understood that the yield strength of aluminum alloys can be expressed as follows [28]: where σ 0 is the yield strength of pure aluminum, ∆σ gb , ∆σ d , ∆σ ss , and∆σ ppt represent the strength increments attributed to grain boundaries, dislocations, solid solution, and precipitation, respectively.Precipitation strengthening, in particular, serves as the predominant factor in elevating the strength of Al-Zn-Mg-Cu alloys [27,28].The reinforcement by precipitation phases primarily revolves around hindering the movement of dislocations through various interaction mechanisms.The central mechanism of precipitation strengthening involves the interference of precipitation phases with dislocation mobility, accomplished using a range of interaction mechanisms.Depending on the properties of the precipitate phase and its crystallographic orientation relationship with the matrix, diverse interactions may take place between the precipitate particles and dislocations-these interactions could encompass either dislocation cutting through or bypassing the precipitate phase.For small-sized precipitate phases that exhibit complete coherence with the matrix, the dislocation shearing mechanism comes into effect.However, when the dimensions of coherent precipitate phases surpass a critical threshold (2 nm), or if the precipitate phase is incoherent, the strengthening mechanism transitions toward dislocation bypassing [29].
Studies have shown that in Al-Zn-Mg-Cu alloys treated with relatively short natural aging durations, the predominant form of precipitation phases is in the shape of GP zones.These structural entities have small dimensions, typically only a few nanometers, and maintain a fully coherent relationship with the Al matrix.As a result, during deformation, the mobile dislocations tend to shear through them.The strengthening effect of the alloy is influenced by the interplay between the size and volume fraction of the strengthening phases [30], as outlined below: where c represents the alloy constant, ϕ and r represent the volume fraction and size of the strengthening phase, respectively.The strengthening effect of the alloy is enhanced as both the volume fraction and size increase.
During the initial stages of natural aging, with the passage of time, solute atoms continuously precipitate from the matrix.This leads to an increasing number of GP zones within the matrix, which gradually grow in size, thereby amplifying the alloy's strengthening effect.Consequently, in this study, samples subjected to solution treatment exhibited a progressive increase in both strength and hardness during the natural aging process spanning from 0.5 h to 3 h.Concurrently, due to the inhibiting effect of GP zones on dislocations, the ductile deformation of the alloy becomes more challenging, resulting in decreased ductility.Thus, as the duration of natural aging extends, the elongation rate of the samples noticeably diminishes.
Figure 11 illustrates TEM bright-field images showcasing the precipitate phases in samples subjected to T6 and T74 treatments within this study.Within the images, a plethora of plate-like precipitate phases can be observed, uniformly dispersed throughout the aluminum matrix.Due to their relatively larger dimensions, these precipitates pose a challenge for dislocations to shear through during deformation.Consequently, dislocations tend to bypass these precipitate phases, with the alloy's increased strength primarily attributed to Orowan strengthening [27].This strengthening effect can be quantified as [31,32]: where M denotes the Taylor factor, G represents the shear modulus of the Al matrix, b represents the Burgers vector, D p represents the average radius of the precipitate phase, f v signifies the volume fraction of the precipitate phase.Equation (3) illustrates that a higher number of precipitate phases and smaller sizes result in a more pronounced strengthening effect of the precipitates on the alloy.
Metals 2023, 13, x FOR PEER REVIEW 10 of 13 spanning from 0.5 h to 3 h.Concurrently, due to the inhibiting effect of GP zones on dislocations, the ductile deformation of the alloy becomes more challenging, resulting in decreased ductility.Thus, as the duration of natural aging extends, the elongation rate of the samples noticeably diminishes.Figure 11 illustrates TEM bright-field images showcasing the precipitate phases in samples subjected to T6 and T74 treatments within this study.Within the images, a plethora of plate-like precipitate phases can be observed, uniformly dispersed throughout the aluminum matrix.Due to their relatively larger dimensions, these precipitates pose a challenge for dislocations to shear through during deformation.Consequently, dislocations tend to bypass these precipitate phases, with the alloy's increased strength primarily attributed to Orowan strengthening [27].This strengthening effect can be quantified as [31,32]: where M denotes the Taylor factor, G represents the shear modulus of the Al matrix, b represents the Burgers vector,  represents the average radius of the precipitate phase, f signifies the volume fraction of the precipitate phase.Equation (3) illustrates that a higher number of precipitate phases and smaller sizes result in a more pronounced strengthening effect of the precipitates on the alloy.During the artificial aging of Al-Zn-Mg-Cu alloy, precipitation follows the sequence of αAl-GP zone-η'-η.The final types, quantities, and sizes of precipitate phases are closely linked to the aging temperature and time.Generally, within a specific temperature range, higher aging temperatures lead to faster precipitation rates of solute atoms and increased quantities of precipitate phases.However, excessively high temperatures during prolonged aging can lead to the coarsening of precipitates.When the aging temperature remains constant, during the initial stage of aging, a large number of precipitated particles will rapidly precipitate inside the alloy as the aging time prolongs.Therefore, the strength will gradually increase with the extension of aging time at this stage.When the alloy reaches peak aging state, the predominant form of precipitation phases is in the shape of η' phases, which are in a semi-coherent relationship with the matrix.As the aging time continues to extend, the precipitated particles will gradually coarsen and transform into a discontinuous distribution η phase.However, it takes quite a long time for the coarsening process when the aging temperature is low.Bai et al. [33] studied the effect of singlestage aging treatment on the mechanical properties of 7B04 aluminum alloy and found that the higher the aging temperature, the faster the strengthening response speed of the alloy and the shorter the time to reach peak aging.Additionally, when aged at 120 °C, the strength of 7B04 aluminum alloy gradually increases with time, reaching its peak after 24 h, and then the strength of the alloy remains basically unchanged.It is consistent with the results obtained in this article.During the artificial aging of Al-Zn-Mg-Cu alloy, precipitation follows the sequence of α Al -GP zone-η'-η.The final types, quantities, and sizes of precipitate phases are closely linked to the aging temperature and time.Generally, within a specific temperature range, higher aging temperatures lead to faster precipitation rates of solute atoms and increased quantities of precipitate phases.However, excessively high temperatures during prolonged aging can lead to the coarsening of precipitates.When the aging temperature remains constant, during the initial stage of aging, a large number of precipitated particles will rapidly precipitate inside the alloy as the aging time prolongs.Therefore, the strength will gradually increase with the extension of aging time at this stage.When the alloy reaches peak aging state, the predominant form of precipitation phases is in the shape of η' phases, which are in a semi-coherent relationship with the matrix.As the aging time continues to extend, the precipitated particles will gradually coarsen and transform into a discontinuous distribution η phase.However, it takes quite a long time for the coarsening process when the aging temperature is low.Bai et al. [33] studied the effect of single-stage aging treatment on the mechanical properties of 7B04 aluminum alloy and found that the higher the aging temperature, the faster the strengthening response speed of the alloy and the shorter the time to reach peak aging.Additionally, when aged at 120 • C, the strength of 7B04 aluminum alloy gradually increases with time, reaching its peak after 24 h, and then the strength of the alloy remains basically unchanged.It is consistent with the results obtained in this article.
Cao et al. [25] studied the effect of two-stage aging on the microstructure and properties of extruded 7075 aluminum alloy and found that with the increase in the second-stage aging temperature and the extension of aging time, the hardness of the alloy showed a trend of initial increasing, reaching a peak, and then gradually decreasing.They believe that during the second stage of aging, as the aging time prolongs, the critical size G.P. zone formed during the first stage of aging will serve as the nucleation core of the η' phase, and then the precipitation phase gradually transforms into the transition η' phase.Additionally, over time, the η' phase will continue to coarsen and form the equilibrium phase η that is coherent with the matrix, causing a decrease in the strengthening effect of the precipitate phase.It is similar to the results obtained in this work that with the increase in the second stage aging temperature and the extension of time, the strength and hardness of the thin plate gradually decrease.
After the two-stage aging treatment of Al-Zn-Mg-Cu alloy, the precipitates mainly exist in the form of large-sized semi-coherent transition phases η' or equilibrium phases η that are coherent with the matrix.At this time, the strengthening mechanism of the alloy is mainly the Orowan dislocation bypass mechanism.When the movement of dislocations encounters obstruction from precipitate particles, the dislocations circumvent these particles, leaving behind a dislocation loop around them [34].During more pronounced ductile deformation of the alloy, a multitude of interactions emerge between dislocations and precipitate particles.This leads to the formation of numerous dislocation loops encircling the precipitate particles, which further hinder the subsequent movement of dislocations, resulting in a noticeable enhancement of the alloy's strength.However, the extensive formation of dislocation loops can easily lead to dislocation pile-ups, causing a sharp reduction in the alloy's ductility.This phenomenon also accounts for fractures occurring during the bending deformation of the alloy subsequent to the artificial aging treatment.

Conclusions
This study primarily explores the impact of natural aging and artificial aging treatments on the microstructure and mechanical properties of Al-Zn-Mg-Cu alloy thin sheets.
The key findings are summarized as follows: 1.
Following the solution treatment of Al-Zn-Mg-Cu alloy, rapid natural aging occurs at room temperature, leading to a swift increase in strength and hardness.Moreover, even within 3 h after solution quenching, the alloy's ductility diminishes noticeably; however, it retains favorable formability, allowing for significant cold deformation.

2.
As the single-stage aging temperature rises, the alloy demonstrates a pattern of strength and hardness, initially increasing and then decreasing, with yield strength being more responsive to changes in aging temperature.Conversely, the elongation of the alloy exhibits an inverse trend compared to its strength with respect to aging temperature.

3.
During the dual-stage aging of Al-Zn-Mg-Cu alloy, an elevation in the second-stage aging temperature leads to a gradual reduction in the alloy's strength and hardness while elongation notably increases.With an extended duration of the second-stage aging, the alloy's strength and hardness gradually diminish, although elongation remains relatively unchanged.4.
Artificial aging treatment significantly boosts the strength and hardness of Al-Zn-Mg-Cu alloy thin sheets, yet it also triggers a sharp decline in ductility, limiting their ability to undergo extensive bending deformation and resulting in reduced cold formability.

Figure 1 .
Figure 1.Single-stage aging heat treatment scheme for Al-Zn-Mg-Cu alloy sheets.(a) Single-stage aging heat treatment scheme with different aging temperatures, (b) Single-stage aging heat treatment scheme with different aging time.

Figure 1 .
Figure 1.Single-stage aging heat treatment scheme for Al-Zn-Mg-Cu alloy sheets.(a) Single-stage aging heat treatment scheme with different aging temperatures, (b) Single-stage aging heat treatment scheme with different aging time.

Figure 3
Figure 3 illustrates the changing trends in tensile performance, hardness, and bending behavior of the thin sheet samples as a function of natural aging time.Within the first 0.5 h following solution treatment, the samples exhibit a tensile strength of 341 MPa, a yield strength of 153 MPa, an elongation of 20.90%, and a Vickers hardness of 84.52 ± 3.2.As the natural aging time extends, the strength and hardness of the samples gradually increase while the elongation rate diminishes rapidly.At a natural aging time of 3 h, the tensile strength rises to 367 MPa, the yield strength reaches 179 MPa, and the elongation drops to 18.31%.

Figure 3 .
Figure 3. Variation of mechanical properties of Al-Zn-Mg-Cu alloy thin sheets with natural aging time.(a) Tensile properties, (b) Hardness

Figure 3
Figure3illustrates the changing trends in tensile performance, hardness, and bending behavior of the thin sheet samples as a function of natural aging time.Within the first 0.5 h following solution treatment, the samples exhibit a tensile strength of 341 MPa, a yield strength of 153 MPa, an elongation of 20.90%, and a Vickers hardness of 84.52 ± 3.2.As the natural aging time extends, the strength and hardness of the samples gradually increase while the elongation rate diminishes rapidly.At a natural aging time of 3 h, the tensile strength rises to 367 MPa, the yield strength reaches 179 MPa, and the elongation drops to 18.31%.

Figure 3 .
Figure 3. Variation of mechanical properties of Al-Zn-Mg-Cu alloy thin sheets with natural aging time.(a) Tensile properties, (b) Hardness

Figure 3 .
Figure 3. Variation of mechanical properties of Al-Zn-Mg-Cu alloy thin sheets with natural aging time.(a) Tensile properties, (b) Hardness.

Figure 4 .
Figure 4. Macroscopic appearance (a) and microstructural characteristics of the bending deformation zone (b) in Al-Zn-Mg-Cu alloy thin sheet samples after bending.

Figure 5 .
Figure 5. EBSD analysis of T4N1 sample.Grain map (a) and distribution map (b) of grain boundary misorientation.

Figure 4 .
Figure 4. Macroscopic appearance (a) and microstructural characteristics of the bending deformation zone (b) in Al-Zn-Mg-Cu alloy thin sheet samples after bending.

Figure 4 .
Figure 4. Macroscopic appearance (a) and microstructural characteristics of the bending deformation zone (b) in Al-Zn-Mg-Cu alloy thin sheet samples after bending.

Figure 5 .
Figure 5. EBSD analysis of T4N1 sample.Grain map (a) and distribution map (b) of grain boundary misorientation.

Figure 5 .
Figure 5. EBSD analysis of T4N1 sample.Grain map (a) and distribution map (b) of grain boundary misorientation.

3. 3 .
Figure 6 illustrates the patterns of change in tensile performance and hardness during the single-stage aging of Al-Zn-Mg-Cu alloy thin sheets across different aging temperatures.As depicted, the variation in tensile strength remains relatively minimal within the temperature range of 100 • C to 130 • C.However, a noticeable decrease in tensile strength occurs after aging treatment at 140 • C (a reduction from 556 MPa at 120 • C to 537 MP).Conversely, the changes in yield strength and elongation are more pronounced with increasing aging temperature.With the rising temperature, the yield strength initially increases before showing a subsequent decline, while the elongation rate initially drops before increasing again.At an aging temperature of 120 • C, the yield strength (481 MPa) reaches the peak value, coinciding with the lowest elongation rate (12.55%).From Figure6b, it can be seen

Figure 6 .
Figure 6.Tensile performance (a) and hardness (b) of Al-Zn-Mg-Cu alloy thin sheets under different single-stage aging treatments at various temperatures.

Figure 7
Figure 7 depicts the curves showing the changes in tensile performance and hardness of Al-Zn-Mg-Cu alloy thin sheets during single-stage aging treatments as a function of aging time.From Figure7a, it can be seen that with the extension of aging time, both the yield strength and tensile strength of the alloy remain basically unchanged, while the elongation shows a trend of gradually decreasing and then increasing.From Figure7b, it can be seen that the hardness of the thin plate shows a trend of first increasing and then gradually decreasing with the extension of aging time, but the change amplitude is relatively small.Based on these experimental results, it can be concluded that, compared to aging time, aging temperature has a more significant influence on the alloy's overall properties.

Figure 7 .
Figure 7. Tensile performance (a) and hardness (b) of Al-Zn-Mg-Cu alloy thin sheets under different single-stage aging treatments at various time intervals.

Figure 6 .
Figure 6.Tensile performance (a) and hardness (b) of Al-Zn-Mg-Cu alloy thin sheets under different single-stage aging treatments at various temperatures.

Figure 7
Figure 7 depicts the curves showing the changes in tensile performance and hardness of Al-Zn-Mg-Cu alloy thin sheets during single-stage aging treatments as a function of aging time.From Figure7a, it can be seen that with the extension of aging time, both the yield strength and tensile strength of the alloy remain basically unchanged, while the elongation shows a trend of gradually decreasing and then increasing.From Figure7b, it can be seen that the hardness of the thin plate shows a trend of first increasing and then gradually decreasing with the extension of aging time, but the change amplitude is relatively small.Based on these experimental results, it can be concluded that, compared to aging time, aging temperature has a more significant influence on the alloy's overall properties.

Metals 2023 ,
13,  x FOR PEER REVIEW 6 of 13 197.6HV.Additionally, when the aging temperature is 120 °C, the hardness of the thin plate is 197.56 ± 3.65 HV.

Figure 6 .
Figure 6.Tensile performance (a) and hardness (b) of Al-Zn-Mg-Cu alloy thin sheets under different single-stage aging treatments at various temperatures.

Figure 7
Figure 7 depicts the curves showing the changes in tensile performance and hardness of Al-Zn-Mg-Cu alloy thin sheets during single-stage aging treatments as a function of aging time.From Figure7a, it can be seen that with the extension of aging time, both the yield strength and tensile strength of the alloy remain basically unchanged, while the elongation shows a trend of gradually decreasing and then increasing.From Figure7b, it can be seen that the hardness of the thin plate shows a trend of first increasing and then gradually decreasing with the extension of aging time, but the change amplitude is relatively small.Based on these experimental results, it can be concluded that, compared to aging time, aging temperature has a more significant influence on the alloy's overall properties.

Figure 7 .
Figure 7. Tensile performance (a) and hardness (b) of Al-Zn-Mg-Cu alloy thin sheets under different single-stage aging treatments at various time intervals.

Figure 7 .
Figure 7. Tensile performance (a) and hardness (b) of Al-Zn-Mg-Cu alloy thin sheets under different single-stage aging treatments at various time intervals.

Metals 2023 ,
13,  x FOR PEER REVIEW 7 of 13 small, as shown in Figure8d.This observation underscores the subpar bending performance of the sample treated with the T6 treatment.

Figure 8 .
Figure 8. Macroscopic appearance (a) and microstructure of deformation zone (b) in T6N2 sample after bending and high magnification micrographs of the outer (c) and inner (d) selected areas of the deformation zone.

Figure 8 .
Figure 8. Macroscopic appearance (a) and microstructure of deformation zone (b) in T6N2 sample after bending and high magnification micrographs of the outer (c) and inner (d) selected areas of the deformation zone.

3. 4 .
Figure9illustrates the curves depicting the changes in tensile performance and hardness during the dual-stage aging treatment of Al-Zn-Mg-Cu alloy thin sheets concerning the temperature and duration of the second-stage aging.As observed in the figure, as the temperature of the second-stage aging rises, there is a noticeable reduction in the strength and hardness of the samples, accompanied by a significant increase in the elongation rate.Furthermore, extending the duration of the second-stage aging results in a gradual decline in both strength and hardness while the elongation rate remains relatively stable.This phenomenon can be attributed to the different morphologies of precipitates formed during the distinct stages of the two-stage aging treatment for Al-Zn-Mg-Cu alloy.During the initial phase of low-temperature aging in the first stage, the alloy's grain interiors experience the development of numerous fine and dispersed GP zones or coherent precipitate phases.Subsequently, in the second stage, characterized by higher temperature and prolonged aging, the precipitate particles formed during the first stage transform into semi-coherent η' precipitates or equilibrium η precipitates[25,26].Consequently, in the second-stage aging process, the size and quantity of precipitates within the alloy undergo changes influenced by aging temperature and duration, leading to variations in the alloy's strength and hardness with changing aging conditions.Moreover, a comparative analysis highlights that aging temperature exerts a more pronounced influence on the alloy's mechanical properties.

Figure 9 .
Figure 9. Variation curves of mechanical properties with aging time during the second-stage aging treatment at different temperatures for Al-Zn-Mg-Cu Alloy thin sheets.

Figure 9 .
Figure 9. Variation curves of mechanical properties with aging time during the second-stage aging treatment at different temperatures for Al-Zn-Mg-Cu Alloy thin sheets.

Figure 10
Figure10illustrates the macroscopic appearance and microstructure at the fracture surface of the sample (T74N3) after undergoing the process of solution treatment at 470 • C for 38 min, followed by two-stage aging treatments at 115 • C for 8 h and 165 • C for 16 h.In Figure10a, it is evident that prominent cracks have developed in the region of bending deformation within the sample.These cracks align precisely with the ridge line of the bending angle, displaying a relatively straight configuration that spans the entire width of the sample.Additionally, as depicted in Figure10b, the microstructure in the bending deformation zone showcases marked grain deformation.The grains on the outer side of the bending deformation zone (left side in Figure10b) are noticeably elongated, resembling a fibrous pattern, while those on the inner side are compressed, adopting a polygonal shape.Noteworthy is the comparison with the morphology of the T6N2 sample after bending, as shown in Figure8.This reveals that the T74N3 sample displays a more pronounced level of grain deformation.This observation suggests that, in comparison to the forming performance of the thin sheet after the T6 treatment, the forming performance of the thin sheet following the T74 treatment has been moderately enhanced, although it still falls short of achieving a significant degree of bending deformation.

Figure 10 .
Figure 10.Macroscopic appearance (a) and microstructure of deformation zone (b) in T74N3 sample after bending deformation.

Figure 10 .
Figure 10.Macroscopic appearance (a) and microstructure of deformation zone (b) in T74N3 sample after bending deformation.

Table 1 .
Double aging heat treatment scheme for Al-Zn-Mg-Cu alloy sheets.

Table 1 .
Double aging heat treatment scheme for Al-Zn-Mg-Cu alloy sheets.

Table 2 .
Mechanical properties of annealed samples.

Table 2 . Mechanical properties of annealed samples. Properties Yield Strength/MPa Tensile Strength/MPa Elongation/% Hardness/HV
3.2.Effects of Natural Aging on the Performance of Al-Zn-Mg-Cu Alloy Thin Sheets

Table 2 .
Mechanical properties of annealed samples.