The Influence Mechanism of Deformation on the Precipitation Behavior and Mechanical Properties of 7075 Aluminum Alloy During Hot Forming-Quenching Integrated Process

Abstract

The hot forming-quenching integrated process (HFQ^®) organically combines the deformation and heat treatment, which can improve the forming performance of aluminum alloy while ensuring the final strength of formed parts. Thermal deformation in HFQ^® has a non-negligible influence on precipitation behavior in subsequent artificial aging treatment and affects the mechanical properties of the formed parts. In this study, the relationship between the precipitation behavior and thermal deformation ratios was investigated. Results indicated that the formation temperatures of η′ and η decreased with an increasing deformation ratio; however, the former decreased more than the latter. The activation energy of η′ precipitation decreased linearly with increasing deformation ratio. Additionally, the phase transition fraction of η′ increased with the deformation ratio, leading to shorter times required to reach equivalent phase transition fractions. Deformation accelerated the phase transition of η′, and greater deformation resulted in a shorter transition time. The aging time required for peak Vickers hardness decreased with increasing deformation ratio, reflecting the promotion of precipitated phase formation and transformation by introduced dislocations. Consequently, peak hardness and yield strength were achieved in shorter aging times. In terms of industrial applications, this discovery offers significant advantages for shortening the production cycle of the hot stamping process and reducing production costs.

1. Introduction

High-strength aluminum alloys are widely used in lightweight fields such as aerospace and transportation owing to their high specific strength attributes [1,2]. However, their poor forming performance often leads to failures such as fracture and springback for complex-designed structural parts. To address this limitation, a hot forming-quenching integrated process (HFQ^®) has been proposed. In this process, the solid solution sheet metal is quickly transferred to a cold die for hot stamping, which not only improves the formability but also avoids thermal distortion, and the following artificial aging can confirm the final strength of the formed part. However, the long process cycle results in multiple factors affecting the final performance of the formed parts, making performance control necessary.

Significant progress has been made in the research of HFQ^® in recent years [3,4,5]. Ford Motor Co. (Detroit, MI, USA) [6] successfully trial-produced AA7075 B-pillar using HFQ^®. It was found that after 475 °C solid solutioned + hot stamped and held for 15 s +120 °C/24 h artificially aged, the microstructure and mechanical properties of the formed components met the requirements and were similar to those of those subjected to conventional artificial aging heat treatment, verifying the feasibility of HFQ^®. Fan et al. [7] investigated the strengthening behavior of Al-Cu-Mg alloy sheets in HFQ^®. In their study, the hot-stamped sheets were naturally aged to obtain enough strength. Ma et al. [8] analyzed the influence of solution temperature and time on mechanical properties for AA6082 by using a central composite design. Liu et al. [9] researched the formability and lubrication of AA6061 and AA7075 in HFQ^®. The formed AA7075 B-pillar was then subjected to 120 °C/24 h artificial aging. Xiao et al. [10] studied the heat transfer coefficient under gap conditions for AA7075 during the stamping process in HFQ^®. Wang et al. [11] established a modified Zerilli-Armstrong constitutive model of plastic deformation behavior for AA7075-H18, and then this model was employed for numerical simulation, and the simulation results were in good agreement with the experimental data, indicating the accuracy of the model. However, almost all of the studies focused on the formality and mechanism of the stamping process, and few studies on the precipitation behavior of thermal deformation on subsequent artificial aging. Lu et al. [12] identified the effect of thermal deformation on the microstructure and mechanical properties of Al-Zn-Mg-Cu alloys. Results demonstrated that the peak aging time was shortened with the increase of the thermal deformation ratio. It is evident that the hot stamping deformation during HFQ^® inevitably impacts the precipitation behavior and final mechanical properties. Therefore, it is of great significance to reveal the influence of deformation on the precipitation behavior during artificial aging and on the final mechanical properties.

AA7075 belongs to Al-Zn-Mg-(Cu) heat-treatable aluminum alloy, whose precipitation behavior is closely related to the deformation between solid solution and artificial aging heat treatment [13]. Han et al. [14] discovered that pre-deformation can improve the aging precipitation rate and density of Al-Zn-Mg-Cu alloys, thereby enhancing the strength and elongation of the alloys. Wang et al. [15] found that the judicious combination of appropriate pre-deformation and aging treatment can effectively enhance the mechanical properties and stress corrosion resistance of Al-Zn-Mg-Cu alloys. YU et al. [16] proved that the introduction of 20% pre-deformation increased the hardness of 7050 aluminum alloy by 30% at a slower quenching rate than that of the non-pre-deformed specimen, but had no effect on the hardness of 7050 aluminum alloy with relatively fast quenching rate, such as water quenching. Zou et al. [17] demonstrated that the effects of pre-deformation on the mechanical properties and microstructure of Al-Zn-Mg-Cu alloys are closely relevant to the alloy composition and Zn/Mg ratio. Specifically, the introduction of pre-deformation reduced the hardness of Al-Zn-Mg-Cu alloys with low Zn/Mg ratios while significantly increasing the hardness of Al-Zn-Mg-Cu alloys with high Zn/Mg ratios. Besides, some studies have also concluded that pre-deformation cannot ameliorate the mechanical properties of 7xxx alloys. For instance, the introduction of 10% tensile pre-deformation reduced the yield strength of AA7108 and AA7030 alloys in the T6 state by 7–10% [18]. Although the aging response of 7A04 aluminum alloy was expedited by extensive cold deformation after solution treatment, the increase in tensile strength was not noteworthy, and the elongation even decreased sharply [19].

In summary, a large number of literature studies on the precipitation behavior and final mechanical properties of aluminum alloy mainly focus on cold deformation. However, there are few studies on hot deformation, and the influence of hot deformation on the precipitation behavior and mechanics of aluminum alloy HFQ^® formed parts has not been fully revealed. Therefore, in this paper, isothermal uniaxial tensile tests, DSC, mechanical properties testing, and microstructure observation are used to study the influence of thermal deformation on the formed parts of the AA7075 alloy sheet in HFQ^®. The non-isothermal and isothermal transfraction of η′ phase during aging are analyzed. The relationship between different thermal deformation and the peak Vickers hardness and aging time of AA7075 is deeply investigated.

2. Experimental Details

2.1. Material

The material utilized in this study is a heat-treatable 7075 aluminum alloy sheet provided by Southwest Aluminum Company (Kunshan, China). It was a rolled sheet with a thickness of 1.5 mm and a state of T6. The tensile and yield strengths of the material were 571.2 MPa and 489.6 MPa, respectively. The chemical compositions are listed in

Table 1. The chemical compositions of the supplied AA7075-T6 sheet.

Element	Si	Fe	Cu	Mn	Mg	Cr	Zn	Ti	Al
Wt%	0.073	0.34	1.50	0.075	2.62	0.21	5.39	<0.1	Balance

. The metallographic image is illustrated in Figure 1. Furthermore, the average grain size of the supply was about 15.6 μm.

2.2. Isothermal Uniaxial Tensile Tests

Isothermal uniaxial tensile tests were performed on the Gleeble-3500 thermal simulation machine (Dynamic Systems Inc., Poestenkill, NY, USA). The deformation temperatures and strain rates in this process were automatically controlled by a computer, and the temperature accuracy was controlled within ±0.5 °C. The profile and size of the test specimens in isothermal uniaxial tensile tests are shown in Figure 2. The specific experiment routine is shown in Figure 3 and is described as follows: the thermocouple was welded in the middle of the tensile sample to measure and control the temperature. Notably, the two thermocouple wires should not be too close or in direct contact to avoid a short circuit. The tensile sample was first heated to 400 °C at a heating rate of 10 °C/s and then heated to 475 °C at a heating rate of 5 °C/s for 30 min to ensure a complete solid solution for the samples. Samples were cooled to 400 °C at a cooling rate of 2 °C/s (to simulate the temperature drop during the transfer in actual production), and tensile deformation was carried out at different deformation ratios (0, 10%, 15%, 20%, and 25%) and different deformation rates (0.001 s⁻¹, 0.01 s⁻¹, 0.1 s⁻¹ and 1 s⁻¹) according to the experimental scheme shown in

Table 2. Experimental scheme for isothermal uniaxial tensile test of AA7075-T6.

Solid Solution	Deformation Temperature	STRAIN Rate/s−1	Deformation Ratio	Aging Temperature	Aging Time/h
475 °C 30 min	400 °C	0.1	0, 10%, 15%, 20%, 25%	120 °C	0~28
		0.001, 0.01, 0.1, 1	10%

. As soon as the deformation was completed, cooling water was sprayed onto the samples quickly (within 3 s). Subsequently, the samples were stored in liquid nitrogen and subjected to artificial aging at 120 °C with different aging times within 2 h.

2.3. Microstructural and Properties Characterization

After the completion of each experimental group, according to Table 2, specimens of each group were sampled for individual subsequent processing. As illustrated in Figure 4a, two dog-bone tensile pieces (marked by red solid lines) were wire-cut from the center of the specimens, and three blocks were obtained from the residual material between these two tensile pieces, labeled as No. 1 to No. 3, respectively. No. 1 and No. 2 were marked with yellow solid rectangles, while No. 3 was marked with a yellow solid circle. Furthermore, the specific dimensions of the tensile pieces and the three blocks are depicted in Figure 4b–d. Next, the tensile pieces and No. 1 block were subjected to artificial aging treatment at a temperature of 120 °C at different times simultaneously. The tensile pieces were used to identify the yield strength of the specimens under different deformation ratios and aging times, and the No. 1 block was used to identify the Vickers hardness of the specimens under different deformation ratios and aging times. Upon completion of the aging treatment, room temperature tensile tests were conducted on an electronic universal testing machine to measure the yield strength variations under different deformation ratios and aging times using the tensile pieces from each experimental group. The room temperature tensile test measurement referred to “GBT228.1-2010 metal material tensile test part 1: room temperature test method” as a standard [20]. The yield strength was taken as the average value of two tensile pieces. Meanwhile, the Vickers hardness test was employed to measure the hardness variations using the No. 1 block on HVS-100 with a load of 200 g and a time of 10 s. The measurement load was set to 200 g, with a holding time of 10 s. The No. 2 block was used for TEM (Transmission electron microscope) observation under different conditions. TEM samples were mechanically thinned to 80–100 μm. A disc with a diameter of 3 mm was punched out from the thinned sample and then subjected to twin-jet electropolishing in a mixture of 30% HNO₃ and 70% CH₃OH at a temperature of −25 °C and a voltage of 20 V. The TEM observations were performed on the Tecnai G2 F30 Transmission Electron Microscope (FEI Corporation, Columbia City, MD, USA). DSC analysis was conducted on a TA-DSC 2500 differential scanning calorimeter (TA Instruments Corporation, Newcastle, DE, USA) using the No. 3 disc with a diameter of 5 mm and a weight of 25 ± 5 mg.

3. Theoretical Model

It is widely recognized that the precipitation behavior of heat-treatable aluminum alloys is intricately linked to the interplay between deformation and aging processes [21]. While the non-isothermal method offers continuous calculation of kinetic parameters throughout the entire temperature range from the onset to the culmination of precipitation, the isothermal approach necessitates the sample to reach a specific temperature and exhibit discernible reflections before measurement. However, it is challenging to precisely control the initiation and termination states of these reflections in the isothermal process, and one non-isothermal thermal analysis curve is equivalent to numerous isothermal thermal analysis curves, making experimental supplies limited. Thus, the DSC combined with the Kissinger and Isoconversional method was employed to investigate the relationship between precipitation behavior and deformation ratios, unveiling the precipitation kinetics during the aging process in the HFQ (Hot Form Quench) process.

In both isothermal and non-isothermal heating processes, the precipitation state at the target aging temperature is determined by the heating rate [22]. According to the assumptions in the Johnson–Mehl–Avrami–Kolmogorov (JMAK), the phase transition mechanism remains constant in consideration of the time–temperature (t-T) range, i.e., the precipitate growth rate β is just a function of temperature, then in a non-isothermal process $\beta$ can be described as [23]:

$$\beta ={\displaystyle \int K\left(T\right)}dt$$

(1)

in which $K\left(T\right)$ is the heating rate (K·min⁻¹); if the heating rate remains constant, set $\varphi =dT/dt$, the following equation can be obtained:

$$\beta ={\displaystyle \frac{T^{2}R}{\varphi E}}K\left(T\right)$$

(2)

where E is the apparent activation energy (J·min⁻¹) and R is the molar gas constant (8.31 J·mol⁻¹·K⁻¹). The heating rate $K\left(T\right)$ can be expressed by the Arrhenius equation as follows:

$$K\left(T\right)=K_0\exp \left(-{\displaystyle \frac{E}{RT}}\right)$$

(3)

In this paper, only the maximum transition rate at the peak temperature (T_p) was taken into account. Combining Equations (2) and (3), logarithms of both sides of the equation can be obtained as follows:

$$\ln \left({\displaystyle \frac{\beta }{T_p^2}}\right)=\ln {\displaystyle \frac{AR}{E}}-{\displaystyle \frac{E}{R{T}_p}}$$

(4)

where A is the pre-exponential factor (s⁻¹), and T_p is the temperature corresponding to the endothermic or exothermic peak on the DSC curve (°C).

The values of $\ln {\displaystyle \frac{T_p^2}{\varphi }}$ are used as the x-axis, and the values of ${\displaystyle \frac{1}{T_p}}$ are used as the y-axis for linear fittings. The activation energy E of the precipitation can be obtained by the slopes of the fitting lines, and then the pre-exponential factor A can be calculated based on the intercepts of the fitting lines.

According to Equations (1) and (3), the time required for isothermal aging (corresponding to the aging temperature T₀ = 120 °C in this subject) with the same conversion rate as continuous heating is as follows:

$${\displaystyle \frac{\beta }{K\left(T_0=120 °\mathrm{C}\right)}}={\displaystyle {\int }_0^{t}dt}=t$$

(5)

Therefore, under the condition that the aging temperature is constant, and assuming that the time required for the precipitation to reach a certain transition α is t_α, then the current transition $Y_{\alpha }$ can be expressed as follows:

$$Y_{\alpha }=1-\exp \left(-k{t}_{\alpha }^n\right)$$

(6)

where $Y_{\alpha }$ is the current transition, and k and n are the JMAK kinetic parameters.

4. Results and Discussion

4.1. Curves of Differential Scanning Calorimetry

Figure 5 illustrates the DSC heat flow curves at different heating rates of solid solution 7075 aluminum alloy subjected to various deformation ratios. As can be seen from the figure, the DSC curves show a similar trend. Specifically, there are three distinct endothermic peaks labeled 1, 2, and 3 in all curves. As the temperature increases, the curves initially ascend to peak 1, followed by a gradual decline, then a rapid rise to peak 2, subsequently dropping sharply, and finally rapidly rising to peak 3. The response peaks of these DSC curves are consistent with those in the published literature [24]. Based on the aging precipitation sequence of 7xxx aluminum alloys, the endothermic peaks in Figure 5 correspond to the precipitation of the GP, η′, and η, respectively.

Moreover, the greater the deformation ratio (0~25%), it is evident that the lower the temperature at which endothermic peaks 1, 2, and 3 occur. For instance, when the heating rate is 10 °C/min, for specimens with a deformation ratio of 0, the formation temperatures of the η′ and η phases are 210 °C and 245 °C, respectively, as shown in Figure 5a. However, for specimens with a deformation ratio of 10%, the formation temperatures of the η′ and η are 206 °C and 244 °C, respectively, as shown in Figure 5b. Similarly, for specimens with a deformation ratio of 25%, the formation temperatures of the η′ and η are 197 °C and 239 °C, respectively, as shown in Figure 5e. Notably, the formation temperatures of η′ for 25% deformation are reduced by 9.5% compared to that of the undeformed specimen, while the formation temperatures of η phase for 25% deformation are reduced by 2.4% compared to that of the undeformed specimen. In summary, the formation temperatures of η′ and η gradually decrease with the increase in deformation ratio (0 to 25%), with the formation temperature of η′ decreasing more significantly than that of η.

It is widely acknowledged that the metastable η′ plays a key role in the strengthening of 7xxx aluminum alloys [25,26], making it the primary focus of this study. Figure 6 shows the DSC curves of the η′ of the specimens with different deformation ratios at different heating rates, which are derived from the curves in Figure 5. It can be observed that all the curves exhibit the same distribution trend along with the x-axis direction. Specifically, they exhibit a sharp rise to the peak with increasing temperature, followed by a rapid decline. Under an identical deformation ratio, the temperatures corresponding to the peaks gradually increase with rising heating rates. For instance, considering Figure 6c, where the deformation ratio is 15%, the temperatures of the peaks for the η′ at different heating rates (5 °C/min, 10 °C/min, 15 °C/min, and 20 °C/min) are 193 °C, 204 °C, 211 °C, and 216 °C, respectively. Notably, the temperature progressively rises with an increasing heating rate, accompanied by a corresponding increase in the peak area. This phenomenon primarily stems from the heightened phase transition intensity per unit of time due to the escalated heating rate, leading to a gradual increase in the generation rate of the η′ precipitated phase. Moreover, when compared to the deformation-free condition (Figure 6a), the integrated peak areas under deformation conditions in Figure 6b–d are notably larger, as illustrated in Figure 6f. This enhancement can be attributed to the introduction of deformation promoting the increase of η′ transition rate [27].

4.2. Evolution of the Non-Isothermal Transformation Fraction

Figure 7 shows the phase transformation fraction change curves of η′ during continuous heating of solid solution 7075 aluminum alloy under different deformation ratios. It is evident that the curves obtained exhibit an S-shaped profile for all deformation conditions, indicating that the transformed fraction increases as the temperature increases at a fixed heating rate. However, at a specific temperature, the transformed fraction decreases with an increasing heating rate. Taking Figure 7b as an example, when the deformation is 10% and the temperature is 200 °C, the transformed fractions of η′ at heating rates of 5 °C/min, 10 °C/min, 15 °C/min, and 20 °C/min are 57.5%, 38.7%, 25.7%, and 25.6%, respectively. This is mainly due to the influence of temperature and time on the diffusion of solute atoms. An increase in temperature will accelerate diffusion, thereby increasing the transformed fraction. On the contrary, an increase in heating rate leads to insufficient time and insufficient diffusion of solute atoms, thereby reducing the transformed fraction [28].

Based on the precipitation kinetic theory in Section 3, the activation energy of η′ under different deformation ratios was shown in Figure 8, which was obtained by fitting the corresponding temperature values of η′ when the transformed fraction is 50%. From Figure 8b, it can be seen that the activation energy of η′ decreases almost linearly with the increase of the deformation ratio. For the deformation-free condition, the activation energy of the η′ is 120.8 kJ·mol⁻¹, while the activation energies of the η′ at deformations of 10%, 15%, 20%, and 25% are 115.3 kJ·mol⁻¹, 111.1 kJ·mol⁻¹, 107.5 kJ·mol⁻¹, and 102.7 kJ·mol⁻¹, respectively, as shown in

Table 3. The activation energies of the η′ at different deformations.

Deformation Ratio	Activation Energy (kJ·mol−1)
0	120.8
10%	115.3
15%	111.1
20%	107.5
25%	102.7

.

4.3. Evolution of the Isothermal Transformation Fraction

Based on the DSC curves and the theory outlined in Section 3, the precipitation behavior of non-isothermal aging under continuous heating with different deformation ratios can be transformed into that of under-isothermal aging. The theoretical change of η′ transformed fraction with aging time can be obtained at an isothermal aging temperature of 120 °C, as shown in Figure 9. The curves obtained during isothermal exhibit S-shaped, indicating that at a fixed aging temperature, the transformed fractions increase with aging time. Additionally, for a given aging time, the transformed fractions increase with increasing deformation ratios. For instance, at an aging time of 5 h, the transformed fractions of the η′ for deformation ratios of 0–25% are 35.1%, 57.8%, 73.8%, 86%, and 93%, respectively. Furthermore, it is observed that the time required to reach the same transformed fraction decreases with an increasing deformation ratio. Specifically, the time required for the η′ to reach a 50% fraction decreases from 6.8 h to 1.9 h as the deformation ratio increases from 0% to 25%. This reduction indicates that the transition of η′ is accelerated with increasing deformation ratio, signifying that deformation promotes the precipitation, with the larger deformation ratios, the earlier the transition.

The k and n values in the isothermal kinetic equation of the η′ under different deformation ratios can be deduced according to Figure 9 and Equation (6), as shown in Figure 10. As shown in the figure, during the isothermal aging process, the k value of η′ increases with the increase of the deformation ratio, while the n value decreases with the increase of the deformation ratio, further indicating that the introduction of deformation promotes the precipitation.

4.4. Aging Hardening Behavior

The hardness measurement of solid solution hot tensile specimens with different deformation ratios and aging times was used to verify the precipitation kinetic behavior mentioned above. Figure 11a shows the variation curves of Vickers hardness with the deformation ratio and aging time under the condition of a deformation temperature of 400 °C and a strain rate of 0.1 s⁻¹. It can be seen that both the deformation ratio and aging time have a significant influence on the Vickers hardness. Irrespective of the deformation presence, the hardness of each tensile specimen experiences a rapid increase followed by a gradual reduction as aging time progresses from 0 to 28 h. This is because, in the early stage of aging, the matrix has a high degree of supersaturation, which leads to a strong force for the precipitation and a correspondingly fast precipitation rate, resulting in a rapid increase in hardness. The precipitation in the early stage of aging maintains a coherent relationship with the matrix, resulting in a large number of lattice distortion areas within the matrix. With the extension of aging time, the precipitation grows, and the coherent relationship is disrupted, resulting in a decrease in the degree of lattice distortion and, thus, the decrease in hardness [29]. Moreover, it is evident that the aging time required for Vickers hardness to reach the peak is less in the presence of deformation compared to without deformation. At 0% deformation ratio, the peak Vickers hardness is attained at 17 h of aging time, measuring 177.7 HV. At a 10% deformation ratio, the peak Vickers hardness is reached at 11 h of aging time, measuring 166.2 HV. At a 15% deformation ratio, the peak Vickers hardness reached 8 h of aging time, measuring 166.2 HV. At 20% deformation ratio, the peak Vickers hardness is reached at 7 h of aging time, measuring 178.2 HV; and at 25% deformation, the peak Vickers hardness is reached at 7 h of aging time, measuring 181.2 HV. When the deformation ratio is 10%, 15%, 20%, and 25%, the time required for Vickers hardness peak is 6 h, 9 h, 10 h, and 10 h earlier than that when the deformation ratio is 0, respectively. Remarkably, the time required for peak Vickers hardness at a 25% deformation ratio is 58% shorter than that at a 0% deformation ratio. The variations of yield strength with deformation ratios and aging time mirrors that of Vickers hardness, wherein yield strength under different deformation ratios rapidly increases and then gradually decreases with aging time, as illustrated in Figure 11b.

AA7075 belongs to the heat-treatable Al-Zn-Mg-Cu alloy, whose yield strength is mainly contributed by precipitation strengthening, dislocation strengthening, solid solution strengthening, and grain boundary strengthening [30,31]. The contributions of each mechanism on yield strength can be expressed as follows [32,33]:

$${\sigma }_y=\Delta {\sigma }_{gb}+\Delta {\tau }_0+\Delta {\tau }_{ss}+{\left(\Delta {\tau }_D^2+\Delta {\tau }_{ppt}^2\right)}^{1/2}$$

(7)

in which $\Delta {\sigma }_{gb}$ denotes the strength increment caused by grain boundaries, $\Delta {\tau }_0$ represents the intrinsic strength of the aluminum alloy matrix, $\Delta {\tau }_{\mathrm{ss}}$ signifies the strength introduced by solid solution, $\Delta {\tau }_D$ accounts for the strength introduced by dislocations, and $\Delta {\tau }_{ppt}$ reflects the strength contributed by precipitation. As mentioned in prior investigations [34,35], grain boundary strengthening plays a negligible role in the strength during HFQ, while precipitation strengthening and dislocation strengthening are the key factors affecting mechanical properties.

Figure 12 presents TEM observations of samples formed at a deformation of 10% with an aging time of 11 h and at a deformation of 20% with an aging time of 7 h (corresponding to the peak aging times for different deformations in Figure 11). The morphologies and dimensions of precipitation have been extensively studied for commercially available heat-treatable 7075 aluminum alloys [36,37,38,39]. Typically, the Guinier–Preston (GP) zones are too diminutive to be discerned. η′ is disk-shaped within 50 nm, while η appears as laths and slats, nearing 200 nm and averaging around 50 nm. Figure 12(a1,a2) shows the TEM observations inside the grain; there are many disc-shaped precipitates about 10–20 nm in the grain, which correspond to the η′ [36]. Additionally, larger lath-like precipitates around 50 nm are also evident, which represent the η [36]. Notably, the average size of the η at a 10% deformation ratio is around 30 nm, whereas, at a 20% deformation ratio, it measures approximately 50 nm, which is 67% larger than that at a 10% deformation ratio. This discrepancy is attributed to the heightened energy available with an increased deformation ratio, facilitating larger precipitate particle sizes.

Figure 12(b1,b2) shows the TEM observations in the grain boundaries. At a lower deformation ratio (as depicted in Figure 12(b1)), the dislocation density is relatively lower, and the precipitate-free zone (PFZ) is not obvious. The precipitates at the grain boundary are distributed discontinuously along the grain boundary, and the size of the η distributed along the grain boundary is slightly larger than that inside the grain. As the deformation ratio increases to 20% (as depicted in Figure 12(b2)), the morphology becomes characterized by wider grain boundaries and clearer PFZ. In addition, the coarsening of intragranular precipitates and the dislocation pile-up near the grain boundaries can be observed. During HFQ, an abundance of dislocations occurs, rendering the alloy thermodynamically unstable and significantly enhancing the driving force for phase transformation. Consequently, the homogeneous nucleation changes to heterogeneous nucleation due to the deformation providing nucleation particles, reducing nucleation energy barriers, thereby promoting the formation of small and dispersed precipitates within the grain. However, excessive dislocation may lead to rapid precipitation and subsequent coarsening. Hence, the introduction of deformation is beneficial for achieving peak hardness and yield strength in a shorter aging time.

5. Conclusions

This study employed uniaxial hot tensile tests, differential scanning calorimetry (DSC), mechanical property tests, and transmission electron microscopy (TEM) observations to investigate the impact of deformation during hot stamping on subsequent precipitation behavior during aging. The research aimed to elucidate the mechanism of different deformations that influence precipitation behavior during aging and offered theoretical insights for designing aging processes in hot stamping systems. The principal findings were summarized as follows:

The influence of thermal deformation on precipitation behavior in HFQ^® was revealed. As the deformation ratio increases (0~25%), the formation temperatures of η′ and η phases gradually decreased; however, the former decreased more than the latter. Simultaneously, deformation during hot stamping accelerated the transformation rate of the η′ phase to the η phase.

In the deformation range of 0~25%, the activation energy of η′ precipitation phase decreased linearly with increasing deformation. The k value in the kinetic equation of η′ precipitation phase increased with deformation, while the n value decreased, indicating that the introduction of deformation promoted the precipitation of the η′ phase.

The volume fraction of η′ increased with aging time during isothermal aging. With increasing deformation, the phase transition fraction of η′ increased at the same aging time, and the time required to reach the same phase transition fraction decreased. Compared with samples without deformation, the transition time of η′ phase was advanced with deformation. Moreover, the larger the deformation, the more accelerated the transition time, indicating that deformation accelerated the precipitation of the phase.

It was found that the peak Vickers hardness and aging time of AA7075 were significantly affected by thermal deformation during HFQ^®. Under deformation conditions, the time required for the Vickers hardness of the formed parts to reach its peak value was notably shorter compared to those without deformation. Moreover, the required aging time decreased with increasing deformation. Peak strength could be achieved within 8 to 10 h of artificial aging in the HFQ^®, representing a reduction of 58% to 67% compared to conventional artificial aging time. In terms of industrial applications, this discovery offered significant advantages for shortening the production cycle of the hot stamping process and reducing production costs.