4.1. Mechanism of Pre-Deformation Effect on Precipitation Behavior
In order to investigate the precipitation behavior of the alloy under different deformation heat treatments, DSC experiments were carried out, as shown in
Figure 10. It can be observed that there are two typical heat-absorbing peaks—A and C—on the curve, indicating the dissolution of some phases. Meanwhile, the other inverse peaks accompanying the heat absorption peaks are exothermic peaks—labeled as B and D—indicating the precipitation of a phase. According to the previous study [
8], the formation of the inspiral peak A corresponds to the dissolution of the GP (Guinier–Preston) region. The exothermic peaks of B and D represent the precipitation of the η′ and η phases, respectively, and C represents the back dissolution of the η′-enhanced phase. With the increase in temperature in
Figure 10a, at peak B1, the reinforced phase η′ of Al-Zn-Mg-Cu alloy precipitates along the GP zone at about 56.7 °C, which is the main precipitated phase in the peak aging stage. After a longer aging treatment, the η phase precipitates after η′ starts to dissolve back at peak C1. It can be seen that the intensity of the dissolution peak in the GP zone increases in the pre-stretched samples with deformations of 2%, 5%, and 7% compared to the un-pre-stretched samples, and the intensity of the precipitation peak of the η′ phase is higher than that of the un-pre-stretched samples, which implies that the volume fractions of the GP zone and the η′ phase in the peak aging samples of the pre-stretched samples are more than those of the un-pre-stretched samples. It is worth noting that the temperature of the dissolution peak (146.4 °C) corresponding to the dissolution of the η′ phase of the 7% sample in
Figure 10d is significantly shifted toward higher temperatures, suggesting that the η′ phase of the artificially aged tissues has a larger size at 7% pre-stretching, and higher temperatures are required for dissolution. In addition, the areas of the precipitation and coarsening peaks of the η phase in the 7% sample are significantly larger than those in the 2% and 5% samples. This is due to the presence of a large number of dislocations as the non-uniform nucleation points of the η phase, which is rapidly precipitated and undergoes significant coarsening. Therefore, and thus the areas of the exothermic and heat-absorbing peaks corresponding to the precipitation of the η phase are larger in the sample with a pre-stretching of 7%. It is also observed that the pre-stretching has no effect on the order of precipitation of the reinforced phase in the figure.
Numerous researchers [
16,
17,
18,
19,
20,
21,
22] have studied the precipitation kinetics and yield strength calculation models for aluminum alloys during aging in some depth, and after generalizing and fitting the experimental data, Luiggi et al. [
23] obtained an empirical equation for the precipitation kinetics of the organization (JMAK equation). In this section, the JMAK equation is used to calculate the volume fraction and precipitation rate of the second phase precipitated from Al-Zn-Mg-Cu alloy during aging and to calculate the precipitation activation energy and constant of the second phase.
Hardness measurements were used to study the precipitation kinetics of the second phase by considering the hardness values from the start of aging (HSA) to the peak aging hardness (HP). Thus, Equation (1) can be used to calculate the volume fraction
f from the hardness (H) values, and the phase transition can then be investigated using the JMAK model shown in Equation (2), which has been successful in calculating the activation energy of the phase transition for different alloys, Q, based on Equation (3):
The implicit function containing the volume fraction of the precipitated phase is:
It can be found that the kinetic equation for tissue precipitation is:
where dT/dt is the heating rate during isothermal aging.
Before calculating the kinetic equations for the GP zone, η′ phase, and η phase under different deformation heat treatment conditions, the value of
n must be discussed. The relationship between the phase transformation mechanism of the alloy and the value of n is given in
Table 4, which provides a reference for the selection of
n. From
Table 4, it can be seen that the value of
n increases with the increase of the precipitation rate. According to the JMAK equation [
23] in the transformation mechanism of long-range diffusion-controlled growth, the phase transition constants of the GP region and η′ and η phases are selected as
n = 2/3,
n = 1, and
n = 1, respectively.
Taking the unstretched DSC curve
Figure 10a as an example, it can be analyzed that the formation temperature range interval of the GP zone is 53.982 °C to 86.09 °C, the formation temperature range interval of the η′ phase is 153.023 °C to 176.763 °C, and that of the η phase is 198.033 °C to 226.147 °C. The overlapping peaks were separated, and the precipitation rate was determined by Equation (3), which is the same as that of the second phase after 475 °C+3 h solid solution quenching treatment. After the separation of the overlapping peaks, the precipitation volume fraction and precipitation rate of the second phase after 475 °C + 3 h solid solution quenching treatment can be obtained according to the separation results and Equation (3), and the results are shown in
Figure 10. During the aging process, the volume fraction of the precipitated phase will fluctuate with the changes in the size and quantity of the second phase. In
Figure 11a, the volume fraction of the precipitated phase gradually increases as the amount of pre-deformation increases, but when the amount of deformation reaches 7%, this adversely affects the properties of the alloy. During the aging process, due to the increase in dislocations introduced under 7% pre-deformation conditions, it provides a channel for the precipitation of solute atoms, which results in the fastest precipitation of the second phase after the same aging time, as shown in
Figure 11b.
In Equation (5), n needs to be chosen according to the nucleation and growth mechanism. After that, the expression of F was derived to make a plot of ln[(df/dt)(v/F)] versus 1/T, whose slopes are the activation energies of precipitation in the GP region and η′ and η phases, as shown in
Figure 12. Based on the slopes of
Figure 12 and Equation (5), the kinetic parameters of precipitation in the GP region and η′ and η′ phases are derived as shown in
Table 5, with the activation energies of the precipitation in the GP region and η′ and η phases of 102.059 kJ·mol
−1, 178.147 kJ·mol
−1, and 241.622 kJ·mol
−1, respectively. The constants k0 are 4.14 × 1018 s
−1, 2.87 × 1023 s
−1, and 2.64 × 1028 s
−1, respectively.
4.2. Strengthening Mechanism of Pre-Deformed Al-Zn-Mg-Cu Alloys
Precipitation strengthening is the main strengthening mode of Al-Zn-Mg-Cu alloy, and its high strength and toughness mainly depend on the type, quantity, size, volume fraction, and distribution of intracrystalline and grain boundary precipitation phases (GP zone, η′ phase, η phase) after aging, etc. GP zone, η′ phase, and η phase are the main strengthening phases of Al-Zn-Mg-Cu alloy.
The precipitation strengthening mechanism is mainly categorized into dislocation cut-through mechanism and dislocation bypassing mechanism. The critical size of the dislocation cut-through mechanism and bypassing mechanism is about 2 nm [
24]. The TEM statistical results are displayed in
Figure 10, which shows that the average size distribution of the precipitated phase shows a normal distribution, and the average size of the particles in all states is larger than the critical value, so the Orowan dislocation bypassing mechanism is used to explain the reinforcement mechanism [
25,
26].
The expression is as follows [
7]:
where β is a constant. f and diameter are the precipitated phase volume fraction and radius. It can be seen that the volume fraction of the precipitated phase is proportional to the strength, and the size of the precipitated phase is inversely proportional to the strength. The average size statistics of the precipitated phase under different pre-deformation conditions are shown in
Table 6 below:
Combined with the DSC curves under different pre-deformation conditions in
Figure 10, it can be seen that the intensity of the dissolution peaks in the GP region increases in the pre-stretched samples with deformation amounts of 2%, 5%, and 7% compared with to the un-pre-stretched samples, and the intensity of the precipitation peaks of the η′ phases is higher than that of the un-pre-stretched samples, which means that the volume fractions of the GP region and the η′ phases in the pre-stretched peak-ageing samples are more than that of the un-pre-stretched samples. The area of the dissolution peak of η′ phase also increases with the increase of deformation, which suggests that many fine η′ phases are mainly formed when peak aging is reached, and these η′ phases are unstable and undergo dissolution during the DSC warming process. It is noteworthy that the temperature of the dissolution peak corresponding to the dissolution of the η′ phase of the 7% sample is significantly shifted to the high temperature direction, which indicates that the η′ phase of the artificially aged tissue is larger in size at 7% pre-stretching, and a higher temperature is required for its dissolution, which is in agreement with the results demonstrated in the previous TEM images. Moreover, the areas of the precipitation and coarsening peaks of the η-phase in the 7% sample are significantly larger than those in the 2% and 5% samples. This is due to the presence of a large number of dislocations as the non-uniform nucleation points of the η-phase, which leads to rapid precipitation of the η-phase and significant coarsening of the η-phase. Therefore, the areas of the exothermic and heat-absorbing peaks corresponding to η-phase precipitation of the η-phase precipitation are larger in the pre-stretched sample.
In conclusion, according to the dimensions of the precipitated phases and the DSC curves in
Table 6, it can be concluded that the average size of the precipitated phases is the smallest, and the volume fraction reaches a high level when the pre-deformation is 2%. According to Equation (6), it can be analyzed that the precipitated phases in the structure of 2% specimens contribute the most to the strengthening of the alloy, which corresponds to the value of the tensile mechanical properties.
In order to investigate the effect of different pre-deformation amounts on the peak aging dislocation density,
Figure 13 and
Figure 14 demonstrate the average grain orientation difference plots and the statistical histograms of KAM values of the alloys after different pre-deformation treatments. At 0% to 7% pre-deformation, the local orientation difference of the peak aging state samples is significantly higher, and its average KAM value increases from 0.61 to 1.02. According to the histogram of KAM distribution, it can be seen that the KAM values have higher peaks at smaller angles, and with the increase of pre-deformation, the KAM values are more distributed at smaller angles. Theoretically, the local orientation difference angle and geometrically necessary dislocation density (ρ_GND) can be calculated using Equation (7) [
7,
27], and the EBSD scanning step value for all the specimens 1.2 μm.
The calculated geometrically necessary dislocation densities of the specimens under different pre-deformation conditions are shown in the histogram of the KAM distribution in
Figure 14. As the amount of pre-deformation rises from 0% to 7%, which shows that the geometrically necessary dislocation densities of the samples in the peak aging state increase as the amount of pre-deformation rises from 0% to 7%, implying that the total dislocation densities within the alloys increase. This is partly due to the increase in the amount of pre-deformation, which introduces a large number of dislocations and increases the dislocation density, which is—on one hand—due to the fact that the increase of the amount of pre-deformation introduces a large number of dislocations that increase the dislocation densities., On the other hand, this is due to the introduction of a large number of dislocations by the increase of pre-deformation, which increases the dislocation density, and on the other hand, it is due to the pinning effect of the fine precipitates relative to the dislocations in the aging process, which inhibits the dislocations’ movement and increases the dislocation density.
The quantitative calculation of dislocation density allows the magnitude of its contribution to the strength of the material [
14]:
where M is the mean orientation factor (M = 3), α is a constant, G is the shear modulus (G = 26.1 GPa), b is the Parker vector (=0.286 nm), and ρ is the dislocation density. It can be seen that the alloy strength is proportional to the dislocation density. Thus, the contribution of dislocation strengthening to the strength of alloys in different deformed heat-treated states increases with the amount of pre-deformation. These dislocations not only provide channels for the precipitation of solute atoms but also increase the strength of the alloy. When the amount of deformation is increased to a certain extent, the dislocations inside the alloy are entangled and piled up with each other, and they impede the dislocation slip, so the strength of the alloy increases.