The Effect of Pre-Deformation on the Microstructure and Hardness of Al-Zn-Mg-Cu Alloy

Abstract

In this paper, hot rolling pre-deformation treatment was applied to an Al-Zn-Mg-Cu alloy after solid solution treatment, followed by peak aging treatment. The effect of the degree of deformation was discussed. The microstructure of the alloy after treatment was observed and the mechanical properties were tested. The results indicate that after hot rolling pre-deformation, high-density dislocations are introduced within the grains of the Al-Zn-Mg-Cu alloy, and the dislocation density gradually increases with the degree of deformation. At the same time, with the increase of rolling deformation, the alloy hardness first increases and then decreases. When the deformation is 40%, the alloy hardness reaches a peak value of 101.7 HV. In the subsequent aging process, with the increase of deformation, the time required to reach peak aging is gradually shortened, and at 40% deformation, the alloy hardness reaches a peak of 99.7 HV after 12 h of aging. Moreover, the dislocations generated by pre-deformation can become entangled around the grain boundary and the coarse quenching precipitated phase, providing nucleation particles for the subsequent precipitation of the strengthened phase, effectively improving the precipitation strengthening effect of the alloy during aging, and thus improving the hardness of the alloy at the peak aging state. This study provides a research idea for improving the hardness of the alloy and expands the application of the deformation aging process in Al-Zn-Mg-Cu alloys.

1. Introduction

The Al-Zn-Mg-Cu alloy, as a typical heat-treatable strengthened aluminum alloy, has attracted significant attention due to its ultra-high strength and excellent overall performance [1,2,3,4,5]. It is an essential structural material in aerospace, vehicles, and weaponry, holding an important strategic position in national defense and economic development. However, as the demands for material performance continue to rise, research on how to further enhance the overall performance of high-strength, high-toughness Al-Zn-Mg-Cu aluminum alloy materials through relevant processes has become a topic of study for scholars.

Firstly, regarding the strength enhancement of Al-Zn-Mg-Cu alloys, researchers have developed a series of improved alloys by adjusting the content of alloying elements. Tang et al. [6] found that as the Zn content increased from 20% to 40%, the grain size of the aluminum alloy decreased from 45 μm to 20 μm, and the volume fraction of the non-equilibrium phase α + η on the grain boundaries increased from 9% to 28%, which results in a significant increase in the alloy’s tensile strength and yield strength. Chen et al. [7] found that after adding trace amounts of Gd to the 7056 series aluminum alloy, the grains were significantly refined, a result of the pinning effect of the dispersed phase Al₃(Gd,Zr) suppressing the growth of recrystallized grains, thereby significantly improving the alloy’s performance.

In addition to the alloy composition, researchers have also controlled the alloy microstructure by applying rolling pre-deformation treatment to Al-Zn-Mg-Cu alloys. Liu et al. [8] studied the effects of deformation degree and rolling temperature on the recrystallization degree of Al-Zn-Mg-Cu alloys, which reveals that as the deformation degree increases, the recrystallization degree continuously increases, while the recrystallization degree gradually decreases with increasing rolling temperature. Ultimately, a best parameter of 440 °C and 80% deformation degree was determined. Chen et al. [9] synchronized solution treatment and hot rolling, which achieved good strength–plasticity matching in Al-Zn-Mg-Cu alloys. Detailed research has found that deformation can affect the morphology, orientation, size, quantity, and density of precipitate in Al-Zn-Mg-Cu alloys. Sha et al. [10] employed ECAP, which promotes the formation of equiaxed η precipitates, effectively optimizing the strengthening effect of precipitates. Mei et al. [11] found that after pre-aging treatment, cold rolling and warm rolling of Al-Mg-Si alloys followed by subsequent aging treatment promote the formation of strengthening phases, thereby increasing the strength of the alloy. Additionally, the presence of nanocrystalline grains and the reduction of dislocation density resulting from deformation also play important roles in the ductility of the material. Yu et al. [12] discovered that as the hot rolling deformation degree increases, the grains become finer and flatter, the grain orientations tend to become more uniform, and the density of low-angle grain boundaries increases. Furthermore, hot rolling can significantly alter the proportion of internal texture in aluminum alloys, refining the recrystallized structure of the alloy and improving its properties. Despite numerous studies on the effects of thermal deformation on the microstructure and properties of alloys, the underlying mechanisms remain unclear.

Therefore, in this paper, the effects and mechanisms of hot rolling deformation on the precipitation behavior of Al-Zn-Mg-Cu alloys are studied. The relationship between dislocation density and deformation degree is explored. Then, artificial aging treatment is conducted to investigate the optimal aging time under different deformation degrees. By establishing the relationship between hot rolling deformation degree and dislocation density, as well as its impact on the corresponding microstructure and mechanical properties, the effects and mechanisms of hot rolling deformation on the precipitation behavior of Al-Zn-Mg-Cu alloys are explored.

2. Materials and Methods

2.1. Materials

The material used in this paper is spray-formed Al-Zn-Mg-Cu alloy provided by CITIC Dicastal, and the alloy composition is listed in

Table 1. Chemical composition of Al-Zn-Mg-Cu Alloy (Mass Fraction, %).

Alloy	Al	Zn	Cu	Mg	Zr	Others
Al-Zn-Mg-Cu Alloy	90.50	6.14	1.70	1.36	0.17	0.13

.

2.2. Heat Treatment and Deformation Processing

The alloy underwent the following experimental procedures. One set of samples remained untreated. The other samples underwent a solid solution treatment at 475 °C for 6 h. A portion of these samples was then subjected to a hot rolling treatment at a temperature of 450 °C, with a holding time of 1 h. The holding time for each pass in the furnace was 2 min, and the rolling amounts were 20%, 40%, and 60%, respectively. Then, they underwent artificial aging treatment at 120 °C for 6 h, 12 h, 18 h, and 24 h. For convenience, the samples are labeled as: the original sprayed state (SP), the solid solution state (SS), the solid solution + pressing amount 20% state (SS + 20%HR), the solid solution + pressing amount 40% state (SS + 40%HR), the solid solution + pressing amount 60% state (SS + 60%HR), the solid solution + aging state (SS + AA), the solid solution + pressing amount 20% + aging state (SS + 20%HR + AA), the solid solution + pressing amount 40% + aging state (SS + 40%HR + AA), and the solid solution + pressing amount 60% + aging state (SS + 60%HR + AA).

2.3. Testing and Characterization

The samples were subjected to etching using Keller’s reagent, and then the microstructure of the three surfaces of the sample (ND × RD surface, TD × ND surface, RD × TD surface, as shown in Figure 1) were observed using an Axio Observer 3M inverted metallurgical microscope.

Using the Hitachi S-3400N Field Emission Scanning Electron Microscope (SEM, manufactured by Hitachi, Tokyo, Japan), the corroded samples were observed. The scanning electron microscope has a large depth of field and can continuously magnify within the range of 2 to 1,000,000×, with an acceleration voltage of 30 kV. At the same time, energy dispersive X-ray spectroscopy (EDS) was used for qualitative analysis of the composition of the precipitates. Phase detection was performed using a D/MAX-2500/PC X-ray diffractometer (XRD, manufactured by Rigaku International Corporation, Tokyo, Japan), with a tube voltage of 40 kV and a tube current of 150 mA, utilizing CuKα radiation. The scanning angle 2θ was set from 20° to 100°, with a scanning step size of 2°/min. In addition, the half-peak width was calculated from the XRD data to determine the dislocation density, thereby establishing the relationship between deformation and dislocation density.

Using the FM-ARS 9000 microhardness tester (manufactured by Future-Tech, Kawasaki, Japan), the hardness of the sample was tested under a load of 200 g and a holding time of 10 s. For each 10 mm × 10 mm × 10 mm sample, continuous rectangular matrix tests were conducted at 15 different positions in the same area, and then the average value was calculated to obtain the average hardness of each sample [13].

3. Results and Discussion

3.1. The Impact of Hot Rolling Deformation on the Microstructure and Properties of Alloys

3.1.1. Metallographic Observations

The metallographic microstructure of the spray-formed Al-Zn-Mg-Cu alloy is shown in Figure 2, which contains solid solution, intermetallic compounds, and a small amount of non-metallic inclusions. The grains are equiaxed and their sizes are uniform.

Figure 3 shows the microstructure after rolling. When the pressing amount is 20%, it can be seen from Figure 3a–c that the grain size of the Al-Zn-Mg-Cu alloy is smaller than that in the original sample. The grains are elongated along the ND × RD direction, with only small equiaxed grains observed. When the pressing amount is 40%, it can be seen from Figure 3d–f that compared to the spray-formed Al-Zn-Mg-Cu alloy sample, the grains are significantly refined, and some small equiaxed grains are formed near the refined grains. This phenomenon is particularly evident in the ND × RD direction. When the pressing amount is 60%, it can be seen from Figure 3g,h,r that there is also a significant change in grain morphology along the ND × RD direction, where the grains become longer and thinner, and are being fragmented. The grain boundaries become blurred, which results in a fibrous structure accompanied by the formation of clustered structures [14]. In summary, all grains are elongated perpendicular to the rolling direction after rolling. The grain boundaries become discontinuous and unclear. As the pressing amount increases, the degree of grain fragmentation increases, and more small equiaxed grains appear on the elongated grain boundaries.

During hot rolling deformation, dynamic recovery and recrystallization occur. With the increase in pressing amount, the dislocation density in the Al-Zn-Mg-Cu alloy structure increases, with dislocation pile-ups at the grain boundaries. When the pressing amount is 40%, dynamic recovery occurs, and the dislocation entanglement gradually evolves into sub-grains and dislocation cells. When the pressing amount is 60%, the dislocation cells grow, but their boundaries remain pinned by precipitates. The hot rolling pre-deformation promotes the expansion and interaction of dislocations, leading to the refinement of elongated grains and the formation of small sub-grains [15].

3.1.2. Scanning Electron Microscopy Observations

Figure 4 shows the SEM morphology of the surface of the “SS + 60%HR” sample. From the image, it can be seen that there is a large amount of precipitates distributed in the alloy. A precipitate marked by the red square was analyzed, and its elemental content is listed in Figure 4b. The content percentage of Mg is 4.53%, and the content percentage of Zn is 4.28%. It can be seen that the Mg/Zn ratio is between 1 and 1.4. According to the previous literature, the precipitate is ita.

3.1.3. X-Ray Diffraction Detection

Figure 5 shows the XRD diffraction patterns after hot rolling deformation. The five diffraction peaks of each curve correspond to the (111), (200), (220), (311), and (222) planes. According to the XRD pattern, it can be seen that the rolling amount has no effect on the phase composition, but the peak width varies. It can be observed that as the pressing amount increases, the peak intensity decreases and the full width at half maximum (FWHM) also increases, indicating that the internal defects in the Al-Zn-Mg-Cu alloy are increased [16].

The dislocation density of the samples after hot rolling deformation were calculated using the Williamson–Hall model. Based on the X-ray diffraction spectrum in Figure 5, the Gaussian fitting of the refined results for each diffraction peak is conducted using Origin 2021 software. From the fitting results, the corresponding 2θ and FWHM for each diffraction peak can be obtained. According to the Scherrer formula, the grain size can be related to the FWHM, leading to the following formula:

$$\begin{array}{c}d={\displaystyle \frac{k\lambda }{\left(\delta 2\theta \right)\mathit{cos}{\theta }_0}}\end{array}$$

(1)

In the formula, d represents the grain size (the size of the XRD coherent diffraction region), k is a constant, λ is the wavelength of CuKα radiation (with a value of 1.54056 Å), $\delta 2\theta$ is the full width at half maximum (FWHM), and θ₀ is the θ angle corresponding to the peak value of the diffraction 0 peak.

The dislocation density can be obtained by examining the grain size and lattice microstrain within the material. According to the Williamson–Hall method, data on the grain size and lattice microstrain of the alloy after certain processing can be obtained from the integral broadening of the XRD spectrum of the test material. The broadening caused by grain size and lattice microstrain can be approximated using Cauchy and Gaussian functions, respectively [17,18], leading to the following formula:

$$\begin{array}{c}{\displaystyle \frac{{\left(\delta 2\theta \right)}^2}{{\mathit{tan}}^2{\theta }_0}}=25<e^2>+{\displaystyle \frac{\lambda }{d}}\left({\displaystyle \frac{\delta 2\theta }{tan{\theta }_0\mathit{sin}{\theta }_0}}\right)\end{array}$$

(2)

In the formula: $<e^2>$ represents the lattice microstrain, and the terms in the formula are replaced with the following letters.

Assumption:

$$\begin{array}{c}Y={\displaystyle \frac{{\left(\delta 2\theta \right)}^2}{{\mathit{tan}}^2{\theta }_0}}\end{array}$$

(3)

$$\begin{array}{c}A=25<e^2>\end{array}$$

(4)

$$\begin{array}{c}B={\displaystyle \frac{\lambda }{d}}\end{array}$$

(5)

$$\begin{array}{c}X={\displaystyle \frac{\delta 2\theta }{\mathit{tan}{\theta }_0\mathit{sin}{\theta }_0}}\end{array}$$

(6)

Thus, the Formula (2) is simplified to:

$$\begin{array}{c}Y=A+BX\end{array}$$

(7)

Substituting the peak positions obtained from the Gaussian fitting and their corresponding half-peak widths (converted to radians) into Equations (3) and (6), we can obtain the coordinates of the five diffraction peaks, denoted as (X1, Y1), (X2, Y2), (X3, Y3), (X4, Y4), and (X5, Y5). By inputting the coordinates of these five points into Origin for plotting and performing linear fitting (the fitting results at 20%, 40%, 60% are shown in Figure 6a,b,c, respectively.), we can obtain the intercept A and slope B of the fitted line. By substituting the values of A and B into Equations (4) and (5), we can calculate the lattice microstrain and the size of the XRD coherent diffraction region. Then, substituting into Equation (8), we can calculate the dislocation density (ρ):

$$\begin{array}{c}\rho ={\displaystyle \frac{2\sqrt{3}{⟨e^2⟩}^2}{db}}\end{array}$$

(8)

In this, b is the Burgers vector of Al, taken as 0.286 nm. The dislocation density at various pressing amounts is shown in Figure 6d. According to Figure 6d, it can be observed that as the pressing amounts increase, the dislocation density continues to increase, indicating that hot rolling deformation introduces a large number of dislocations. At the same time, as the dislocation density increases, the volume fraction of supersaturated solute atoms and vacancies diffusing into dislocations through short-circuit diffusion increases. This leads to a corresponding decrease in the uniformly precipitated GP zones and η’ phases in the matrix, while the formation of the η equilibrium phase at dislocations increases [19,20]. This reveals the reason for the decrease in alloy hardness shown in Figure 7 when the deformation amount exceeds 40%.

3.1.4. Hardness Testing in Various Rolling Directions

Hardness tests were conducted on three different surfaces of the samples after rolling, as shown in Figure 7. It can be observed that the hardness initially increases and then decreases with the increase in the pressing amounts [21,22]. The hardness reaches its maximum value when the deformation amount is 40%. When the deformation amount is constant, the hardness in the ND×RD direction is always the highest.

The hardness of the spray-formed Al-Zn-Mg-Cu alloy samples SS and “SS + HR” on the ND×RD surfaces are shown in Figure 8. It was found that the hardness of the samples generally decreased after rolling. This is because dynamic recrystallization occurred during the hot rolling process, where work hardening and recrystallization softening coexist, and the effect of softening is greater than that of hardening.

3.2. The Effect of Deformation on the Aging Microstructure and Properties of Alloys

3.2.1. Age Hardening

Figure 9 shows the aging hardening curves for SS + 20%HR + AA, SS + 40%HR + AA, and SS + 60%HR + AA, with hardness measured on the ND × RD surface. For the sample with a deformation of 20%, the hardness first decreases and then increases with aging time, which reaches peak hardness at 18 h. For the samples with 40% and 60% deformation, the hardness shows a trend of first increasing and then decreasing with aging time, with the former reaching its peak hardness at 18 h and the latter at 12 h. The analysis indicates that the reason for this behavior is that oversaturated solid solution atoms and vacancies form G P zones and η′ phases, significantly enhancing the strength properties of the alloy. However, when dislocations are present in the matrix, elastic interactions between dislocations, solute atoms, and vacancies cause the solute atoms and vacancies around the dislocations to diffuse into the dislocations, forming coarse equilibrium phases, thereby reducing the aging strengthening effect of the alloy. Furthermore, as the amount of deformation increases, leading to an increase in dislocation density in the matrix, the hardness shows a gradual decreasing trend.

Comparing the aging curves under the three pressing amounts, increasing the deformation not only increases the hardness but also causes the samples to reach the peak aging hardness earlier. During the aging process, the increase in hardness is mainly due to the formation of precipitates. Dislocations serve as the core for the heterogeneous nucleation of the precipitates, which promotes nucleation. Therefore, the increase in dislocation density accelerates the precipitation process, leading to an increase in hardness and allowing the samples to reach the peak aging hardness sooner. Ultimately, the peak artificial aging time is set at 18 h for the sample with 20% deformation, 18 h for the sample with 40% deformation, and 12 h for the sample with 60% deformation.

3.2.2. Metallographic Structure

The microstructure after aging was observed, and its metallographic structure is shown in Figure 10. Figure 10a shows the metallographic structure of “SS + AA”, while Figure 10b–d show the aged structures with pressing amounts of 20%, 40%, and 60%, respectively. With the change in pressing amounts, the size of the grains also changes, and fine grains are observed to form within the structure. The majority of the crystal boundaries are unclear. As shown in Figure 10c, there are a large number of new grains smaller than the original grains at the boundaries of the elongated grains. In Figure 10d, a large number of disc-like sub-grain clusters oriented perpendicular to the compression axis can be seen [8], and this structure increased after aging.

3.2.3. Characterization Results of Scanning Electron Microscopy

Figure 11 shows the surface morphology of the “SS + 60%HR + AA” sample. According to the scanning results, it is observed in Figure 11a that the grain boundaries are no longer continuous. To further determine whether the precipitates are changed after aging, a compositional analysis of one precipitate (marked by a circle) in Figure 11a was conducted, revealing that the composition of the precipitated phase is still MgZn₂.

3.2.4. Hardness Test

The hardness of the “SS,” “SS + HR,” “SS + AA,” and “SS + HR + AA” samples on the ND×RD surface are shown in Figure 12. The hardness value of the “SS + AA” sample is the highest. The hardness decreased after rolling deformation.

The analysis of the hardness of the “SS + HR” samples after rolling shows that as the amount of deformation increases, the hardness first increases and then decreases. This can be explained by the calculation of dislocation density mentioned earlier. When the deformation is at 20% and 40%, the dislocation density increases gradually with the amount of deformation. However, when the deformation reaches 60%, the increase in dislocation density becomes significantly pronounced. Therefore, the reason for the change in hardness can be analyzed as follows: when the rolling deformation is small, an appropriate amount of dislocation is introduced into the crystal. The accumulation of dislocations creates a stress field within the grains, which hinders the slip and deformation of the lattice, thereby increasing the hardness of the alloy. However, as the amount of deformation increases, when the dislocation density becomes too high, the dislocations will begin to entangle and block each other, increasing the defect density in the material and affecting the intrinsic structure of the crystal, which in turn reduces the hardness of the alloy.

The reason for the decrease in hardness after aging may be that during the aging process, the metastable MgZn₂ phase gradually transforms into the stable MgZn₂ phase [23,24,25,26]. Since the main strengthening phase in Al-Zn-Mg-Cu alloys is the metastable MgZn₂, the stable precipitate phase after aging will affect the hardness of the alloy. Additionally, during the aging process, changes in the state of the precipitate phase and the coarsening of some stable precipitate phases can further reduce the hardness of the Al-Zn-Mg-Cu alloy. Furthermore, rolling may cause uneven chemical composition and microstructure within the alloy, thereby affecting the aging effect and leading to a decrease in alloy hardness. The coarsening of the GP zone during aging is also one of the reasons for the decrease in alloy hardness [27].

4. Conclusions

The Al-Zn-Mg-Cu alloy was subjected to hot rolling pre-deformation treatment. The microstructural characteristics after hot rolling treatment were characterized and studied, and its mechanical properties were tested, leading to the following conclusions:

(1)

After hot rolling pre-deformation treatment, the microstructure changes. The grains are elongated, fragmented, the grain boundaries become blurred, and a second phase appears at the grain boundaries in a banded distribution, with small, recrystallized grains forming. As the amount of deformation increases, the degree of grain elongation increases, the degree of grain fragmentation increases, and more recrystallized grains are present. At a deformation amount of 60%, the structure exhibits a fibrous distribution. After aging, the amount of the second phase increases.

(2)

Hot rolling pre-deformation will introduce dislocations. As the amount of deformation increases, the dislocation density increases. Dislocations provide energy and particles for the nucleation of the second phase during the aging process, promoting the precipitation of the precipitate phase.

(3)

During the hot rolling process, the main precipitated phase is metastable MgZn₂, which is the primary strengthening phase in the alloy, appearing in a plate-like shape and classified as a nanoscale precipitate, dispersed throughout the crystal. During the aging process, the precipitated phase gradually transitions to stable MgZn₂, which also exhibits a dispersed distribution within the grains.

(4)

During hot rolling, dynamic recrystallization occurs, and work hardening also takes place. At the hot rolling temperature of this study (450 °C), softening predominates, which explains the observed hardness of the samples after “solution treatment-rolling-aging”.

Hot rolling involves both hardening and softening, making it a complex high-temperature, dynamic, and instantaneous process. Various deformation mechanisms collectively determine the characteristics of aluminum alloy hot rolling deformation. Therefore, studying the pre-deformation in hot rolling presents certain challenges. In actual production, reasonably controlling the hot rolling temperature and time is of great significance for improving the performance of Al-Zn-Mg-Cu alloys.