The Influence of Creep Ageing on the Hardening Behavior and Microstructure of 7050 Aluminum Alloy

Abstract

The creep ageing process can have a significant influence on the mechanical properties of aluminum alloys. In the present work, microstructural analysis and mechanical testing were implemented to characterize the age hardening effect and microstructure evolution, and to investigate how the stress applied under creep ageing conditions can affect a material’s microstructure. The curves depicting yield strength in relation to creep ageing time suggested that the stress applied in creep ageing can result in a reduction of the strength of aluminum alloy 7050; the yield strength decreases with increasing applied stress. Microstructural analysis by transmission electron microscopy (TEM) revealed that by applying stress, the growth and coarsening rate of the Guinier-Preston (GP) zones and η′ precipitates can be sped up. Even after pure/creep ageing for 8 h, there are still some GP zones in the aluminum matrix, demonstrating that the GP zones’ nucleation is a continuous process.

1. Introduction

There is always an increasing demand for strong, lightweight materials for cars, airplanes, and other manufacturers. Aluminum alloys play a crucial part in this field. Age hardening contributes the greatest strengthening effect for these alloys [1]. For instance, when aluminum alloy AA7050 is in the annealed form, the tensile strength is at around 200 MPa, but with an ideal ageing process, it is possible to raise the strength up to 600 MPa [2,3,4]. For the purpose of strengthening the aluminum parts by age hardening, regular manufacturing procedures usually necessitate additional procedures of ageing treatment to enhance the mechanical properties of a part. Such extra processing, however, could raise production costs and make the product more difficult to make.

Creep age forming has a unique characteristic that makes it a promising approach for producing aluminum parts. For heat-treatable alloys, particularly aluminum alloys, the forming operation and age heat treatment can be performed simultaneously at a specific increased temperature.

The precipitation process of Al-Zn-Mg alloys often starts from the generation of Guinier-Preston (GP) zones [5,6]. Two forms of GP zones, GPI zones and GPII zones can typically coexist in the matrix. A GPI zone is one layer of highly concentrated Mg and Zn/Al metastable clusters that are coherent with the matrix [7,8,9]. GPI zones can occur anywhere between ambient temperature and 150 °C, regardless the temperature of quenching [10]. A GPII zone is two-layer clusters of Zinc that are packed by three-layer aluminum [11]. GPII zones form after solution treatment at above 450 °C, and at above 70 °C after ageing in Al-Zn-Mg alloys [8,12,13,14].

During the ageing process, those GP zones gradually convert to platelet-shaped η′ precipitates with a metastable hexagonal phase [15,16,17]. Finally, the η′ precipitates continue to grow and become the equilibrium $\eta$ precipitate phase (commonly referred to as MgZn₂, a Laves phase), with a morphology that is either plate-like or rod-like [18] based upon the plane of observation. The $\eta$ precipitates have a different crystal structure from the aluminum matrix and lower free energy than the GP zones and ${\eta }^{\prime }$ precipitates.

Due to the higher free energy the GP zones and ${\eta }^{\prime }$ precipitate possess, they can easily form $\eta$ phase precipitates—the energy barriers at the interface boundary between the ${\eta }^{\prime }$ phase and the matrix is low and can be overcome easily because of high coherency [19]. The sequence of precipitation phase transformation is shown below [20]:

Solid solution → GP zones(I and II) → ${\eta }^{\prime }$ precipitates → $\eta$ precipitates

It is not always possible to obtain the required mechanical properties with just one quenching and ageing treatment. For instance, despite the fact that optimally sized and dense precipitates may form within the grains in a one-step ageing treatment, precipitate-free zones (PFZ) may also form near the sites of vacancy sinks and adjacent to the grain boundaries. Such PFZs can be detrimental to a material’s mechanical properties [21]. To address the aforementioned issue, two-step ageing processes have been developed [22]. For Al-Zn-Mg alloys, the two-step ageing process can produce a more refined and dispersed precipitate microstructure with superior properties.

The two-step ageing process generally begins after a solution treatment. In the first step of this process, GP zones start to form, and some of them gradually convert to ${\eta }^{\prime }$ precipitates. In the next step, the transformation rate accelerates, and therefore the time required can be reduced for reaching the optimum strength from the change of GP zones to ${\eta }^{\prime }$ precipitates. In the two-step ageing process, while the GP zones continue to transform into ${\eta }^{\prime }$ precipitates, the ageing temperature needs to be carefully managed so as to maximize the profuse growth of nanometer-sized precipitates.

With this two-step ageing process, the aluminum alloys have been heat-treated with the second step temperature higher than the first step temperature. They were reportedly found to possess better strength than those heat-treated with both temperatures [23,24,25]. This is due to the fact that GP zones have low surface energy, so at low ageing temperatures, the likelihood increases for large number densities of GP zones or extremely fine and disperse transition phases to occur, which is advantageous for the ${\eta }^{\prime }$ precipitates to nucleate at the second step ageing [22].

When the temperature in the first step of ageing is higher than that in the second step, it may be high enough for the stable phases to form. It is not likely that, at the second step with a lower temperature, they would dissolve and form GP zones again [26]. Therefore, it is fair to draw the conclusion that to produce very dense GP zones in the matrix, the temperature of the first step of ageing should be lower than that of the second so that substantial amounts of dispersed ${\eta }^{\prime }$ precipitates can be formed in the second step.

It has been reported that in alloy AA7475, the creep deformation rate rises during the second step of ageing in the forming operation [27]. This is presumably due to the GP zones that formed previously in the first step dissolved in the second and nucleated into ${\eta }^{\prime }$ precipitates [28]. It should be noted that a higher creep deformation rate is advantageous in the metal forming operation for it can reduce springback after forming.

In Al-Zn-Mg alloys, GP zones (coherent to the matrix) and ${\eta }^{\prime }$ precipitates (semi-coherent to the matrix) are the main strengthening phases. At the start of the ageing process, the lattice distortion that occurs between the matrix and GP zones generates lattice stresses. These stresses lead to coherency hardening that can impede dislocation movement as they attempt to cut through. The alloy’s strength is thereby increased. As the ageing process goes on, the GP zones continue to grow and become ${\eta }^{\prime }$ precipitates (semi-coherent to the matrix) that are between 5 and 20 nm in size. ${\eta }^{\prime }$ precipitates themselves and the lattice strain fields around them can inhibit the mobility of dislocations, which results in the alloys being strengthened.

Dislocations are often able to penetrate and cut through smaller particles with coherent boundaries because these coherent boundaries are easier to pass through. When facing large incoherent particles, on the other hand, dislocations tend to bow and bypass so as to avoid the disordered interphase boundaries. This is known as the Orowan dislocation bowing mechanism. Such a mechanism occurs when the particles reach a particular peak size or when there is a high lattice disorder in the interphase boundary. The influence of age hardening on yield strength is an incremental process that depends on whether the dislocation mechanism of the material is predominantly “cutting” or “bowing”.

As the ageing process goes to a later point, the ${\eta }^{\prime }$ precipitates grow into less dense and large-sized $\eta$ precipitates (incoherent to the matrix), which are roughly 50 nanometers or more in size. The large size and incoherency of the $\eta$ precipitates block the dislocations from cutting through them. Instead, the dislocations can only move by bending between and bowing around the precipitates, which can result in an abrupt change in the curvature of the dislocation line. Dislocation loops, as a result, develop around precipitates in the same manner as they would at a Frank-Read source. After passing by those $\eta$ precipitates, the dislocations may continue moving, but they will leave dislocation loops encircling the precipitates. The stress fields of the loops increase the lattice distortion, making the subsequent dislocation movement more difficult.

As the $\eta$ precipitates continue to coarsen and grow, the material enters the over-ageing condition; the spacing between each precipitate increases and the number density decreases. Consequently (according to the dispersion strengthening mechanism), the stress required to move the dislocation drops. Since the yield strength can be seen as the shear stress necessary to bend dislocations among precipitates, this stress is inversely related to the distance of two precipitates, as expressed in the formula $\tau =Gb/L$, where $\tau$ is the stress applied, $G$ is shear modulus, $b$ is burger vector, and $L$ is the distance between particles.

Due to the influence of stress, creep aged alloy is expected to have more dislocations than the pure aged alloy. The dislocations can serve as nucleation sites for the GP zones and ${\eta }^{\prime }$/$\eta$ precipitates. The relevant studies, however, have rarely discussed about how the influence of stress under the creep ageing process can affect the microstructures of an aluminum alloy and its strength. The present study—by employing ageing treatment, creep ageing tests, and transmission electron microscopy (TEM)—investigated and characterized the age hardening effect, the microstructure evolution, and the ways the stress levels in creep ageing conditions can affect the microstructures for the AA7050-T6 alloy.

2. Materials and Methods

2.1. Materials

Aluminum alloy AA7050-T6 was used for this investigation. This alloy was made following the AMS 4050 standard [29]. It has high fracture toughness and exceptional mechanical properties.

Table 1. Average chemical composition of AA7050-T6 (wt%).

Other	Si	Fe	Cu	Mg	Zn	Al
0.15	0.03	0.05	2.23	2.14	6.25	Bal.

lists the chemical composition of this alloy.

The heat treatment history and manufacturing procedure of this material follows the AMS 2770 standard [30] with some slight adjustments to the second step ageing time. Such adjustments were intended to make the alloy achieve an intermediate state for the subsequent creep ageing process.

Based on AMS 2770, the alloy underwent a solution treatment at 477 °C for 1 h and was then rapid quenched in water. After that, it was pure aged in two stages: the first for 8 h at 121 °C and the second for 1 h at 165 °C. When the alloy reached this state, it required an additional 8 h of ageing at 174 °C so as to achieve a state comparable to the T74 temper condition [31].

2.2. Pure Ageing and Creep Ageing

The mechanical testing included pure ageing (without stress applied) and creep ageing tests (with stress applied). Pure ageing tests were performed in an oven with fan-forced air circulation. Creep ageing tests were performed in a uniaxial tensile constant load creep test apparatus at 174 °C, the optimum ageing temperature. The creep test apparatus applies a fulcrum loading mechanism with upper and lower fixture assemblies, and a furnace is integral to its structure. By lowering the lower fixture to raise the weight, the test specimen can be loaded by the fulcrum loading mechanism. The fulcrum arm has a load ratio of 10:1.

For the creep ageing tests, single-electrode capacitance gauges were employed to measure strain. Capacitance gauges can measure the changes of electric fields between the probe of the gauges and a conductive specimen surface, with no contacting the specimen physically. These gauges can measure the voltage differences as tiny as 1 mV, which translates to a displacement of around 0.5 μm. Two linear capacitance gauges were attached to the ridges of the specimens to detect the displacement of the creep strain.

Uniaxial creep ageing and pure ageing tests were performed at 174 °C in air over a predetermined period of time. The durations and temperatures of these tests were marginally altered from AMS 2770 [30] to maximize the strength of the alloy. Uniaxial creep ageing tests were performed at an intermediate temperature of 174 °C for 2, 4, 6, and 8 h, and at eight different stress levels from 75 to 175 MPa, respectively.

2.3. Tensile Tests

Instron 5584 was used to perform the tensile tests on the pure and creep aged specimens at ambient temperature, and the strain rate was set at 10⁻⁴ s⁻¹. The objective of the tensile tests was to assess the influence of pure/creep ageing time and the stress applied in creep ageing on the yield strength variation of the alloy. To capture the strain during tests, an extensometer was attached to the specimen during the test.

2.4. Transmission Electron Microscopy (TEM)

Transmission electron microscopy (TEM) was used to investigate the microstructural evolution of the specimens after pure/creep ageing tests. Since the pure/creep ageing treatment is designed for the AA7050-T6 alloy to attain T7 temper after 8 h of ageing, TEM was only carried out on specimens that had been pure/creep aged for up to 8 h. The TEM specimens were first sliced along the applied stress axes from gage sections. They were then stamped into 3-mm diameter discs and milled mechanically to thin foils with a thickness of 0.08 mm. Then, they were electropolished to perforation using a twin-jet electro-polishing apparatus in a solution of 1/3 nitric acid and 2/3 methanol at 10 V and −20 °C. Finally, the electropolished specimens were analyzed with a 200 kV Tecnai F20 field emission gun (FEG) scanning TEM fitted with an energy dispersive X-ray (EDX) spectrometer.

3. Results

3.1. Creep and Tensile Properties

Figure 1 illustrates the curves of creep strain vs. ageing time at various load conditions. Primary creep typically occurred within 1 h under all stress conditions. In this alloy, secondary creep was found to occur under creep ageing conditions of less than 162.5 MPa and less than 8 h. Tertiary creep occurred, however, after 5 h when the applied stress reached 175 MPa.

The room temperature tensile test was performed at a strain rate of 10⁻⁴ s⁻¹. Figure 2 depicts the stress vs. strain curves for the pure aged (2–8 h) and as-received specimens. It can be observed that, with increasing ageing time, the elongation increased but the flow stress decreased. The as-received specimen possessed the maximum yield strength at around 574 MPa and ultimate tensile strength at around 612 MPa, and both of the strength decreased with increasing ageing time.

The elongation at break conversely decreased from the as-received condition of 18.7% to 18.1% after pure ageing for 2 h, and later increased to 18.4% after pure ageing for 4 h, 19.1% after pure ageing for 6 h, and 20.6% after pure ageing for 8 h. Figure 3 depicts the relationship of yield strength (0.2% offset) vs. creep ageing time. In general, the material presented an over-ageing behavior; the yield strength dropped as ageing time rose. The highest yield strength of the alloy, about 574 MPa, was achieved at the as-received. After 4, and 8 h of pure ageing, this number rapidly decreased to 534 MPa and 455 MPa, respectively.

As the stress applied reached 162.5 MPa under the creep ageing condition, the yield strength dropped dramatically to 493 MPa after 4 h and to 462 MPa after 8 h of creep ageing. The decline in yield strength was accelerated by an increase in the applied stress during creep ageing. As the creep ageing time increased to 8 h, the yield strengths of specimens subjected to various stresses levels steadily approached a similar value.

3.2. Microstructural Characterization

3.2.1. As-Received Condition

The results from creep ageing and tensile tests presented in the earlier section intriguingly revealed that, the strengths of this alloy rely on not only the ageing time but also the magnitude of stress applied in creep ageing. To clarify this phenomenon and to elucidate the relationship between yield strength, applied stress, and creep ageing time, a further TEM microstructural analysis was performed.

The grain shapes and overall morphologies of the as-received material were first characterized. Figure 4 presents a montage of TEM micrographs of the alloy demonstrating elongated and partially recrystallized grain structures. These structures resulted from the prior processing of heat treatment and hot rolling. The microstructures are actually composed of very fine newly developed aluminum grains. The fine-grained matrix of the alloy with a face centered cubic (fcc) lattice features ${\eta }^{\prime }$ precipitates and GP zones. Some of the ${\eta }^{\prime }$ precipitates come from the separated nucleation as the GP zones dissolved [9]. At this stage, these ${\eta }^{\prime }$ precipitates and GP zones are the main phases that contribute to the material’s strength [31]. Other intermetallics and larger precipitates, such as MgZn₂ ($\eta$ phase), Mg₂Si, Al₂CuMg, and Al₇Cu₂Fe, have also been reportedly found in this alloy system [32,33].

Figure 5 shows more detailed TEM micrographs and analysis for the as-received alloy. It can be observed in Figure 5a that ${\eta }^{\prime }$ and $\eta$ precipitates are dispersed throughout the aluminum matrix, and very few dislocations can be found at the as-received condition. The size distribution of these precipitates is illustrated in Figure 5b. Figure 5c is a high-resolution transmission electron microscopy (HRTEM) image revealing the GP zones, which are characterized by the multi-atomic layers structure (marked with 1 to 4). Figure 5d illustrates the respective diffractograms by fast Fourier transform (FFT) of those GP zones.

In addition to the primary diffraction patterns from the FCC Al matrix, distinctive diffuse streak patterns can be also detected. Those streak patterns represent the atomic layer structure of the GP zone. Notably, the direction of those layer-structured GP zones is perpendicular to the streak patterns; i.e., it is the reciprocal lattice signal of a typical layer structure. Figure 6a is an HR-TEM image of an ${\eta }^{\prime }$ precipitate, and the corresponding FFT diffractogram is illustrated in Figure 6b, which was acquired from the moiré fringes found in the matrix and ${\eta }^{\prime }$ precipitate. Figure 6c is the schematic diffraction patterns for ${\eta }^{\prime }$ precipitates and aluminum matrix. EDX data presented in Figure 6d suggest that the ${\eta }^{\prime } \mathrm{precipitate}$ is rich in Al, Cu, Mg, and Zn elements.

3.2.2. Precipitation Behavior of Stress Ageing

To evaluate the impact of ageing duration and stress applied on the precipitation behavior of GP zones and ${\eta }^{\prime }$ precipitates, the specimens subjected to pure/creep ageing treatment were studied through TEM. In both ageing processes, GP zones easily converted into ${\eta }^{\prime }$ precipitates as a result of ageing. This conversion reduced the GP zones’ stress field energy, which is caused by the Zn and Mg atoms’ inherent reordering. For the ${\eta }^{\prime }$ precipitates, there are four distinct orientation correlations with regard to the aluminum matrix: ${(0 0 0 1)}_{{\eta }^{\prime }}$//${ (1 1 1)}_{\mathrm{Al}}$ and ${[1 \overline{1} 0 0]}_{{\eta }^{\prime }}$//${ [0 \overline{1} 1]}_{\mathrm{Al}}$; ${(0 0 0 1)}_{{\eta }^{\prime }}$//${ (\overline{1} 1 1)}_{\mathrm{Al}}$ and ${[1 \overline{1} 0 0]}_{{\eta }^{\prime }}$//${ [0 \overline{1} 1]}_{\mathrm{Al}}$; ${(0 0 0 1)}_{{\eta }^{\prime }}$//${ (1 \overline{1} 1)}_{\mathrm{Al}}$ and ${[1 \overline{1} 0 0]}_{{\eta }^{\prime }}$//${[0 1 1]}_{\mathrm{Al}}$; ${(0 0 0 1)}_{{\eta }^{\prime }}$//${ (1 1 \overline{1})}_{\mathrm{Al}}$ and ${[1 \overline{1} 0 0]}_{{\eta }^{\prime }}$//${[0 1 1]}_{\mathrm{Al}}$. Therefore, through the lens of TEM along the zone axis of ${[1 \overline{1} 0 0]}_{{\eta }^{\prime }}$//${ [0 \overline{1} 1]}_{\mathrm{Al}}$, two variants of ${\eta }^{\prime }$ precipitates appear to be rod-shaped, both of which are in an edge-on formation on the ${ (\overline{1} 1 1)}_{\mathrm{Al}}$ and ${\left(111\right)}_{\mathrm{Al}}$ habit planes. Two other variants that are not in edge-on formation, on the other hand, appear to be disk-shaped. Figure 7 presents the microstructures of the pure/creep aged specimens via TEM along the ${ [0 \overline{1} 1]}_{\mathrm{Al}}$//${[1 \overline{1} 0 0]}_{{\eta }^{\prime }}$ zone axis. The precipitation behavior of the two aged specimens exhibited very different morphologies.

Figure 7b,d,f demonstrate the direction of tensile stress (162.5 MPa) applied at the creep ageing condition. It has been found that the stress applied had a significant effect on the creep aged specimens, causing the precipitates to align normally to the direction of stress in comparison with the pure aged specimens in Figure 7a,c,e. TEM analysis on the specimens after pure ageing for 4 h (Figure 7c) and creep ageing for 4 h (Figure 7d) revealed that the precipitates were somewhat bigger than the ones in the specimens after pure/creep ageing for 2 h, as shown in Figure 7a,b. Pure specimens aged for 8 h (Figure 7e), and creep aged for 8 h (Figure 7f) exhibited coarser phases, which were identified as $\eta$ precipitates. It is reasonable to deduce that ${\eta }^{\prime }$ to $\eta$ transformation took place throughout the duration of the pure/creep ageing processes. The creep ageing process, moreover, led to the elongation of the $\eta$ precipitates, as indicated by TEM examination along the zone axis of ${[0 \overline{1} 1]}_{\mathrm{Al}}$.

The TEM micrographs presented in Figure 8 depict a horizontal view of the precipitates lying on their habit planes after pure and creep ageing treatment, where the zone axis of the specimen for TEM analysis was adjusted to an orientation parallel to the <111>Al in order to determine the GP zones and ${\eta }^{\prime }$ precipitates’ habit planes. Different shapes of precipitates (spheroid, plate, rod-like shapes) can be observed. Through analyzing electron diffractograms, ${\eta }^{\prime }$ precipitates and GP zones can be recognized. The morphologies in Figure 8 also show no discernible differences in the form and allocation of ${\eta }^{\prime }$ precipitates and GP zones between the ageing process with or without the influence of stress.

To analyze the differences in growth of 𝜂′ precipitates and GP zones under the influence of stress in the ageing process, the sizes of 𝜂′ precipitates and GP zones has been measured. As mentioned previously, ${\eta }^{\prime }$ precipitates and GP zones are both the key contributors to strength in precipitation hardening, but bigger precipitates, such as $\eta$ phase, have a less impact on the alloy’s strength. Thus, the size of $\eta$ precipitates was not measured.

In HR-TEM evaluation, fast Fourier transform (FFT) diffractograms can be used to identify the lattice structures of the ${\eta }^{\prime }$/$\eta$ precipitates and GP zones. Because the habit planes tend to expand in tandem with the increasing size of ${\eta }^{\prime }$ precipitates and GP zones, the longitudinal lengths of the ${\eta }^{\prime }$ precipitates and GP zones can be used to illustrate their overall dimensions. The mean longitudinal lengths of the ${\eta }^{\prime }$ precipitates and GP zones for each parameter were determined by taking fifty measurements on each specimen from the TEM micrographs.

Figure 9 presents the measured longitudinal lengths of ${\eta }^{\prime }$ precipitates and GP zones as a function of ageing time, respectively, for the pure and creep aged specimens. For the pure ageing condition, the GP zones’ longitudinal length, as shown in Figure 9a, grew from 3.1 nm before the ageing process to 4.4 nm after 4 h of ageing, and then they subsequently ceased increasing in size. For the creep ageing condition, alternatively, the GP zones’ longitudinal length of developed more quickly, from 3.1 nm to around 6 nm, before the growth stopped, as shown in Figure 9a. The sizes of GP zones in pure aged specimens, in general, were shorter than those in creep aged specimens, with around 1.5 nm difference in longitudinal length, as illustrated in Figure 9a.

Under the creep ageing condition, the applied stress expanded the habit planes for GP zones to grow—the GP zones’ lengths of pure aged specimens were, as a result, shorter than those of creep aged specimens. The elongated GP zones were presumed to slowly convert into ${\eta }^{\prime }$ precipitates, and the latter ${\eta }^{\prime }$ precipitates were therefore longer than the former GP zones. As shown in Figure 9b, the ${\eta }^{\prime }$ precipitates were typically shorter for the pure aged specimens than the creep aged specimens. After 4 h of creep ageing, the longitudinal length of the ${\eta }^{\prime }$ precipitates reached 8.2 nm, while it was only 6.7 nm in specimens that were pure aged. After 8 h of ageing, the difference in length between the creep aged specimens and pure aged specimens fell lower, i.e., 8.3 nm (creep aged) versus 7.8 nm (pure aged).

3.2.3. TEM Characterization

Since ${\eta }^{\prime }$ precipitates and GP zones are very fine and dispersed, they provide the majority of strength for the material. Both of these phases are transitional phases with coherent or semi-coherent interfaces, which provide great lattice coherency with the matrix. In contrast, $\eta$ precipitates, the final equilibrium phase, with incoherent interfaces with the matrix, has no any lattice coherency with the matrix.

For the as-received material, the predominant phases are ${\eta }^{\prime }$ precipitates and GP zones. The minor distortion between the aluminum matrix and GP zones (${\eta }^{\prime }$ precipitates) produces coherency strengthening. Consequently, the material’s strength at this stage is high. When the material becomes over aged, these transition phases convert into larger precipitates that are incoherent with the matrix and fewer in number. The hardening mechanism of these precipitates is no longer contributed by coherency strain but instead by dispersion hardening—the mechanism wherein dislocations bypass precipitates by bending and bowing.

The transformation of GP zones into ${\eta }^{\prime }$ and $\eta$ precipitates is a continuous process [28]. Figure 10 is a high-resolution TEM (HR-TEM) image of a GP zone after 6 h of ageing. The GP zone vanishes abruptly after 5 s of examination. Such a phenomenon demonstrates that the GP zone is extremely unstable and tends to disintegrate throughout the ageing process so as to facilitate the transition of ${\eta }^{\prime }$ precipitates.

Figure 11 is an HR-TEM micrograph revealing the transformation of ${\eta }^{\prime }$ to $\eta$ precipitates, presumably MgZn₂. It is clearly shown that the ${\eta }^{\prime }$ and $\eta$ precipitates were linked with one another, as revealed in Figure 11a. Their respective diffractograms by (FFT) are shown in Figure 11b,c; Figure 11d,e illustrate the analyzed diffraction patterns for confirming the ${\eta }^{\prime }$ and $\eta$ precipitates.

The transformation of ${\eta }^{\prime }$ to $\eta$ precipitates is presumed to have been happening as the micrographs were being taken. The $\eta$ nuclei, which developed at the interface between the aluminum matrix and ${\eta }^{\prime }$ phase, slowly formed and took over the ${\eta }^{\prime }$ precipitate. In the end, the ${\eta }^{\prime }$ precipitate vanished, leaving an $\eta$ precipitate with ${\eta }^{\prime }$ morphology. It must be emphasized that the transformation of ${\eta }^{\prime }$ to $\eta$ includes the redistribution of alloy constituents; $\eta$ is presumably MgZn₂.

On the other hand, GP zones are also observable at every level of ageing condition. Therefore, when the alloy reaches the state of over ageing, it is plausible to presume that the remaining solute atoms in the aluminum matrix are still enough to form GP zones. As a result, coherency strain hardening can still contribute to the alloy’s strength. In the over ageing condition of aluminum alloys, the primary factor of strength softening is the coarsening of $\eta$ precipitates, since it decreases the number of GP zones and ${\eta }^{\prime }$ precipitates significantly. In the creep aged alloy, the coarsening effect of $\eta$ precipitates can be expedited by the applied stress.

For aluminum alloys, the mechanical properties are also affected by the precipitate free zones, which usually form during the pure/creep ageing process. Grain boundaries serve as effective vacancy sources and sinks that separate solute atoms, causing stable phases to emerge at the grain boundaries. Figure 12 compares the TEM images of precipitate free zones in the specimens after 8 h of pure and creep ageing. In both instances, the widths of the precipitate free zones were roughly identical. The observable impact of creep on precipitate free zones does not seem to be substantial.

4. Conclusions

This work has been done to characterize the strength, deformation, and microstructural development of the AA7050-T6 alloy throughout the creep ageing processes. Evaluations from mechanical testing demonstrated that the material exhibits typical over ageing behavior, as characterized by the dependency of applied stress and ageing time.

Tensile tests on pure aged specimens indicated that as ageing time rises, there is an increase in material elongation. This finding is consistent with the understanding that, during the forming process, ageing treatment can increase the ductility of the alloy [34]. The relationship between yield strength and creep ageing time suggested that the stress applied in the creep ageing process may reduce the alloy’s strength. After 8 h of ageing; however, the yield strength disparity of the specimens under various stress levels was marginalized.

Another focus of this study was the microstructural analysis of the AA7050-T6 alloy under different creep ageing process, with an emphasis on the development of GP zones and ${\eta }^{\prime }$/$\eta$ precipitates, the growth of which can be affected by the stress applied in creep ageing.

To address the issue, TEM analysis was performed. It has been discovered that the external stress can speed up the growth rate of GP zones and ${\eta }^{\prime }$ precipitates, which also accelerates the coarsening of $\eta$ precipitates. It was discovered that even after 8 h of creep ageing treatment, GP zones could still exist in the matrix, indicating that the nucleation of GP zones is a continuous process.

Precipitate free zones were observed in both pure and creep aged specimens, and the widths of the precipitate-free zones were roughly identical. This illustrated that the effect of creep ageing on precipitate free zones did not seem to be substantial. The effect on hardening by dislocation density evolution, however, could not be observed with confidence. Deeper research on the measurement of dislocation density is necessary. The main findings and contributions are listed as follows:

• Applying stress can speed up the growth rate of ${\eta }^{\prime }$ precipitates and GP zones, which also leads to faster coarsening of the η precipitates.
• The stress applied in creep has been found to reduce the overall strength of the AA7050-T6 alloy; the higher the applied stress, the lower the yield strength.
• TEM microscopy indicated that the drop in yield strength could be due to the faster coarsening rate of the η precipitates under higher applied stress in creep.
• GP zones are still observable when the alloy reaches the state of over ageing. It is therefore plausible to presume that the remaining solute atoms in the matrix are still enough to form GP zones.