Retrogression and Reaging Applied to Warm Forming of High-Strength Aluminum Alloy AA7075-T6 Sheet

Two AA7075-T6 sheet materials, referred to herein as materials A and B, were obtained from two different suppliers in the peak-aged condition. Both materials had an as-received sheet thickness of 2 mm. Elemental compositions were determined using X-ray fluorescence (XRF) analysis with a Horiba XGT-7200™ X-ray analytical microscope, and are shown in Table II together with specification limits of the AA7075 alloy.[38,39] Elemental compositions were determined after calibrating with a pure copper reference because an AA7075 reference standard was not available. This may explain why some elemental compositions for materials A and B appear outside of the designated composition limits for AA7075. The relative differences in elemental composition between the materials are less likely to be affected. Compared to material B, material A has lower concentrations of Zn, Mg, Cu, and Si, and therefore might be expected to exhibit a lower peak-aged hardness.

Specimens were prepared for polarized light optical microscopy by mechanical grinding, polishing, and then electrolytic etching (20 V) with Barker’s reagent (1.8 pct fluoboric acid in 98.2 pct distilled water)[40] to produce color contrast between grains when viewed under polarized light. Figure 2 presents the microstructures of materials A and B in both the as-received T6 and the solutionized conditions. The solutionizing treatment involved annealing at 480 °C for 2 hours, followed by water quenching. Grain sizes were measured along the short-transverse direction (STD) and the rolling direction (RD) in the as-received and the solutionized conditions using the lineal-intercept method described in ASTM E112.[41] Average grain sizes are provided with standard deviations of the measurements in Table III. Both materials exhibit the elongated grain structures expected of rolled sheet material, as demonstrated in Figure 2. Grain size after solutionizing changed little for either material, as demonstrated in Table III; also compare Figures 2(a) and (b) and compare Figures 2(c) and (d). Some difference between grain sizes at the sheet surface and sheet center is noted for each material, with finer grains near the surface, and this difference is most pronounced in material B.

Heat treatment experiments were conducted on rectangular specimens with dimensions of 2 × 10 × 30 mm, excised from materials A and B in the as-received T6 condition. These specimens were used to evaluate the effects of several different heat treatments involving retrogression. Retrogression treatments were performed in a molten salt bath at temperatures of 200 °C, 225 °C, 250 °C, 300 °C, and 350 °C for times ranging from 10 to 150 seconds. The salt bath controlled specimen temperature to within approximately ± 3 °C of that desired. Following retrogression, specimens were held at room temperature (approximately 21 °C) for 1 week prior to reaging. This hold period was implemented to reflect conditions where formed components are typically trimmed, cleaned, and applied to a vehicle body-in-white over the course of several days before paint baking. Specimens were then either retained in the retrogressed condition or subjected to further heat treatment in a resistance tube furnace to evaluate reaging. The tube furnace controlled temperature to within approximately ± 3 °C of that desired. Three heat treatments were used to study reaging behavior. The first employed the typical reaging treatment of 120 °C for 24 hours.[19] The second was a simulated paint-bake treatment of 180 °C for 30 minutes, which reflects one industry practice for paint baking.[29] The third employed the standard reaging treatment followed by the simulated paint-bake treatment. Specimens were thus produced in the following four conditions: retrogressed only (R); retrogressed and reaged (RA); retrogressed and paint baked (RP); or retrogressed, reaged, and paint baked (RAP). These four heat treatments are summarized in Table IV with the labels used to reference each. The results of these experiments were used to determine the retrogression conditions from which peak-aged strength can be recovered during subsequent reaging.

At each heat-treating step, temperature was monitored with one or more K-type thermocouples. Specimens were quenched to room temperature (approximately 21 °C) in water immediately following each heat treatment. The final quenching step was applied for consistency among experiments but is unlikely to be needed in the practical application of retrogression forming. The material response to heat treatment was measured using Vickers hardness tests (HV10) at room temperature; hardness can be used as a reasonable surrogate for tensile strength in aluminum alloys.[42] Hardness tests were conducted using a load of 10 kg and a load hold time of 15 seconds. Each hardness value reported is an average of three to five individual measurements, and error bars used in the figures represent an average of the standard deviations across all hardness measurements. Specimens in all heat-treated conditions were stored at room temperature (approximately 21 °C), and hardness tests were intermittently performed over a period of three years to measure the effects of natural aging. These tests were conducted because AA7075 is susceptible to natural aging at room temperature in some temper conditions.[19,37]

Tensile tests were conducted at room and elevated temperatures using computer-controlled electromechanical and servo-hydraulic mechanical testing machines. Tests were performed in displacement control mode with a computer adjusting displacement rate based on cross-head or piston motion to maintain an approximately constant true-strain rate during testing. Rates for this control were calculated assuming uniform specimen deformation and constant specimen volume. Strain is reported from an extensometer for room-temperature tests and from piston motion for elevated-temperature tests; an elevated-temperature extensometer was not available for this study. Room-temperature (approximately 21 °C) tests were conducted with the electromechanical testing machine. Elevated-temperature (200 °C and greater) tests were conducted with the servo-hydraulic mechanical testing machine; specimen temperature during these tests was controlled using a furnace coupled to the mechanical testing machine. Two different furnaces were used, one for testing at low temperature (200 °C) and the other for testing at high temperatures (≥ 300 °C). Dog bone-shaped tensile specimens were designed according to ASTM standards E8, E21, and E2448.[43‐45] Specimens were water-jet cut from the as-received materials. Different specimen geometries were required by the furnace and gripping fixture combinations available for tensile testing at different temperatures and are described in Table V. The rolling direction was oriented either parallel or perpendicular to the tensile direction and is noted for each individual test result presented. All tensile specimens retained the as-received sheet thickness of 2 mm for materials A and B prior to testing.

Tensile tests were conducted at 200 °C in order to evaluate ductility under warm-forming conditions. A convective box furnace maintained this test temperature during testing. It was determined from the heat treatment experiments that exposure to the testing temperature at 200 °C should not exceed 150 seconds to guarantee the recovery of peak-aged strength during subsequent reaging. Specimen exposure time to this testing temperature was limited by using the following rapid-heating technique. Specimens were first preheated to 110 °C in less than 20 minutes in a resistance tube furnace. Data available in the literature demonstrate that the microstructure of AA7075-T6 is not significantly altered by annealing at temperatures below 120 °C for this short time.[19] This was confirmed by experiments that measured hardness before and after such low-temperature soaks. Specimens were then placed between two copper plates held at the desired test temperature, 200 °C, within the convective box furnace for 60 seconds to reach the testing temperature. Specimens were then transferred to tensile grips in the convective furnace within approximately 30 seconds. Tensile testing began within a maximum of 60 seconds and was completed within an additional 20 seconds. Tested specimens were then extracted and quenched to room temperature (approximately 21 °C) in water within 20 to 40 seconds. The longest time from specimen insertion into the furnace at 200 °C through specimen quenching is 210 seconds. Because the time to reach 200 °C, estimated as 60 seconds, is included in this time, the experiments met the goal of limiting total specimen time at temperature to approximately 150 seconds. Specimen temperature during testing was controlled to within ± 3 °C of the desired test temperature. The temperature of the specimen was monitored during testing using two K-type thermocouples in contact with the surface of the specimen. Specimens were tested in tension at approximately constant true-strain rates of 1 × 10⁻¹, 5 × 10⁻², and 1 × 10⁻² s⁻¹.

Tensile tests at 300 °C, 460 °C, 480 °C, 500 °C, and 520 °C were conducted to evaluate ductility under hot-forming conditions, for which partial or full solutionization of AA7075 is expected.[19] Tensile specimens were solutionized at 480 °C for 2 hours immediately prior to testing to prevent the precipitate dissolution process from influencing deformation behavior. Test temperature was controlled using a split-tube three-zone resistance furnace. The furnace was preheated to the desired temperature prior to testing. The specimen was then transferred to the furnace and brought to the desired temperature; temperature during testing was controlled to within ± 3 °C. The temperature of the specimen was monitored during testing with two K-type thermocouples in contact with the surface of the specimen. Tests were conducted at approximately constant true-strain rates from 3 × 10⁻⁴ to 3 × 10⁻¹ s⁻¹. Specimens were immediately quenched to room temperature (approximately 21 °C) in water upon the completion of tensile deformation.

Vickers hardness measurements after retrogressing at 200 °C, 225 °C, 250 °C, 300 °C, and 350 °C are presented in Figure 3(a) for material A and Figure 3(b) for material B as functions of retrogression time. The retrogression behaviors in Figure 3 can be grouped into two categories based upon the potential to recover hardness to peak-aged values through reaging. Either the hardness lost during retrogression is recoverable to nearly peak-aged values upon reaging or the hardness lost during retrogression cannot be recovered. Retrogression treatments conducted at 200 °C, 225 °C, and 250 °C demonstrate a local minimum in hardness, as shown in the schematic of Figure 1(a). Note that the local minimum in hardness at 200 °C is not readily determined from Figure 3 and might occur at retrogression times greater than those considered in this study. Within the vicinity of this local hardness minimum, there is a potential to fully recover hardness through subsequent reaging. The extant literature suggests that after retrogression of AA7075-T6 material at these temperatures and times, peak-aged hardness and strength can be recovered through an appropriate reaging treatment.[30‐33] As retrogression temperature increases from 200 °C to 250 °C, the severity of the initial reduction in hardness increases, and the time required to reach the local minimum hardness value, \( t_{\text{r}}^{*} \), decreases. For both materials A and B, \( t_{\text{r}}^{*} \) occurred at 10, 30, and 150 seconds or longer for retrogression at 250 °C, 225 °C, and 200 °C, respectively. The reduction in hardness at \( t_{\text{r}}^{*} \) increased from 12 to 20 pct for material A and from 9 to 15 pct for material B as retrogression temperature increased from 200 °C to 250 °C. It is noted from Figure 3 that material A typically exhibits a 3 to 5 pct greater reduction in hardness from retrogression than does material B for a given retrogression temperature and time. The reason for this discrepancy is not explored here, but potential causes are later discussed. The absence of local hardness minima in Figure 3 at 300 °C and 350 °C demonstrates that the hardness decreases in both materials A and B are unrecoverable after retrogression at these temperatures. Although shorter retrogression times might allow some hardness recovery during reaging, such short times are not of practical importance to retrogression forming. For this reason, retrogression at temperatures of 300 °C and higher is not further considered.

To better represent and understand the combined effects of temperature and time on hardness reduction during retrogression heat treatments, a method for constructing a single-master curve from hardness data is proposed. This method uses the Arrhenius relationship[46] to combine time and temperature into a temperature-compensated retrogression time, τ, against which hardness can be plotted to produce a single curve among data from the different retrogression temperatures investigated. This technique assumes that the dissolution of precipitates during retrogression depends primarily on a single diffusion process, likely either Zn or Mg in Al, governed by a single activation energy. The temperature-compensated retrogression time is defined aswhere Q is the activation energy for dissolution of η′ and η precipitates (J/mol); R is the universal gas constant (J/mol-K); T_R is the retrogression temperature (K); and t is the retrogression time (s). The value for Q is taken as 95 kJ/mol, based upon data reported by Balderach et al. for precipitate dissolution in Al-Zn-Mg alloys, obtained using differential scanning calorimetry.[37]

Vickers hardness after retrogression at temperatures from 200 °C to 350 °C is presented in Figure 4 as a function of the temperature-compensated retrogression time, τ, for materials A and B. Hardness data from Park[31] and from Kanno et al.[32] at the local hardness minimum, i.e., \( t_{\text{r}}^{*} \) in Figure 1(a), are included for comparison. All data from materials A and B fall reasonably well onto the single dashed curve drawn in Figure 4. The most important feature of Figure 4 is that a single value of temperature-compensated retrogression time, \( \tau_{\text{r}}^{*} \), describes the location of the local hardness minimum for all these data, including data from the literature, across a range of retrogression temperatures. The bracket shown above the \( \tau_{\text{r}}^{*} \) label in Figure 4 indicates the approximate uncertainty in the local hardness minimum for the data presented. The average value for \( \tau_{\text{r}}^{*} \) in Figure 4 is 3.6×10⁻⁹ seconds, and its minimum and maximum values are 2.3×10⁻⁹ and 5.4×10⁻⁹ seconds, respectively. Note that this value for the temperature-compensated retrogression time can be used with Eq. [1] to describe the retrogression time necessary to reach the local hardness minimum for any given retrogression temperature. For example, the \( t_{\text{r}}^{*} \) value for retrogression at 200 °C is predicted as 110 seconds with minimum and maximum values of 70 and 165 seconds, respectively, using Figure 4. For both materials A and B, Figure 3 demonstrates for retrogression at 200 °C that \( t_{\text{r}}^{*} \) occurred at 150 seconds, which falls within the predicted range of values for \( t_{\text{r}}^{*} \) using Figure 4. The hardness values at \( \tau_{\text{r}}^{*} \) vary between 135 and 181 HV10 with material and with retrogression temperature, as might be expected. Note that in Figure 3, for example, the local minimum hardness value is different at each retrogression temperature. Despite the differences in the local hardness minimum between the materials and retrogression temperatures, the local minimum reliably occurs at a consistent value of temperature-compensated retrogression time, \( \tau_{\text{r}}^{*} \), of approximately 3.6 × 10⁻⁹ seconds, and this provides a useful means of predicting optimal retrogression treatments across a range of retrogression times and temperatures.

The viability of recovering hardness through a simulated paint bake after retrogression was investigated. Specimens were naturally aged for one week following retrogression, prior to application of a simulated paint-bake treatment of 180 °C for 30 minutes. Reaging experiments were conducted with the simulated paint bake on specimens retrogressed at 200 °C, 225 °C, and 250 °C. Figure 5 presents Vickers hardness measurements after paint-bake reaging, which demonstrates that peak-aged hardness was not fully recovered in either material A or B. Hardness, however, did increase by 3 to 10 pct from the retrogressed condition for specimens of material A that were retrogressed at 200 °C and 225 °C and for specimens of material B that were retrogressed at 200 °C. The hardnesses of these specimens after paint-bake reaging were within 5 to 10 pct of the peak-aged (T6) condition. Material B retrogressed at 225 °C did not display any appreciable recovery of hardness through reaging with the simulated paint bake. Both materials retrogressed at 250 °C recovered significant hardness through the simulated paint bake compared to their retrogressed conditions, but the final hardnesses never approached peak-aged values. Retrogression treatments conducted at 250 °C were not investigated further because hardness was not sufficiently recovered at this temperature to make it of interest.

The hardnesses of materials A and B after retrogression and reaging with the simulated paint bake are compared with hardness values produced by other reaging procedures shown in Figure 6. Vickers hardness measurements from the following conditions are presented as a function of retrogression time for retrogression temperatures of 200 °C and 225 °C: retrogressed only (R); retrogressed and paint baked (RP); retrogressed and reaged (RA); and retrogressed, reaged, and then paint baked (RAP). The effectiveness of a reaging procedure can be evaluated by comparing the hardness produced to the peak-aged (T6) hardness and to the hardness after retrogression.

Figure 6(a) presents the hardness of material A after retrogression at 200 °C (curve R) and the three different reaging procedures investigated in this study (curves RA, RP, and RAP). Note that T6 hardness, measured from material A in the as-received, peak-aged condition, is marked in Figure 6(a). Curve RA demonstrates that the standard reaging procedure increases hardness by 10 to 20 pct from the retrogressed condition (curve R) and recovers the full T6 hardness of material A. Curve RP illustrates that hardness in material A increased by 3 to 10 pct from the retrogressed condition after reaging with the simulated paint bake, with the exception of the specimen retrogressed for only 10 seconds. Hardness remains between 5 and 10 pct below peak-aged hardness for the RP specimens. The RAP reaging procedure produced hardness increases from the retrogressed condition that are similar to those produced by the RP reaging procedure, as demonstrated by Figure 6(a). Figure 6(b) presents the hardness of material B after retrogression at 200 °C and after subsequent reaging. T6 hardness is marked in Figure 6(b) as measured from material B in the as-received, peak-aged condition. The standard reaging procedure, curve RA, recovered hardness to within 3 pct of the peak-aged value, with the exception of the specimen retrogressed for 150 seconds at 200 °C. The paint-bake reaging, curve RP, increased hardness by 5 to 8 pct over the retrogressed condition for retrogression times of greater than 10 seconds. However, the hardnesses of the RP specimen remained approximately 5 to 10 pct less than the peak-aged (T6) condition. The RAP procedure did not increase hardness significantly over the retrogressed condition.

Figure 6(c) presents the hardness of material A after retrogression at 225 °C and after subsequent reaging. The standard reaging procedure recovered hardness in material A to nearly peak-aged values after retrogression at 225 °C for 10 and 30 seconds, but not for 90 and 150 seconds, as illustrated by curve RA in Figure 6(c). The RP and RAP procedures generally increased hardness by 5 to 12 pct over the retrogressed condition in material A after retrogression at 225 °C. Figure 6(d) presents the hardness of material B after retrogression at 225 °C and after reaging. Similar to material A after retrogression at 225 °C, Figure 6(d) demonstrates that the standard reaging procedure, curve RA, recovered hardness to nearly peak-aged values for retrogression times of 10 and 30 seconds, but not for 90 and 150 seconds in material B. The RP and RAP procedures did not improve hardness in material B over the retrogressed conditions, as demonstrated in Figure 6(d).

The effect of natural aging on the hardness of materials A and B after retrogression at 200 °C and subsequent reaging with the simulated paint-bake cycle was monitored for over three years. Vickers hardness measurements are shown as a function of the logarithm of natural aging time for four different retrogression times in Figure 7. The hardnesses of materials A and B in the retrogressed condition, shown in Figures 7(a) and (b), respectively, changed with natural aging time. There is a slight increase in hardness over time, most pronounced in the materials retrogressed for 10 seconds. Figures 7(c) and (d) present the hardnesses of materials A and B, respectively, after retrogression at 200 °C and reaging with a simulated paint bake at 180 °C for 30 minutes. Natural aging did not significantly affect either material in the retrogressed and paint-baked (RP) condition. The natural aging of materials A and B retrogressed at 225 °C was also monitored. Because results from the materials retrogressed at 225 °C mimic the materials retrogressed at 200 °C, those data are not presented. The hardnesses of the materials retrogressed at 225 °C also increased with natural aging time, but the RP condition is similarly stable to the RP materials retrogressed at 200 °C.

Uniaxial tensile tests at 200 °C were conducted to examine the tensile ductilities possible from AA7075-T6 under deformation conditions suitable for retrogression from which peak-aged strength can be recovered during subsequent reaging (see Figure 6(a) for material A and Figure 6(b) for material B). Deformation behavior was examined at 200 °C because the retrogression time to the local minimum in hardness, \( t_{\text{r}}^{*} \), occurred at 150 seconds or longer for retrogression at 200 °C in both materials studied (see Figure 3), which maximizes the time available for retrogression-forming operations compared to the other retrogression temperatures investigated. Additionally, deforming at the lowest retrogression temperature studied, 200 °C, suggests less energy expenditure without sacrificing significant gains in ductility compared to higher temperatures. Tests were conducted at approximately constant true-strain rates of 1 × 10⁻¹, 5 × 10⁻², and 1 × 10⁻² s⁻¹. Rapid-heating procedures, described in the materials and experimental methods section, limited the specimen’s time at 200 °C to a maximum of approximately 150 seconds, which is within the time \( t_{\text{r}}^{*} \) indicated by the data shown in Figure 3. Only data from tensile tests on material A are presented in this study; the behaviors observed for material B are not appreciably different from those of material A. Figure 8 presents plots of true stress against true strain for tests conducted on material A at 200 °C. Data from specimens with the rolling direction oriented parallel and perpendicular to the tensile direction are presented in Figures 8(a) and (b), respectively. Neither flow stress nor rupture strain varied markedly with strain rate or the orientation of the rolling direction relative to the tensile direction. Flow stress remained between 364 and 396 MPa for all tests, and rupture, defined as the complete failure of the specimen into two distinct pieces, occurred at an average true strain of 0.2 and did not vary by more than 0.03 between the three strain rates tested. Failure was by flow localization, i.e., necking, quickly followed by rupture. The downward trend in flow stress near rupture is an artifact from the calculation of true stress that results from necking behavior; local geometry in the necked region could not be measured during the test. The benefit of presenting these data as true values is in demonstrating the absence of any significant strain hardening prior to neck development.

Tensile data from the present study are presented in Figure 9 on a plot of true stress against true strain for material A at temperatures of 21 °C, 200 °C, and 300 °C. Data from Wang et al.[16] and Sotirov et al.[17] for AA7075-T6 tested at temperatures from 180 °C up to 260 °C are also shown in Figure 9. These data together demonstrate the tensile ductilities possible from AA7075-T6 at temperatures up to 300 °C for strain rates between 5×10⁻² and 1 s⁻¹. Table VI summarizes the results from tensile tests conducted in this study. Elongation-to-rupture (e_r (pct) in Table VI) in material A at 200 °C, a temperature suitable for retrogression, is double that achieved at room temperature. Data from the literature likewise demonstrate that ductility in AA7075-T6 is increased by a factor of 2, approximately, at temperatures of 180 °C to 230 °C. It is noteworthy that increasing temperature from 230 °C up to 300 °C does not significantly improve tensile ductility. Even when comparing data from within one set of tests, the ductility at 300 °C is only slightly greater than that at 200 °C, for a strain rate of 5 × 10⁻² s⁻¹.

Tensile behavior of material A between 460 °C and 520 °C is summarized in Figure 10 and Table VI. Ductility between 460 °C and 520 °C is much greater than that at 200 °C. Flow stress data at a true strain of 0.1 are presented in Figure 10 as a plot of the logarithm of true-strain rate against the logarithm of true stress normalized by the dynamic, unrelaxed, temperature-dependent elastic modulus, i.e., σ/E. Elastic modulus data from Köster for aluminum were used for the stress normalization.[52] The stress exponent of 5, measured from the slope of the data in Figure 10, is that expected of dislocation-climb-controlled creep deformation.[53] Data from different temperatures, taken at constant values of σ/E, were used to calculate the activation energy for creep using the following relationship,[53]where Q_C is the activation energy for creep deformation; R is the universal gas constant; \( \dot{\varepsilon } \) is the true-strain rate; and T is absolute temperature. The activation energy for creep, is calculated to be within the range of 166 to 170 kJ/mol. for σ/E values of 4.9 × 10⁻⁴, 5.2 × 10⁻⁴, and 5.5 × 10⁻⁴.

The retrogression behaviors of materials A and B, as presented in Figure 3, can be grouped into two categories based on whether hardness is (1) recoverable or (2) unrecoverable by reaging after retrogression. Among retrogression treatments conducted at 200 °C, 225 °C, and 250 °C, each demonstrates a local minimum in hardness and hardness lost during retrogression is 80 to 100 pct recoverable from the vicinity of these local minima by the standard reaging treatment (24 hours at 120 °C). This is illustrated by comparing hardness after retrogression in Figure 3 to hardness after reaging with the simulated paint bake in Figure 5 and by comparing the RA and RP curves to the R curves in Figures 6(a) through (d). Data in the literature demonstrate through TEM and SAED that retrogression at these temperatures either partially dissolves the fine (5 to 10 nm) η′ and η precipitates characteristic of peak-aged AA7075-T6 or transforms and coarsens these into larger (> 20 nm) η precipitates, particularly at grain boundaries.[31‐36]

Because the recovery of hardness by reaging is possible among the conditions studied here only after retrogression temperatures less than or equal to 250 °C for times limited by the temperature-dependent values of \( t_{\text{r}}^{\hbox{max} } \), the mechanism of reaging likely depends on limiting species diffusion during retrogression. It is likely necessary to limit the diffusion distances of Zn and/or Mg out of and away from η′ and η precipitates during their dissolution and limit diffusion of Zn and/or Mg into larger η precipitates during coarsening. Retaining high local concentrations of Zn and/or Mg near partially dissolved precipitates will support the reformation of these precipitates during reaging. If full dissolution of a precipitate occurs, a high local concentration of Zn and/or Mg can accelerate nucleation and growth of a new precipitate at the same location. As retrogression time and temperature increase, the increased homogenization of these species because of diffusion within the Al lattice and the possibility of tying up solute elements in significantly coarsened precipitates is likely to prevent effective reaging, and hardness is then unrecoverable in the practical sense. Retrogression at temperatures of 300 °C and higher does not demonstrate a local hardness minimum within the scope of experimental data shown in Figure 3, as any local minimum in hardness occurs at times too short to be readily measured. At these retrogression temperatures, complete dissolution of fine η′ and η precipitates and a fairly homogeneous distribution of the atomic species Mg and Zn likely occur after relatively short times. Large η precipitates and even larger alternative precipitation products, such as the T-phase ((AlZn)₄₉Mg₃₂), are then likely.[34,42] Consequently, strength cannot be recovered after such retrogression treatments by simple reaging because solute elements are consumed by the large precipitation products preventing the formation of fine η′ and η precipitates. A full solutionizing treatment and the usual artificial aging practice would be required after retrogression at temperatures of 300 °C and higher.

Figures 3 and 6 demonstrate that material A exhibits a larger drop in hardness than material B under identical retrogression conditions by approximately 3 to 5 pct in hardness loss. The reason for this was not explored in this study, but it is likely a result of composition differences between these materials. Table II demonstrates that material B contains more Zn, Mg, Si, and Cu than material A. These differences are expected to affect retrogression behavior, but further investigation is needed to confirm this expectation. Regardless, controlling composition is necessary to achieve predictable retrogression behavior and possibly adjust it for a desired performance.

Figure 4 proposes the use of temperature-compensated time, τ, as defined by Eq. [1] to predict retrogression times that produce a local hardness minimum after retrogression at given retrogression temperatures. In this case, the local hardness minimum occurs at an approximately constant value of 3.6 × 10⁻⁹ seconds for temperature-compensated time, \( \tau_{\text{r}}^{*} \), across a range of retrogression temperatures. The activation energy used in calculating this reduced time for Figure 4 is 95 kJ/mol, which is that determined in a previous study[37] for precipitate dissolution in similar alloys at temperatures similar to those used for retrogression in this study. It is hypothesized that this value can be interpreted as approximately the activation energy for Zn diffusion in Al, as in the previous investigation,[37] because the rate of Zn diffusion is faster than Mg in Al.[47‐51] The rapidity of η′ and η precipitate dissolution, from the perspective of strength loss, is likely controlled by the fastest diffusing species among those which comprise these precipitates, Zn in this case. The activation energy for the diffusion of Zn in Al, Al-Zn-Mg, and Al-Zn-Cu materials obtained from traditional tracer diffusion methods is reported in the literature as approximately 120 kJ/mol,[47‐51] but these data are for higher temperatures than those used for retrogression. The extant literature on diffusion provides data only for the diffusion of Mg and Zn in Al at temperatures higher than those used for recoverable retrogression (200 °C to 250 °C) in this study,[47‐51] which leaves a need for additional diffusion data to test the interpretation proposed here. The value for activation energy could also be determined empirically by forcing the data in Figure 4 to collapse onto a single curve. Regardless of these details, Eq. [1] provides a tool of practical utility for predicting appropriate retrogression times and temperatures for a potential retrogression-forming process using AA7075-T6.

Reaging AA7075 after a retrogression treatment is necessary to recover strength lost during that treatment and is ideally performed for 24 hours at 120 °C, the standard artificial aging temperature for AA7075.[19] Figure 6 demonstrates that peak-aged (T6) hardness was fully recovered (within 5 pct of the measured T6 value) in materials A and B after retrogression at 200 °C and 225 °C by reaging at 120 °C for 24 hours, if retrogression time did not pass \( t_{\text{r}}^{\hbox{max} } \), as shown in Figure 1(a). For materials A and B, \( t_{\text{r}}^{\hbox{max} } \) is likely beyond 150 seconds for retrogression at 200 °C and is between 30 and 90 seconds for retrogression at 225 °C (see Figure 6). The literature indicates that this reaging procedure produces a microstructure containing numerous fine η′ precipitates and some η precipitates distributed throughout the aluminum matrix, from which peak-aged strength is derived. Park and Ardell demonstrated that the numbers of η′ and η precipitates present in retrogressed-and-reaged (RRA) AA7075 and in AA7075-T6 are approximately equal.[34]

Reaging was also performed with a simulated paint-bake cycle of 180 °C for 30 minutes. If strength lost during retrogression forming is recovered through the paint bake, no additional heat treatments or changes to current paint-baking practices would be required to provide nearly peak-aged strength in final components. Figure 6 demonstrates that hardness after retrogression at 200 °C and 225 °C for times near \( t_{\text{r}}^{*} \) was recovered to within 5 to 10 pct of the peak-aged value by reaging with the simulated paint bake. The simulated paint-bake temperature of 180 °C, which reflects one current industry practice, exceeds the temperatures recommended for standard peak-aging (120 °C for 24 hours) and the second step of standard over-aging (163 °C for 27 hours) treatments.[19,42] Reaging with the simulated paint bake likely produces fewer fine η′ and η precipitates compared to reaging at 120 °C for 24 hours, and the average precipitate size is thereby expected to be larger. However, hardness after reaging with the simulated paint bake is typically greater than that of over-aged AA7075-T7, nominally 155 HV10.[19] We surmise, based upon characterization of AA7075 microstructures in the peak-aged, over-aged, and retrogressed conditions presented in the literature,[31‐36] that the short duration of the simulated paint bake avoids over-aging while still producing fine precipitate reformation. The short duration of the simulated paint bake may largely avoid (1) the coarsening of large η precipitates, (2) the transformation of η′ to η, and (3) the further dissolution of fine η′ and η precipitates. Further characterization of the microstructures produced would be necessary to more fully understand these effects.

A two-step reaging procedure after retrogression (RAP—see Table IV) was also investigated for materials A and B. For the two-step procedure, the material was first retrogressed at 200 °C or 225 °C, reaged at 120 °C for 24 hours, and then subjected to the simulated paint bake. This two-step procedure suggests one possible manufacturing route in which material is subjected to a standard reaging treatment following retrogression forming and is subsequently processed through a paint-bake cycle. As demonstrated by the data in Figure 6, the RAP process provides no substantive advantage in final hardness, and through inference, final peak strength, over applying only a simulated paint bake after retrogression (RP—see Table IV). The RAP process generally produces slightly less hardness recovery than does the RP process alone. These data suggest that a standard reaging treatment following retrogression, or retrogression forming, should be avoided if a paint bake is to follow.

After retrogression at 225 °C, the simulated paint bake restored significant hardness in material A but not in material B, as illustrated by Figures 6(c) and (d). Figures 6(a) and (b) illustrate that a standard reaging treatment after retrogression at 200 °C in material A recovers approximately 3 pct more hardness than for material B. These differences in reaging behavior can likely be attributed to compositional differences between the materials.

The conditions appropriate to perform retrogression concurrently with warm-forming and to recover nearly peak-aged strength after a simulated paint bake were determined. Hardness can be recovered to within 5 to 10 pct of the peak-aged (T6) condition in both materials A and B by a simulated paint bake following retrogression treatments conducted at 200 °C for up to 150 seconds. This offers a reasonable window for simultaneous warm forming at 200 °C, which is necessary to enable a practical retrogression-forming technology. While not addressed in this study, the effect of strain on RRA behavior in AA7075 sheet material should be explored prior to the implementation of RRA-forming.

Tensile elongation and reduction-in-area at rupture are two of the many measures useful for characterizing sheet material formability. Figure 11(a) presents tensile elongation and reduction-in-area at rupture for material A as a function of test temperature. These data are from uniaxial tensile tests conducted at temperatures from 21 °C to 520 °C. Large differences between tensile elongation and reduction-in-area measurements at a given temperature suggest significant flow localization, i.e., necking, prior to rupture. Maximum true flow stress is presented in Figure 11(b) as a function of test temperature; this value represents a nearly steady-state flow stress for tests at 200 °C and higher temperatures, as is evident from Figures 8 and 9. Figure 12 displays specimens before and after uniaxial tensile tests conducted on material A at 21 °C, 200 °C, 300 °C, 460 °C, and 520 °C. Data from these tests are summarized in Table VI.

The tensile elongation of AA7075-T6, as represented by material A, at room temperature is limited to approximately 10 pct (see Figure 9). Modest flow localization is evident for this specimen in Figure 12 and from the discrepancy between reduction-in-area and tensile elongation at rupture in Figure 11(a). Figure 11(b) demonstrates a maximum true flow stress of 641 MPa at room temperature. This high strength is useful in service but makes room-temperature forming challenging because it produces very high stamping forces and significant spring back. These issues and limited ductility at room temperature severely restrict the complexity of shapes that can be stamped from AA7075-T6 at room temperature. Tensile elongation at room temperature is improved to 17 pct when the material is in the O temper, the softest condition for AA7075.[19] However, the O temper introduces significant problems with temper stability during storage and requires aging at 120 °C for 24 hours after forming to achieve the T6 temper, adding additional complexity and cost to the manufacturing process.

Data from the warm tensile tests, defined as temperatures near 200 °C in this study, suggest significant improvements in ductility and reductions in strength that benefit the retrogression-forming process. Figure 11(a) demonstrates that the ductility of AA7075-T6 is improved twofold, to 20 pct tensile elongation, at 200 °C compared to room temperature. Modest necking is suggested in Figure 12 and from the discrepancy between reduction-in-area and tensile elongation at rupture, shown in Figure 11(a). This degree of necking is similar to that observed at room temperature. The peak true flow stress is reduced from 641 MPa at room temperature to approximately 380 MPa at 200 °C. This flow stress at 200 °C is 40 pct less than that of AA7075-T6 at room temperature, significantly less than that for over-aged AA7075-T7 at room temperature (≈ 500 MPa), and only 130 MPa greater than that for AA7075-O at room temperature (≈ 250 MPa).[19] The specific mechanism responsible for this reduction in flow stress at warm temperatures was not determined during this study, but it is expected to be related to the effects that produce retrogression at these temperatures. Regardless of the active mechanisms, these data indicate that warm temperatures near 200 °C significantly increase the tensile ductility of AA7075-T6 sheet material. The hardness data acquired confirm that strength lost during warm forming can be recovered if warm forming is conducted within the times specified for recoverable retrogression. This study demonstrates that nearly peak-aged strength (within 5 to 10 pct) can be recovered through the paint bake after warm forming.

Tensile ductility at 300 °C is not significantly improved compared to that at 200 °C. Necking at 300 °C is more severe than at 200 °C, as is demonstrated in Figure 12. Furthermore, strength cannot be effectively recovered through the simulated paint bake after retrogression at temperatures greater than 225 °C. Consequently, forming at temperatures between 225 °C and 300 °C is unlikely to be of practical use for sheet forming applications.

Tensile tests at 460 °C and 480 °C provided the greatest tensile ductilities, from 63.5 up to 125 pct. Failure at these high temperatures occurred by cavitation after significant necking, in contrast to the necking-controlled failures observed at lower temperatures. Note from Figure 11(a) that tensile ductility decreases significantly as temperature rises above 480 °C. This is because of a rapid onset in cavitation prior to significant necking above 480 °C. At 520 °C, tensile ductility is comparable to that at 200 °C. The onset of cavitation failure at temperatures above 460 °C for strain rates greater than 10⁻³ s⁻¹ is in agreement with similar data reported in the literature.[20,54] Forming at temperatures between 460 °C and 480 °C might be useful in applications where high formability is desired, the initial AA7075 temper is of no concern and subsequent heat treatments to achieve peak-aged (T6) strength that are not subject to time limits are acceptable. Hot-forming practice may be useful for low-volume production parts, but for high-volume production parts, such as in the automotive industry, high-temperature forming is cost-challenged and retrogression forming may provide a better alternative.