Research on Alloy Design and Process Optimization of Al–Mg–Zn-Cu-Based Aluminum Alloy Sheets for Automobiles with Secured Formability and Bake-Hardenability

Abstract

In this study, the compositional design of high-formability, high-bake-hardening Al–Mg–Zn-Cu-based aluminum alloys was carried out, and process conditions were established to secure mechanical properties under harsh conditions for Al–Mg–Zn-Cu-based alloys. Using JMatPro13.0 for precipitation phase simulation, the optimal pre-aging temperature and time of the design composition were selected. Through the introduction of pre-aging, it was confirmed that no over-aging phenomena occurred, even after bake-hardening, and it was confirmed that it could have mechanical properties similar to those of test specimens subjected to traditional heat treatment. Through DSC (Differential Scanning Calorimetry) and TEM (Transmission Electron Microscope) analyses, it was found that pre-aging provided sufficient thermal stability to the GP (Guinier–Preston) zone and facilitated transformation to the η’-phase. In addition, it was confirmed that, even under bake-hardening conditions, coarsening of the precipitation phase was prevented and number density was increased, thereby contributing to improvements in the mechanical properties. The designed alloy plate was evaluated as having excellent anisotropy properties through n-value and $\overline{\mathrm{r}}$-value calculations, and it was confirmed that a similar level of formability was secured through FLC (Forming Limit Curve) comparison with commercial plates.

1. Introduction

The environmental problems we face are also having a huge impact on the automobile industry, which poses a very big challenge to fuel efficiency and exhaust gas regulation, so the automobile industry is making various efforts to improve fuel efficiency. Among them, reductions in the weight of the car body are considered the most effective way to improve fuel efficiency, and aluminum alloys are considered strong candidates to replace existing steel materials due to their high specific strength, formability, and recyclability [1,2,3,4,5]. Currently, aluminum alloys are applied mainly to high-end vehicles in the automobile industry, but their use is gradually expanding. Al–Mg-based alloys, with excellent formability, are used as high-formability sheets. As automobile structural materials, Al–Mg–Si-based alloys are widely used in sheets and chassis due to medium-strength characteristics and high formability [3,4,5], and Al–Mg–Zn-Cu-based alloys with excellent specific strength are mainly used in chassis that require high-strength properties [3,6,7,8,9,10].

In the automobile manufacturing process, aluminum sheet material is formed a considerable period of time after solution heat treatment, and the mechanical properties are secured through bake-hardening in the paint bake process of BIW (Body In White) after forming [11,12,13]. Therefore, currently, Al–Mg–Si alloys are mainly used, which have high formability and relatively low yield strength, even after natural aging, and excellent bake-hardening characteristics after paint bake treatment. On the other hand, in the case of Al–Mg–Zn-Cu alloys, it is generally known that maximum strength is obtained through artificial aging at a temperature of around 120 °C [14,15,16], and excellent mechanical properties of over 500 MPa can be secured [17,18]. But, despite their high-strength characteristics, they are not preferred due to increased yield strength due to natural aging, increased sheet-forming pressure, and reduced mechanical properties due to grain coarsening at bake-hardening temperatures [7,19,20,21]. It is necessary to review a compositional design and process that can maintain the excellent mechanical properties of Al–Mg–Zn-Cu alloys, prevent adverse effects due to natural aging, and sufficiently secure mechanical properties, even under relatively harsh temperature conditions.

In the case of Al–Mg–Si alloys, many studies are being conducted on the effect of applying pre-aging techniques to control adverse effects occurring during the natural aging process after solution treatment [22,23,24,25,26]. On the other hand, in the case of Al–Mg–Zn-Cu alloys, there have not been many reviews of their application as automotive sheets, including during pre-aging, so it is necessary to study the effects of changes in natural aging and bake-hardening characteristics through the application of pre-aging treatments.

In this study, an Al–Mg–Zn-Cu-based aluminum alloy with high strength and high formability was manufactured through a compositional design that could increase formability while minimizing natural aging reactivity based on an Al–Mg–Zn-Cu aluminum alloy composition with high-strength characteristics. By using ‘JMatPro13.0’ commercial software, precipitation behavior was predicted, suitable process conditions for pre-aging were explored and reflected, and natural aging behavior and bake-hardening reactivity were evaluated when pre-aging treatments were present or absent. Differential scanning calorimetry (DSC) thermal analysis and tunneling electron microscopy (TEM) microstructure analysis were used to compare the differences in precipitation behavior and the shape of the precipitation phase, with or without pre-aging treatments. Formability-related values were calculated based on mechanical property evaluation data and compared with those of commercial materials through FLC/D formability evaluation, through which the formability of the designed alloy was discussed.

2. Materials and Methods

High-forming Al–Mg–Zn-Cu-based alloys were prepared using pure metal Al (99.9%), Zn (99.9%), Al–40Mg, Al–20Si, Al–50Cu, Al–5Fe, Al–5Mn, Al–5Cr, and Al–10Ti as raw materials in the form of a mother alloy. Before alloy manufacturing, commercial software JMatPro13.0 was used to calculate the precipitation phase formed under major heat treatment conditions, and the results of aging conditions and mechanical properties were predicted. The alloys were melted in the ceramic crucible using a high-frequency induction furnace, and the molten metal was maintained at a temperature of 710–720 °C. To minimize the turbulence of molten metal during casting, a trumpet-shaped casting mold was used, and, finally, a slab of 110 × 65 × 20 mm was manufactured. The cast slab was faceted in consideration of a chill zone, Homogenized at 510 °C for 4 h for the purpose of eliminating chemical segregation in the casting structure, and then air-cooled at room temperature. After that, the slab was preheated to 400 °C and hot-rolled at a reduction rate of 10% from 20 mm to 6 mm thickness using a 50-ton rolling machine. After air-cooling to room temperature, it was cold-rolled at a reduction rate of 30% from 6 mm to 1.2 mm in thickness to finally obtain the sound rolled sheets. The cold-rolled sheets were kept at 470 °C for 2 h and then cooled in water, pre-aged at 120 °C for 30 min, and then naturally aged for 7 days. For bake-hardening specimens, 2% pre-strain was applied to the specimens for which the natural aging was completed using a universal material testing machine, maintained at 185 °C for 25 min using an oil bath, and then air-cooled.

The actual stoichiometric composition of the cold-rolled sheets was analyzed using S-OES (Spark-Optical Emission Spectrometry, SPECTROMAXx, Spectro, Kleve, Germany). Mechanical properties were evaluated to determine the effect of each aging condition on bake-hardening, and ASTM E8 standard [27] sub-size specimens were conducted under a strain rate of 10⁻³/s using a universal material tester (MTS810, MTSC, Eden Prairie, MN, USA). To determine the cause of the difference in bake-hardening reactivity in the presence or absence of pre-aging treatments, DSC (1HT, Mettler Toledo, Columbus, OH, USA) was performed and analyzed at a rate of 10 °C/min from 50 to 350 °C in an Ar atmosphere. In addition, a high-resolution transmission electron microscope (HR-TEM, F200, JEOL, Tokyo, Japan) was used to investigate the effect of pre-aging treatments on precipitation behavior. The specimens were manufactured by mechanical surface polishing using 2400-grit SiC paper and then milling with a focused ion beam (FIB, Crossbeam350, ZEISS, Cluj-Napoca, Romania). The TEM analysis was conducted under an acceleration voltage of 200 keV, and a Fast-Fourier Transform (FFT) pattern for HR-TEM images was implemented using Digital Micrograph software(Delta™ NMR Data Processing Software). In order to evaluate the formability of the manufactured sheets, the n-value (Strain Hardening Exponent) and $\overline{\mathrm{r}}$-value (mean plastic strain ratio) of each sheet were calculated and compared with commercial high-formability sheets, and the FLC diagram was derived using a hot-sheet formability tester (forming limit diagram test machine, MTDI). The sheets that were naturally aged for 7 days were evaluated by measuring the change in pattern shape and size by applying a structural pattern to the specimens after processing into the shape of the Nakajima specimens. The test was carried out in accordance with ASTM E2218-02:2008 [28] and ISO-12004-2 standards [29], and the FLD-only experimental die and 100 mm diameter hemispherical punch of the Numisheet96 LDH Test Tooling standard were tested under conditions of a punch speed of 1 mm/s and a blank holding force of 30 t. EBSD (Electron Backscatter Diffraction) analysis was performed to determine the influence of texture components that affect the formability of sheets and was compared with commercially available alloys AA6016 and AA7075 of T4-temper.

3. Results and Discussion

3.1. Alloy and Process Design

In order to design sheet alloys with excellent formability at T4-temper while maintaining the excellent mechanical properties of the Al–Mg–Zn-Cu-based alloys, alloy design was performed based on the AA7075 material, which shows excellent mechanical properties among Al–Mg–Zn-Cu-based alloys. In Al–Mg–Zn-Cu-based alloys, Cu is dissolved in the matrix and plays a role in increasing the speed of natural aging. It was determined that hardening through natural aging would have a detrimental effect on the formability of the sheets, so the Cu content was reduced to 0.5 wt.% in the design. In terms of securing elongation, Zn was set in the range of 3.0–5.0 wt.%. In order to maintain quenching sensitivity similar to AA7075, the alloy was designed by keeping the ratio of Zn/Mg constant. The final designed composition is shown in

Table 1. Nominal chemical composition of AA7075 and design alloys. (unit: wt.%).

Elements	Zn	Mg	Cu	Fe	Ti	Cr	Mn	Si	Zn/Mg (at.%)	Al
AA7075	5.6	2.30	1.3	0.23	0.03	0.20	0.03	0.07	0.91	Bal.
M7-1	5.0	2.10	0.5	0.20	0.10	-	-	-	0.88	Bal.
M7-2	3.5	1.45	0.5	0.20	0.10	-	-	-	0.90	Bal.

.

The AA7075 alloy composition and the two designed compositions were investigated for the formation behavior of major precipitation phases (GP zone, η’, T’) according to heat treatment time using the commercial software JMatPro 13.0. Calculations were performed under conditions of 470 °C/2 h of solution temperature and 120 °C of artificial aging, and, in order to learn more about the formation tendency of the GP zone, which is the nucleation site of the η’ phase, the region between 0 and 1 h was enlarged and is shown in Figure 1. The heat treatment conditions used in the calculation were the commonly performed T6 heat treatment conditions for Al–Mg–Zn-Cu-based aluminum alloys [14,15,16]. It was confirmed that the GP zone was formed rapidly at a temperature of 120 °C, peaked within about 30 min of aging, and then transformed to η’ and decreased. Based on these results, the conditions of 120 °C and 0.5 h, which can maintain the GP zone in a sufficiently formed state, were selected as pre-aging conditions.

3.2. Mechanical Property Evaluation

In order to determine the mechanical properties at T4-temper, where sheet forming is performed, the mechanical properties of each composition were evaluated according to whether or not the prescription proceeded after solution heat treatment. Both conditions were evaluated after natural aging for 168 h. As shown in

Table 2. Comparison of mechanical properties at T4-temper with and without pre-aging treatment (P.A.).

	Non P.A. (T4)			P.A. (T4)
	Y.S. (MPa)	T.S. (MPa)	E. (%)	Y.S. (MPa)	T.S. (MPa)	E. (%)
AA7075	250 ± 4.7	415 ± 5.9	18.2 ± 0.8	252 ± 5.5	416 ± 6.1	17.0 ± 1.1
M7-1	229 ± 6.1	396 ± 6.3	25.5 ± 0.6	219 ± 5.0	380 ± 6.4	25.3 ± 0.6
M7-2	163 ± 4.3	296 ± 5.6	25.2 ± 0.4	166 ± 4.7	300 ± 5.5	24.9 ± 0.8

, no significant difference was found in the mechanical properties in the presence or absence of pre-aging. As the Zn and Mg contents decreased, the mechanical properties tended to decrease as the main precipitation phase decreased. Additionally, as the Cu content decreased compared to AA7075, the mechanical strength slightly decreased and the elongation improved. Although the GP zone and other precipitation phases predicted to be generated through pre-aging for 30 min did not make a significant difference to the mechanical properties at T4-temper, it was determined that additional analysis of bake-hardening reactivity was necessary.

After bake-hardening samples in two conditions, mechanical property evaluation was performed. Each sample was classified according to the presence or absence of pre-aging after solution heat treatment. After 168 h of natural aging and 2% pre-strain application, bake-hardening was performed in an oil bath (silicon oil) at 185 °C for 0.5 h. As shown in

Table 3. Comparison of mechanical properties at B.H.-temper with and without pre-aging. (B.H.: bake-hardening).

	Non-P.A. (B.H.)				P.A. (B.H.)
	Y.S. (MPa)	T.S. (MPa)	E. (%)	Δσ (Y.S.)	Y.S. (MPa)	T.S. (MPa)	E. (%)	Δσ (Y.S.)
AA7075	397 ± 7.4	450 ± 7.5	8.8 ± 0.4	145.7	445 ± 8.1	494 ± 7.7	8.7 ± 0.6	193.5
M7-1	289 ± 6.5	369 ± 6.9	12.1 ± 0.8	60.5	406 ± 7.2	459 ± 6.9	12.3 ± 0.7	186.4
M7-2	135 ± 4.9	246 ± 5.1	27.0 ± 1.2	−28.2	153 ± 4.5	258 ± 5.1	17.9 ± 0.9	−12.8

, it was confirmed that the specimens that underwent pre-aging treatment had higher mechanical properties than the specimens that did not undergo pre-aging. Y.C. Lin et al. [30] reported that pre-aging was applied to Al-Zn-Mg-Cu-based alloys improved corrosion resistance and creep properties, and J. Luo et al. [31] reported the contribution to mechanical properties by forming a larger number of GPII;-zones by applying pre-aging to the AA7075 alloy. Similarly, these results confirmed that applying pre-aging to the Al–Mg–Zn-Cu-based alloy can have a beneficial effect on mechanical properties. Although no significant change in T4-temper through pre-aging was observed, it could be determined that it had an advantageous effect in terms of bake-hardening reactivity. In terms of the amount of change in yield strength, the M7-1 test specimen that underwent pre-aging showed an increase in yield strength of about 186 MPa or more, while the M7-1 test specimen that did not undergo pre-aging showed an increase in yield strength of about 60 MPa. In addition, in the case of the M7-5 specimen that was not subjected to pre-aging, the mechanical properties deteriorated after bake-hardening. As Zn and Mg decreased, the alloy composition became closer to the Al–Mg series (5xxx-series), a non-heat-treatable alloy system. Also, because the amount of precipitation in the Mg–Zn-based precipitation phase to be formed through bake-hardening decreased, the amount of change in yield strength tended to decrease.

3.3. Phase Analysis

In order to analyze the cause of the difference in mechanical properties that appeared after bake-hardening, depending on the presence or absence of pre-aging treatment, DSC analysis was performed on two specimens after natural aging for 168 h. As shown in Figure 2, the GP zone peak observed in a temperature range of 100–150 °C tended to shift to higher temperatures as pre-aging was applied. J. Luo et al. [31] reported that by applying pre-aging to AA7075 alloy, the formation of GPII-zone becomes more fine and uniform, and W. Zhang et al. [32] also reported that by applying pre-aging to AA7075 alloy, a fine GP-zone was formed and that the GP-zone favored the nucleation of fine reinforcing precipitates. It was believed that this result occurred as the GP zone grew more uniformly to a critical size through pre-aging and achieved sufficient thermal stability. In addition, the η’ phase, the main strengthening phase, was observed in a temperature range of 200–225 °C, and the η phase was observed at around 250 °C. In the case of the η’ phase, as pre-aging was applied, the peak tended to shift to lower temperatures, contrary to the GP zone. It was believed that this was because the GP zone, which had a uniform critical size and thermal stability, became more susceptible to transformation into the η’ phase through pre-aging. Therefore, due to the GP zone, which has thermal stability through pre-aging, it was determined that the Al–Mg–Zn-Cu-based alloys would have improved strength characteristics without the over-aging phenomenon, even under a relatively harsh condition of 185 °C.

To analyze the microstructural causes of mechanical properties as dependent on pre-aging, TEM analysis was performed under bright-field conditions. The two specimens used in the analysis were bake-hardened specimens of M7-1 composition, and they were classified according to the presence or absence of prior aging. In the specimen without pre-aging, shown in Figure 3a,b, spherical and needle-shaped precipitation phases were observed, and each precipitation phase was confirmed to have a size of approximately 15 nm. On the other hand, in the case of the specimens with pre-aging, shown in Figure 3c,d, it was confirmed that the size was maintained at a level of 5 nm, which was significantly smaller than that of the specimen without pre-aging. Through TEM analysis, it was confirmed that the two specimens showed significant differences in the size of the precipitation phase and number density per unit area depending on the presence or absence of pre-aging. H. Jiang et al. [33] reported that high-density dispersed precipitation of the η’-phase, the main reinforcing phase of AA7055 alloy, was improved through pre-aging and, X. Liao et al. [34] reported that the main reinforcing phase, the η’-phase, was formed more finely with a higher volume fraction through the application of pre-aging to Al-Zn-Mg-Cu alloys. Also, L. Winter et al. [35] reported that by applying pre-aging to AA6056 alloy, the formation of room-temperature clusters and natural aging were suppressed after solid solution treatment, and the artificial aging reactivity was greatly improved. It was confirmed that pre-aging was advantageous to the mechanical properties of the designed alloy based on the size of the precipitate phases that were maintained at a smaller size through pre-aging and the number density of the precipitate phases per unit area. This is because the GP zone grew uniformly with thermal stability due to pre-aging, as shown in the DSC analysis results. Additionally, for this reason, the Al–Mg–Zn-Cu alloys could effectively serve as a nucleation site for a finer and more uniform precipitation phase under relatively harsh conditions.

3.4. Texture Analysis

To further examine the various factors affecting the formability of sheets, EBSD (Electron Backscatter Diffraction) analysis was performed to determine the influence from the perspective of texture. Additionally, samples were compared with commercially available alloys AA6016 and AA7075. In the case of two types of commercial sheets, T4-temper specimens were used, and the M7-1 alloy was analyzed after solution treatment and pre-aging, and natural aging for one week. From the GB (Grain boundary) Map and average misorientation results in Figure 4, it can be seen that AA6016, known as a commercial high-formability sheet, has a larger grain size than the M7-1 and AA7075 alloys. In addition, it can be confirmed that a significant amount of recrystallized grain growth has occurred, as the fraction of low-angle grain boundaries is relatively small and the fraction of high-angle grain boundaries has increased. Figure 5 shows the ODF (Orientation distribution function) and pole figure used to analyze the crystal orientation distribution of the AA6016, AA7075, and M7-1 alloys. In the case of the AA6016 sheets, mainly, it can be seen that a strong cube orientation developed, and the recrystallization texture also developed strongly. On the other hand, AA7075 was observed to have a mixture of a recrystallized texture and deformed texture. In the case of the M7-1 alloy, deformed textures, such as brass, cupper, and S, were mainly observed [36]. Therefore, in terms of the anisotropy of the sheets, it was judged that the M7-1 alloy would be advantageous compared to AA6016, a commercially available high-formability sheet.

3.5. Formability Evaluation

To compare the formability of cold-rolled sheets, the strain-hardening index n-value and average plasticity ratio $\overline{\mathrm{r}}$-value were calculated and compared with commercial high-formability sheets (

Table 4. n-value and $\overline{\mathbf{r}}$ -value of design and commercial alloys.

	n-Value	$\overline{\mathbf{r}}$-Value
6016/T4	0.270	0.650
6011/T4	0.260	0.560
5182/O	0.310	0.750
7075/T4	0.203	0.650
M7-1/T4	0.240	0.724
M7-2/T4	0.260	0.740

). The n-value and $\overline{\mathrm{r}}$-value were each calculated by the formulas below.

(1)

$\mathsf{\sigma }=\mathrm{K}{\mathsf{ϵ}}^n$

• True stress $\mathsf{\sigma }=\mathrm{S}\left(1+\mathrm{e}\right),\mathrm{T}\mathrm{r}\mathrm{u}\mathrm{e} \mathrm{s}\mathrm{t}\mathrm{r}\mathrm{a}\mathrm{i}\mathrm{n} \mathsf{ϵ}=\ln \left(1+\mathrm{e}\right)$
• (K: strength coefficient, S: engineering stress, e = engineering strain)
• $\mathrm{l}\mathrm{o}\mathrm{g}\mathsf{\sigma }=\mathrm{l}\mathrm{o}\mathrm{g}\mathrm{K}+\mathrm{n}\log ϵ$ [ASTM E646] [37]

(2)

r = $\frac{{\epsilon }_w}{{\epsilon }_t}$ (Where ${\epsilon }_w$ and ${\epsilon }_t$ are the strains in the width and thickness directions),

• $\overline{\mathrm{r}}$ = $\frac{r_0+r_{90}+2r_{45}}{4}$ ($r_0{, r}_{90,} r_{45}$: Along the rolling direction (RD), along 45° to RD and perpendicular to rolling direction) [38].

T4-temper samples of the design alloys were calculated to have n-values in a range of 0.20–0.26 and an $\overline{\mathrm{r}}$-value in a range of 0.715–0.740. In the case of the n-value, which can evaluate the formability of sheets, it showed a higher value than the AA7075 alloy and a similar or slightly lower value than alloy sheets, such as AA5182, AA6011, and AA6016, which are used as high-formability sheets. On the other hand, the $\overline{\mathrm{r}}$-value, which can evaluate the uniformity of the reduction rate in each direction, has a higher value compared to commercial alloys and is predicted to have excellent properties in terms of anisotropy.

For a more precise comparative evaluation of formability, FLC/D formability evaluation of the M7-1 alloy was performed and the results were compared with the FLC curve of commercial sheets (Figure 6). Similar to the n-value calculation results, it was confirmed that the FLC was located in a higher area than the AA7075 sheet, resulting in a significant improvement in formability compared to the commercial Al–Mg–Zn-Cu alloys. In addition, when compared with 5xxx- and 6xxx-series commercial high-formability plates, it was confirmed that a significant level of formability was secured, as the FLC was located in a similar area.

4. Conclusions

In this study, the composition was designed to secure formability, which has previously been identified as the biggest disadvantage in Al–Mg–Zn-Cu-based aluminum alloy. In addition, in order to explore the possibility of use as a sheet material for automobiles, process conditions capable of securing mechanical properties under bake-hardening conditions were investigated.

Based on the composition of AA7075, a representative high-strength Al–Mg–Zn-Cu-based alloy, the Cu content was reduced to minimize natural aging reactivity, and the Zn and Mg contents were changed to observe changes in mechanical properties and formability. As the Zn and Mg contents decreased from 5.6 to 3.5 wt.% and 2.3 to 1.45 wt.%, respectively, the yield strength was observed to decrease by 445 to 153 MPa due to a decrease in the η’-phase, which is the main strengthening phase. Through calculation using the commercial software JMatPro13.0, it was expected that the GP zone would form close to the peak through heat treatment at 120 °C for about 30 min, and, based on this, the pre-aging period was introduced under conditions of 120 °C/30 min. Through the application of pre-aging treatment, bake hardening reactivity was improved to 60.5 to 186.4 MPa based on the change in yield strength, and it was confirmed to be similar to the heat treatment characteristics of the traditional 7xxx-based alloys. Through DSC analysis, it was expected that pre-aging would provide sufficient thermal stability to the GP zone, facilitating transformation into the η’ phase. Through TEM analysis, it was confirmed that pre-aging contributed to an improvement in mechanical properties by preventing the coarsening of the precipitation phases and increasing the number density.

The microstructure and physical properties were compared with those of commercial materials through n-value and $\overline{\mathrm{r}}$-value calculations and EBSD analysis, and it was predicted that the M7-1 composition showed an $\overline{\mathrm{r}}$-value of 0.724, would show excellent properties in terms of anisotropy. In addition, the FLC/D formability evaluation showed that it had a level of formability similar to that of 5xxx- and 6xxx-series alloys, which are high-formability commercial sheet alloys.

In this study, compositional design and an evaluation of mechanical properties were performed to ensure that excellent formability was achieved for Al–Mg–Zn-Cu-based alloys while their high-strength characteristics were maintained. Process conditions that can maintain excellent mechanical properties, even under automobile manufacturing process conditions, were provided, and, through these, new guidelines were provided for the application of Al–Mg–Zn-Cu-based alloy sheets to the automobile industry.