Effect of Trace Er Addition on the Microstructural Evolution and Heat Resistance Properties of an Al-Zn-Mg-Cu Alloy During High Temperature Tensile and Thermal Exposure

Abstract

High temperature tensile properties and long-term thermal stability play an important role in practical applications of Al-Zn-Mg-Cu alloys. In order to evaluate the effect of Er addition on the properties of an Al-Zn-Mg-Cu alloy as potential high temperature structural materials, the heat resistance properties of an Al-Zn-Mg-Cu alloy were investigated at various temperatures. After high temperature tensile testing and long periods of heat exposure testing, the microstructures of Al-Zn-Cu-Mg alloys with and without small Er addition is intentionally investigated by X-ray diffraction (XRD), scanning electron microscopy (SEM), and quantitative transmission electron microscopy (TEM) characterization to explore the potential effect of Er on the tensile properties. The experimental results reveal that the heat resistance of T76-tempered Al-Zn-Cu-Mg alloy is obviously improved by adding trace Er. The Al₈Cu₄Er phase is found to segregate at the localized regions along grain boundaries and strengthens the grain boundaries at elevated temperatures. The η′ and η precipitation is obviously promoted by adding trace Er, and dispersed nano-sized Al₃(Er, Zr) precipitates were formed in the Er-containing alloys after homogenization, thereby enhancing the strength of Al-Zn-Mg-Cu. In addition, precipitates in both alloys gradually coarsen with the increase in thermal exposure temperature and the extension of thermal exposure time. The influence of precipitates on mechanical properties of the investigatived alloy after thermal exposure is also discussed.

1. Introduction

Aluminum alloy is used as the main structural material to prepare important parts and components in the industry due to its low density, high strength, and good stress corrosion resistance [1,2,3]. In the high temperature service environment, the mechanical properties of the alloy will be reduced due to the deterioration of the microstructure, which will cause serious hidden dangers to the safe use of various components [4,5]. Over the last decade, researchers have conducted studies related to heat-resistant aluminum alloys, with the main studies focusing on high temperature thermal stability. However, most researchers only focus on the study of traditional Al-Cu-Mg series heat-resistant aluminum alloys. Relevant researchers have studied the microstructure evolution or mechanical properties loss of 2014, 2A12, and 2219 aluminum alloys after various thermal exposure conditions [6,7,8]. The high temperature environment of the heat-resistant aluminum alloy (Al-Cu-Mg series alloys) studied by the above researchers is mainly concentrated in the range of 200 °C~300 °C [8,9]. The Al-Zn-Mg-Cu series high strength aluminum alloys are widely used due to its outstanding specific strength advantages. Previous research on the 7xxx series of alloys has focused on improving their stress corrosion resistance or toughening, with relatively little research on their heat resistance. However, when the airframe parts or oil drill pipes prepared by an Al-Zn-Mg-Cu alloy are used in the field of aviation industry and oil drilling, they are inevitably used in high temperature service conditions [10,11]. Due to the high temperature environment near the engine, the drilling depth of the oil drill pipe is improved, which will require the aluminum alloy to have better thermal stability to ensure safe service. Presently, relatively few systematic studies have been carried out on the correlation between the microstructure evolution and tensile properties of Al-Zn-Mg-Cu series alloys at different high temperatures and long periods of thermal exposure.

The mechanical properties of aluminum alloys can be significantly improved by adding appropriate amounts of rare Earth elements. Therefore, relevant researchers have extensively explored the effects of common rare Earth elements such as Sc, Y, and La on the microstructure evolution and mechanical properties of aluminum alloys [12,13]. Secure applications of Al-Zn-Mg-Cu alloys require stable microstructure under medium and high temperature service conditions. Since the actual properties of Al-Zn-Mg-Cu alloys are dependent on the microstructure evolution of precipitated phases, it is obvious that the intensive precipitation of precipitated phases with good thermal stability will contribute significantly to the heat resistance of Al-Zn-Mg-Cu alloys. In contrast to the expensive Sc, Yb, and Y elements, some studies have opted for the relatively inexpensive addition of rare Earth Er elements to improve the properties of the alloys. The strength of the Al-Zn-Mg-Cu alloy can be increased by adding trace amounts of Er without reducing the elongation of the alloy [14,15]. Meanwhile, the addition of trace Er can increase the recrystallization temperature of the alloy, which will lead to the refinement of the grain structure of the alloy [16]. Although the effect of Er on the microstructure and properties of aging-state Al-Zn-Mg-Cu alloys was discussed previously, there is no available literature that can elucidate in detail the close relationship between the addition of trace Er and the microstructure evolution during high temperature stretching and prolonged thermal exposure under medium to high temperature service conditions. Summarizing previous studies, it can be found that Er-microalloying has not addressed the microstructure evolution and performance changes in the alloy under high temperature service conditions.

In present work, two Al-Zn-Mg-Cu alloys (named Er-free alloy and Er-added alloy, respectively) were designed to investigate the effect of Er on the heat resistance of the alloys. Thermal stability tests were carried out on both alloys at different high temperatures (150 °C, 165 °C, 180 °C) and long period service conditions. Compared with the traditional Al-Cu-Mg heat-resistant alloy (high temperature service temperature range of 200~300 °C), the heat resistance of the Al-Zn-Mg-Cu alloy is relatively poor. The material used in this study is an Al-Zn-Mg-Cu alloy, and its actual high temperature service temperature is relatively low, so 150~180 °C is selected as the test temperature range.

Meanwhile, the coarsening behavior and phase transformation mechanism of alloy precipitates during thermal exposure were thoroughly investigated. In addition, the alloys were also tested for room and high temperature tensile properties and corresponding microstructural analyses were carried out. Finally, the above research results are summarized and analyzed. This study will provide some experimental data and theoretical analysis for the preparation of a high strength Al-Zn-Mg-Cu heat-resistant alloy and expand the potential application of Al-Zn-Mg-Cu under high temperature service conditions.

2. Experimental Procedure

The actual chemical compositions of the investigated alloys are displayed in

Table 1. Chemical compositions of studied Al-Zn-Mg-Cu-(Er) alloys with different Er contents (in wt%).

Alloy	Zn	Mg	Cu	Mn	Zr	Ti	Er	Fe	Si	Al
Er-free	6.51	1.83	1.28	0.29	0.98	0.07	-	0.04	0.04	Bal.
Er-added	6.48	1.81	1.28	0.30	0.10	0.06	0.10	0.03	0.05	Bal.

, and Er-free and Er-added alloys refer to Er additions of 0 wt% and 0.1 wt%, respectively.

Two self-designed alloys were melted and cast after degassing and standing. The purity of the pure metal (Al, Zn, Mg) and the intermediate alloy (Al-50 wt% Cu, Al-50 wt% Mn, Al-10 wt% Zr, Al-20 wt% Ti, Al-20 wt% Er) used in the smelting ingot is more than 99.99%, so as to ensure that the alloy contains fewer impurities. Two-stage homogenization treatment (425 °C/24 h + 465 °C/48 h, with a heating rate of 20 °C/h) was performed on the two as-cast ingots to eliminate segregation. Subsequently, the ingot was held at 440 °C for 2 h and then hot rolled to 3 mm sheet. The alloy sheets were subjected to solid solution treatment at 470 °C for 1 h and immediately subjected to water quenching and two-stage aging heat treatment (T76: 120 °C/15 h + 165 °C/25 h). The investigated samples after aging treatment were subjected to long-period thermal exposure at high temperatures of 150 °C, 165 °C, and 180 °C, respectively. Then, tensile tests were performed. The potential effect of Er addition on the thermal stability of the alloy was evaluated by a thermal exposure test. The tensile specimens were prepared parallel to the transverse direction (Transverse–Longitudinal, T-L) of the hot rolled sheets, and the gauge length and gauge width are 30 mm and 12.5 mm, respectively. For the tensile test of two alloys after thermal exposure treatment, three samples were tested on an MTS-810 tensile machine (MTS Systems Corporation, Eden Prairie, MN, USA) with a strain rate of 1.1 × 10⁻³ s⁻¹ to ensure the accuracy of the test results. The tensile test at different temperatures (150 °C, 165 °C, and 180 °C) is performed on tensile machine model SANSCMJ5105 (SANS, Shenzhen, China), and the temperature error is stable at ±3 °C.

In order to investigate the intrinsic mechanism for the effect of Er addition on the heat resistance of the alloys, the microstructures of Er-free and Er-added alloys were intensively characterized using scanning electron microscopy and transmission electron microscopy. The morphology and distribution of the second phase in the alloy were observed by FEI Quanta-200 SEM, Whis was provided by FEI Company from the Hillsboro, OR, USA. Fracture analysis of tensile specimens was also carried out using the same scanning electron microscope at 20 KV. In order to identify the nano-scale precipitates and observe the microstructure evolution of the strengthening phase after thermal exposure, the disks were cut into 3 mm diameters and mechanically grinded to a thickness of about 60 μm, and twin-jet electropolishing was performed for transmission electron microscopy (TEM, Tecnai G2 20, This was provided by FEI Company from the Hillsboro, OR, USA). When statistically analyzing the number density, particle size, and volume fraction of the precipitates, at least 10 TEM images were selected to reduce the error to ensure the reliability of the statistical results. Specimens for electron backscattered diffraction (EBSD) observation were mechanically grinded and electrolytically polished, thus ensuring a flat and low-stress test surface. The phase identification and analysis were carried out by XRD testing, and the test surface also needs to be mechanically ground and polished before testing.

3. Results

3.1. The High Temperature Tensile Properties

Figure 1 shows the high temperature tensile properties of Al-Zn-Mg-Cu-(Er) alloys at different high temperature conditions (tensile temperatures of 150, 165, and 180 °C, respectively, with a holding time of 10 min). In Figure 1a, it can be observed that the high temperature tensile properties of Er-added alloys are higher than those of Er-free alloys when the tensile temperature is 150 °C. The tensile strengths of the Er-added alloy and Er-free alloy were 587 and 556 MPa, respectively. The yield strengths of the Er-added alloy and Er-free alloy were 540 and 514 MPa, respectively. The results show that the addition of trace Er can not only improve the tensile properties of the Al-Zn-Mg-Cu alloy at room temperature, but also improve the high temperature tensile properties of the alloy. The high temperature tensile properties of the two alloys at 165 °C are shown in Figure 1b. The tensile strength and yield strength of the Er-added alloy are 517 MPa and 482 MPa, respectively. The tensile strength of the Er-free alloy is 551 MPa and the yield strength is 511 MPa. With the increase in tensile temperature from 150 °C to 165 °C, the high temperature tensile properties of the two alloys decrease further. However, the high temperature tensile properties of the Er-added alloy are still higher than those of the Er-free alloy.

The high temperature tensile properties of the two alloys at 180 °C are shown in Figure 1c. The ultimate tensile strength of the Er-free alloy is 469 MPa, and the yield strength is 425 MPa. For the Er-added alloy, they are 505 MPa and 457 MPa, respectively. Comparing the high temperature tensile properties of the two alloys at various temperatures, it can be seen that the high temperature tensile properties of the two alloys show a continuous decrease, which is attributed to the increase in tensile temperature. Under the same tensile conditions, the high temperature tensile properties of the Er-added alloy are always better than that of the Er-free alloy, which indicates that the addition of Er has a positive effect on improving the high temperature tensile properties of the Al-Zn-Mg-Cu alloy.

The fracture surfaces of high tensile samples of both aging alloys are shown in Figure 2. The fracture morphology of the high temperature tensile specimens of both alloys is typical ductile trans-granular fracture, and there are dimples of different sizes in the fracture structure. It can also be observed from Figure 2 that the number of dimples in the tensile fracture of the Er-added alloy is higher than that of the Er-free alloy at the same temperature, which indicates that the elongation of the Er-added alloy is higher than that of the Er-free alloy. In addition, with the increase in tensile temperature, the number of dimples in the fracture structure of the high temperature tensile samples of the two alloys increases, which indicates that the elongation of the two alloys increases with the increase in tensile temperature. These fracture microstructure analyses are consistent with the elongation changes in both alloys of the high temperature tensile specimens shown in Figure 2.

3.2. Thermal Stability Testing After Exposure to 150 °C, 165 °C, and 180 °C for 600 h

In order to reveal the influence mechanism of trace Er addition on the thermal stability of the alloys after long-term service at high temperature, the two alloys were first subjected to T76 aging heat treatment. Subsequently, the two alloys were subjected to sustained thermal exposure at 150, 165, and 180 °C. The tensile test results of the samples after long-term thermal exposure are presented in Figure 3. It can be observed from Figure 3a,b that the tensile properties of the two alloys show a decreasing trend while the elongation shows a small increase during the increasing thermal exposure time; the elongation results are presented in Figure 3c. The ultimate tensile strength (UTS) of the Er-free alloy decreased from 620 MPa to 443 MPa, which decreased by 28.5%, the yield strength (YS) decreased from 558 MPa to 380 MPa, which decreased by 31.9%, and the elongation increased from 9.8% to 12.1% after 600 h thermal exposure at 150 °C. The tensile strength of the Er-added alloy decreased from 654 MPa to 495 MPa, with a decrease of 7.3%, the yield strength decreased from 586 MPa to 410 MPa, with a decrease of 7.0%, and the elongation increased from 10.3% to 13.3%.

After increasing the thermal exposure temperature to 165 °C, the UTS, YS, and elongation of the Er-added alloy are still higher than those of the Er-free alloy. The detailed test results are shown in Figure 3d–f. After thermal exposure at 165 °C for 600 h, the UTS and YS of the Er-free alloy are 232 MPa and 172 MPa, respectively. The UTS and YS of the Er-added alloy are significantly higher than those of the Er-free alloy after adding trace Er to the Al-Zn-Mg-Cu alloy. In addition, the elongation of the alloy after adding Er is also slightly improved. Comparing the elongation of the two alloys, it is observed that the elongation of the Er-added alloy is still slightly higher than that of the Er-free alloy even if the thermal exposure temperature is further increased.

Increasing the thermal exposure temperature will undoubtedly deteriorate the tensile properties of the alloy, as higher heat exposure temperatures will result in a more significant decrease for the tensile properties of the alloy. It is observed in Figure 3g–i that the tensile properties of the Er-added alloy are still higher than those of the Er-free alloy even after the thermal exposure temperature is increased to 180 °C. Under different high temperature conditions, the Er-added alloy has always shown higher tensile properties than the Er-free alloy, which indicates that the addition of Er can indeed improve the heat resistance of the alloy, which undoubtedly broadens the application of the alloy under high temperature service conditions.

The hardness evolution of the Al-Zn-Mg-Cu-(Er) alloy at different thermal exposure temperatures is shown in Figure 4. The thermal exposure testing of the two alloys at high temperatures will lead to a decrease in hardness, and the higher the thermal exposure temperature, the more significant the decrease in hardness. It should be noted that although the hardness of both alloys shows a decreasing trend, the hardness of the Er-added alloy has always been higher than that of the Er-free alloy.

3.3. XRD Patterns

Figure 5 shows the XRD results of the as-cast and homogenized Er-free and Er-added alloys. The phases in both alloys were identified and analyzed by X-ray diffractometry (D/Max 2500) using Cu-Ka radiation. According to the test results in Figure 5, it can be observed that the diffraction peaks corresponding to Al, MgZn₂, and Al₂CuMg are observed in the as-cast and homogenized Er-free alloy. When analyzing the XRD test results of the as-cast and homogenized Er-added alloy, it can also be observed that there are also diffraction peaks corresponding to Al, MgZn₂, and Al₂CuMg in the Er-added alloy. In addition, there are Al₈Cu₄Er corresponding diffraction peaks in the Er-added alloy.

3.4. EBSD Analysis

The two alloys were characterized using EBSD analysis with the aim of fully investigating the effect of added Er on the crystal structure of the alloys, and detailed results are presented in Figure 6. There is no obvious grain orientation in both alloys, which was confirmed by the IPF results in Figure 6a,b. Figure 7 describes the grain size distribution of the two alloys after T76 treatment. Combined with Figure 6a,b and Figure 7, it can be intuitively found that trace Er was added to refine the grain size of the Al-Zn-Mg-Cu alloy. The average grain sizes of the two alloys were 18.4 μm and 16 μm, respectively. Compared with the Er-added alloy, the grain size distribution range of the Er-free alloy is larger, where individual grain size was more than 100 μm, while the grain size distribution of the Er-added alloy was relatively concentrated. Grains with sizes less than 20 μm (equivalent circle diameter) dominate the grain structure of the two alloys. The histogram of the grain boundary angle distribution in Figure 6c,d shows that the proportion of HAGBs in all alloys was much higher than that of LAGBs, which occupies a dominant position. The misorientation angle reveals an increase in the fraction of LAGBs of the Er-free alloy compared with the Er-added alloy.

3.5. SEM Observation

Figure 8 illustrates the backscattered scanning map of the as-cast and homogenized Er-added alloy and the energy dispersive spectroscopy analysis results of each element, and the EDS analysis of the second phase at the grain boundary is also carried out. Figure 8a is the EDS analysis results of the as-cast Er-added alloy, and Figure 8b is the EDS analysis results of the homogenized alloy. According to the distribution of each element in Figure 8a,b, the phases marked with A and B points at the grain boundary in Figure 8a,b contain Al, Cu, and Er elements. In order to further accurately analyze the elemental composition of the phases at point A and point B, EDS point analysis was performed, and the results are shown in Figure 8c,d. According to the EDS analysis results, the ratio of Cu and Er atoms in the Er-containing phase of point A and point B is about 4:1, which is very close to the ratio of Cu and Er atoms in the Al₈Cu₄Er phase. The EDS analysis results shown in Figure 8 are consistent with the XRD patterns in Figure 5, which further confirms the presence of the Al₈Cu₄Er phase in the Er-added alloy.

3.6. Conventional TEM Characterization

The high temperature tensile strength and heat resistance of the alloy were improved after the addition of trace Er. XRD (Figure 5), EBSD (Figure 6 and Figure 7), and SEM (Figure 8) analyses also revealed potential changes in the microstructure of the alloy. In order to comprehensively analyze the effect of Er addition on the microstructure evolution of the alloy, TEM analysis (provided in Figure 9, Figure 10 and Figure 11) was used to further characterize the nano-precipitates, and quantitative analysis was carried out.

3.6.1. Precipitation Behavior

The bright field TEM images of nano-scale precipitates in the two alloys after T76 aging treatment are presented in Figure 9a,b, while Figure 9c illustrates the dark field TEM images of the Al₃Er phase in the Er-added alloy after two-stage homogenization treatment, and TEM images in a <110> matrix zone axis. The diffraction spots at 1/3{220}α and 2/3{220}α can be clearly observed in the selected area electron diffraction (SAED) pattern, which confirms the existence of a semi-coherent η′ phase in both alloys. The diffraction spots at the 1/2{022}α position indicate that η precipitates appear in the alloy. The fine η′ and η precipitates dispersed in the Al matrix can be clearly observed from the bright field TEM images (Figure 9). For the heat-treatable Al-Zn-Mg-Cu alloy, the η′ and η phases are the main strengthening phases, which play an important role in improving the tensile properties of the alloy. The number density of nano-precipitates in the Er-added alloy is higher than that in the Er-free alloy, which can be clearly observed in Figure 9a,b. The size of precipitated phases is relatively smaller, and the distribution of precipitated phases is relatively more uniform and dispersed. By comparing the microstructure differences, it is found that the Er-added alloy has higher number density, smaller particle size, and more dispersed precipitated phases, which ultimately leads to the fact that the Er-added alloy has better mechanical properties than the Er-free alloy. Figure 9c shows the Al₃(Er, Zr) particles and the corresponding SAED patterns observed in the dark field mode of the homogenized Er-added alloy. These particles were formed during homogenization, and subsequent heat treatment did not change their microstructure, indicating that they have high thermal stability. The particles with high thermal stability will undoubtedly promote the heat resistance of the Al-Zn-Mg-Cu alloy.

3.6.2. After Further Exposure to 150 °C, 165 °C, and 180 °C for 600 h

After thermal exposure at 150 °C for 600 h, the bright field TEM images of the precipitated phases of the two alloys are presented in Figure 10a,d. It can be clearly observed that the morphology of the precipitates of the two alloys changed significantly after thermal exposure, and more spherical and rod-shaped precipitates appeared in the intragranular. Long-term high temperature thermal exposure leads to the coarsening of the intragranular precipitates of the two alloys, and most of the η′ phase is transformed into incoherent η phases. The coarsening rate of precipitates in the Er-added alloy is lower than that in the Er-free alloy. The SAED results in Figure 10 confirm the presence of Al₃(Er, Zr) particles in the Er alloy.

As shown in Figure 10b,e, two alloys were subjected to thermal exposure tests under more severe service conditions (165 °C/600 h), which resulted in a more significant coarsening of the intragranular precipitates. As shown in Figure 10b,e, the strong diffraction spots at 1/3{220}α, 2/3{220}α, and near 1/2{022}α in the SAED pattern also confirm this result.

After the two alloys were subjected to thermal exposure at 180 °C for 600 h, the η phase became the dominant strengthening phase, and few η phases were observed. This can be supported by the bright field TEM images and the corresponding SAED patterns presented in Figure 10c,f. Strong diffraction spots were observed at the 1/2{022}α positions, but the diffraction spots of the η′ phase were not observed. It should be noted that the diffraction spots of Al₃(Er, Zr) particles are still observed in the SAED pattern presented in Figure 10f. This indicates that Al₃(Er, Zr) has excellent heat resistance, so that the thermal stability of the Er-added alloy is higher than that of the Er-free alloy, which is consistent with the test results of the tensile properties of the alloy after thermal exposure described in Figure 3g,h.

The high-resolution TEM (HRTEM) images in Figure 11 display the microstructure evolution of the intragranular precipitates of two alloys after thermal exposure at various temperatures for 600 h. The precipitates indicated by the yellow arrow and the red dotted lines in Figure 11a–f are the η′ and η phases, respectively. The η′ phase and the η phase exist in the two alloys. In comparison, the particle size of the precipitates in the Er-added alloy is relatively fine. The size of the precipitates in the Er-added alloys is still smaller than those in the Er-free alloys, despite the fact that coarsening and growth of the precipitated phase with the increase in thermal exposure temperature. This result can be verified by comparing the HRTEM images of precipitated phases after thermal exposure at 150 °C, 165 °C (Figure 11b,e), and 180 °C (Figure 11c,f), which is also confirmed by the effect of thermal exposure temperature on the microstructure evolution of the alloy described in Figure 10. Interestingly, the aspect ratio of the plate-like η phase and the rod-like η phase also increases with the increase in thermal exposure temperature. In addition, it can also be clearly observed that the number density of precipitated phases decreases significantly with the increase in thermal exposure temperature.

The statistical histogram and average size of the particle size distribution of the intragranular precipitates are shown in Figure 12. The statistical results are obtained by analyzing the TEM images according to the method introduced in the literature [16,17]. In order to obtain the results shown in Figure 12 and ensure that the statistical results can be reliably used to analyze the microstructure evolution of both alloys during thermal exposure, no less than 10 TEM images and more than 1000 precipitates were used for statistical analysis. After thermal exposure at 150 °C for 600 h, the radius of precipitate in the Er-added alloy (9.17 nm) is smaller than that of Er-free (9.72 nm), as shown in Figure 12a,d. The statistical results show that the proportion of precipitates with radii less than 3 nm in Er-added is 10.2%, while that in Er-free is 9.8%. The radius of most precipitates in the Er-free alloy is more than 3 nm, and the maximum radius can reach 20 nm, while that in the Er-added alloy is 15 nm, and the proportion of precipitates with a radius of more than 3 nm is lower than that in the Er-free alloy. The tensile test results (see Figure 3a–c) after thermal exposure at 150 °C for both alloys correspond to a large difference in the radius of the precipitates. The radius distribution and average radius of all alloy precipitates after 165 °C and 180 °C thermal exposure are similar to those after 150 °C thermal exposure. The proportion of small-sized precipitates in the Er-added alloy is always higher than that in the Er-free alloy. It should be noted that there is almost no precipitated phase with a radius of less than 3 nm after thermal exposure at 165 °C and 180 °C for 600 h, and the precipitates are obviously coarsened due to the increase in thermal exposure temperature.

In order to intuitively and comprehensively describe the coarsening behavior of precipitates during high temperature thermal exposure, and also to facilitate the subsequent calculation of the contribution of precipitation strengthening to the yield strength of the alloy, the detailed characteristics of the intragranular precipitates of the two alloys after various high temperature thermal exposure tests are summarized in

Table 2. Precipitate characteristics of two alloys after thermal exposure testing for 600 h (acritical radius is 3 nm).

Alloy	Expose Temperature/°C	Precipitate	f/%	r/nm	Dislocation Mechanism
Er-free	150	η′ + η	10.7	9.72	9.8/% Shear + 90.2% Bypass
	165	η′ + η	7.2	19.63	Bypass
	180	η	6.8	27.89	Bypass
Er-added	150	η′ + η + Al3(Er, Zr)	11.6	9.17	10.2/% Shear + 89.8% Bypass
	165	η′ + η + Al3(Er, Zr)	9.7	16.53	Bypass
	180	η + Al3(Er, Zr)	8.1	24.68	Bypass

, where the f_v of the precipitated phases is obtained by the method described in the refs. [18,19,20]. During the long-period thermal exposure, the main phase transformation in the alloy is the transformation of the metastable η′ phase into the stable η phase. The volume fractions of the precipitated phases were 10.7%, 7.2%, and 6.8% for the Er-free alloys after thermal exposures at 150 °C, 165 °C, and 180 °C for 600 h, respectively. After the same thermal exposure treatment, as illustrated in Table 2, the volume fraction of the precipitated phase of the Er-added alloy exhibited significant difference compared with the Er-free alloy, while a higher volume fraction of the precipitated phase exists in the Er-added alloy.

4. Discussion

4.1. Effect of Er Addition on High Temperature Tensile Properties

Considering the complexity of the service environment of an Al-Zn-Mg-Cu alloy for aerospace applications, it is indispensable to obtain prominent high temperature mechanical properties and thermally stable microstructures while maintaining excellent room temperature tensile properties. As described in Figure 1, the tensile properties of T76 tempered Er-free and Er-added alloys exhibit decreasing trends with the increase in tensile temperature (from 150 °C to 180 °C). The high temperature tensile test results indicate that the yield strength and tensile strength of the Er-added alloy are higher than those of the Er-free alloy at the same tensile temperature. The addition of the Er element can significantly improve the high temperature properties of T76 tempered alloy. In the present work, the investigated alloys are heat-treatable strengthened alloys and the yield strength is mainly determined by the resistance to dislocation motion. For the Er-free and Er-added alloys, the activation energy of dislocation motion and the hindrance of precipitates to dislocations combine to affect the yield strength. The high temperature stretching of the alloy can be considered as a continuous aging process at high temperature. The fine precipitated phase (such as the GP zone and the η′ phase) coarsens rapidly during high temperature tensile process, which is different from the room temperature tensile. The rapid coarsening of the precipitates will inevitably reduce the high temperature tensile properties of the alloy. In addition to the above factors, with the increase in tensile temperature, the α-Al matrix in the Er-free and Er-added alloys continues to soften, which will reduce the activation energy of dislocation motion. Therefore, the yield strength of the alloys will also decrease. During high temperature tensile testing, the alloy undergoes recrystallization behavior, which leads to a decrease in strength. Meanwhile, the high deformation energy stored during stretching can provide the driving force for dynamic recrystallization (DRX). DRX consumes a large number of dislocations, which is also the dominant factor in the decrease in strength. Despite the fact that the increase in tensile temperature leads to a decrease in tensile strength, the high temperature tensile strength of the Er-added alloy is always higher than that of the Er-free alloy. This is due to the Al₃Er phase significantly inhibiting DRX and improving the thermal stability of the microstructure. In addition, the addition of trace Er can inhibit the coarsening of aging precipitates in high temperature environments and still hinder dislocation movement during high temperature tensile process, which further improves the high temperature tensile strength of the Er-added alloy. Different from the decreasing trend of tensile strength, the increase in fracture elongation with increasing tensile temperature is attributed to the fact that the increase in tensile temperature favors dynamic recovery during deformation and promotes dislocation migration in the alloy and activation of more slip systems, thus increasing alloy elongation.

In addition, the tensile properties of alloys are usually affected by grain boundaries. When the alloy is stretched at high temperatures, the grain boundary is a relatively weak position. Therefore, if the microstructure at the grain boundary can be controlled in a targeted manner, the high temperature tensile properties of the alloy can be effectively improved. The addition of trace amounts of Er in the alloy resulted in the formation of the Al₈Cu₄Er phase with a high melting point and distributed at the grain boundaries, as depicted in Figure 5 and Figure 8. According to the research results in refs. [21,22,23], the Al₈Cu₄Er phase has its own high thermal stability and melting point. When the alloy is stretched at high temperature, the Al₈Cu₄Er phase can be effectively “pinned” at the grain boundary, so that the grain boundary is not easy to slip at high temperatures and so that the cracking and deformation of the grain boundary are delayed to a certain extent. The high temperature tensile properties of the alloy are improved due to the presence of the Al₈Cu₄Er phase with high thermal stability.

The fracture behavior of an alloy is closely related to the tensile testing temperature and microstructure evolution. Unlike room temperature conditions, the thermal activation and softening behavior of the alloy have a significant effect on the fracture behavior when the tensile temperature is 150 °C. The strength difference between the grain interior and the grain boundary is still significant, which eventually leads to intergranular fracture as the main fracture mode. Meanwhile, due to the coarsening of the precipitated phase and the softening of the α-Al matrix, the dimple size in the high temperature tensile fracture is larger than that in the room temperature tensile fracture. When both alloys are tensile tested at 165 °C and 180 °C, the formation and growth of voids have a significant effect on the fracture behavior of alloys. As the tensile temperature increases, the size of the grain boundary precipitates and the width of the precipitation-free zone (PFZ) becomes larger. In the process of high temperature stretching, the “soft-hard” coupling between the coarse grain boundary precipitates (GBPs) and the wide PFZ causes dislocations to accumulate around the grain boundary precipitates. With the coarsening of grain boundary precipitates and the increase in PFZ width, the size of dimples becomes larger, which eventually leads to the increase in plasticity of both alloys with the increase in tensile temperature. This is consistent with the elongation and fracture morphology of the alloy after high temperature tensile testing.

4.2. Microstructure Evolution and Thermal Stability During Extended Exposure

Based on the results of the heat exposure tests at different high temperatures, it can be seen that the Er-free and Er-added alloys show a tendency to decrease in tensile strength after heat exposure, and the higher the heat exposure temperature, the greater the loss of tensile strength, as depicted in Figure 3. The decrease in tensile properties of the alloys after thermal exposure is closely related to the microstructure evolution of the precipitated phase during thermal exposure. For the age-hardening Al-Zn-Mg-Cu alloy, the tensile strength of the two alloys involved in this study is the result of the accumulation of multiple strengthening mechanisms. The Er-free and Er-added alloys used in this study are age-hardening alloys, and the yield strength (σ_y) can be calculated by Equation (1) (as suggested by refs. [7,17,24]), as shown below.

$${\sigma }_y={\sigma }_{gb}+{\sigma }_p+{\sigma }_{ss}+{\tau }_D$$

(1)

where σ_gb, σ_p, σ_ss, and τ_D represent grain boundary strengthening, precipitation strengthening, solid solution strengthening, and dislocation strengthening, respectively. Obviously, the strengthening effect of the above microstructure factors varies with the processing conditions of the alloy. In present work, for both alloys after thermal exposure, the strengthening mechanism is mainly σ_gb, σ_p, and σ_ss.

The contribution of σ_gb to yield strength follows the Hall–Petch relationship, as described in Equation (2).

$${\sigma }_{gb}={\sigma }_0+k{d}^{-0.5}$$

(2)

where σ₀ is the inherent frictional stress for pure aluminum and taken as 10 MPa; and k denotes the Hall–Petch coefficient, which is about 0.12 MPa/m^0.5 for the Er-free and Er-added alloys [25,26]. d represents the average grain size, which is 18.4 μm and 16.0 μm for the Er-free and Er-added alloys, respectively (as shown in Figure 10). σ_gb of the Er-free and Er-added alloys is 45 MPa and 48 MPa, respectively, calculated according to Equation (2).

For Al-Zn-Mg-Cu alloys, strength can be significantly improved by effectively hindering the movement of dislocations by various types of precipitates. The contribution of precipitation strengthening to the tensile properties of the alloy is closely related to the microscopic characteristics of the precipitated phase, and the differences in parameters such as the type, radius, number density, and volume fraction of the precipitates will affect the strengthening effect. In this study, the alloys were subjected to thermal exposure, which is equivalent to long-period aging treatment at high temperatures. In the Er-free and Er-added alloys, whether the dislocation cuts or bypasses the precipitated phase mainly depends on the size of the precipitated phase, and ultimately leads to different precipitation strengthening mechanisms. When the size of the precipitates is smaller than the critical size, the precipitates are easily cut by dislocations, and the improvement of the strength of the alloy is mainly attributed to the new surface energy generated by the precipitates [27,28]. If the size of the precipitated phase is larger than the critical size, the precipitated phase is difficult to be cut by the dislocation, and the dislocation can only bypass the precipitate. At this time, greater stresses are required, ultimately leading to increased strength of the alloy. Previous studies have shown that the critical size of the interaction between dislocations and precipitates from the cutting mechanism to the bypass mechanism is 3 nm [29,30]. For aluminum alloys, the contribution of precipitation strengthening to yield strength can be calculated by the following Equation (3):

$${\sigma }_p={\sigma }_{cut}+{\sigma }_{pass}$$

(3)

The contribution of GP zones and fine η′ to yield strength in Al-Zn-Mg-Cu alloys can be calculated by Equation (4) (refer to the method described in the refs. [23,31,32]).

$${\sigma }_{cut}=362.7f^{1/2}{r}^{1/2}$$

(4)

The yield strength increment provided by dislocation bypassing the precipitated phase can be expressed by Equation (5).

$${\sigma }_{pass}={\displaystyle \frac{0.4}{\pi }}{\displaystyle \frac{MGb}{\sqrt{1-v}}}{\displaystyle \frac{\mathit{ln}\left(4r/b\right)}{\lambda }}$$

(5)

where M is the mean orientation factor (3.06 for fcc material). G is the shear modulus (3.06 for Al). b is the Burgers vector of Al, which is 0.286 nm. r is the average radius of the precipitated phase with a size larger than 3 nm. υ represents Poisson’s ratio, which is 0.33 for Al. λ refers to the average free path between the precipitates.

For the Er-free and Er-added alloys after thermal exposure at 150 °C for 600 h, the radius of more than 90% of the precipitates is larger than 3 nm, as shown in Figure 12a and Table 2. Moreover, the main precipitates are η′ and η (as shown in Figure 10a and Figure 11a), and these coarse precipitates are difficult to be cut by dislocations. Consequently, the main strengthening mechanism for the Er-free alloy is the dislocation bypass mechanism, which accounts for 90.2%, and the contribution of the dislocation bypass mechanism to precipitation strengthening is relatively weak. The interaction mechanism between dislocations and precipitates in the Er-added alloy exhibits a similar performance to that of the Er-free alloy.

When the heat exposure temperature is lower, the precipitation behavior regains the precipitation driving force, and secondary precipitation occurs in the alloy [33]. Different from thermal exposure at lower temperatures, when the thermal exposure temperature is increased to 165 °C or 180 °C, the precipitates will be redissolved or coarsened. Whether the precipitated phase dissolves or coarsens mainly depends on the radius of the precipitated phase. When the radius of the precipitated phase is less than the critical radius, the precipitated phase will dissolve, and when it is greater than the critical radius, coarsening will occur [34]. As described in Figure 10b,e and Figure 11b,e, the size of the precipitates and the spacing between the adjacent precipitates are increased when the alloys are subjected to thermal exposure at 150 °C for 600 h. The average radius of the precipitated phase in the Er-free and Er-added alloys is coarsened to 19.63 μm and 16.53 μm, respectively. This indicates that the increase in the strength of the alloy relative to the precipitation contributes to the dislocation bypass mechanism. By further increasing the thermal exposure temperature to 180 °C, the existing η′ phase and η phase will continue to coarsen during this process, which can be confirmed by Figure 10c,f and Figure 11c,f. There is no doubt that the precipitation strengthening is attributed to the dislocation bypassing the precipitated phase. The overall comparison of the strength loss of the two alloys during long-period thermal exposure at three temperatures shows that the loss of strength at thermal exposure temperature is much greater than that of the thermal exposure time. In addition, the coarsening rate of the precipitated phase of the Er-added alloy is always lower than that of the Er-free alloy.

In summary, after T76 aging treatment, the precipitation strengthening mechanism of the Er-free and Er-added alloys includes the dislocation cutting mechanism and the Orowan bypass mechanism. When the thermal exposure time extended from 100 h to 600 h, or the thermal exposure temperature increased from 150 °C to 180 °C, the contribution of the Orowan bypass mechanism plays an increasingly dominant role for influencing the overall precipitation strengthening compared with the shearing mechanism. In order to quantitatively and intuitively analyze the effect of Er addition on the heat resistance of an alloy, alloys after thermal exposure to 150 °C, 165 °C, and 180 °C for 600 h were selected to verify the reliability of the yield strength prediction model. The characteristic parameters of the precipitated phases in the alloys subjected to different heat exposure treatments are listed in Table 2, and the calculated increments for yield strength of the alloys subjected to various heat exposure conditions are shown in Figure 13. During thermal exposure, the precipitates in the Er-added alloy exhibited excellent heat resistance with lower coarsening compared to the Er-free alloy. Consequently, the Er-added alloy shows significantly lower strength loss after thermal exposure than the Er-free alloy.

5. Conclusions

In this paper, the effects of trace Er addition on the microstructure and heat resistance of the Al-Zn-Mg-Cu alloy were quantitatively investigated and analyzed. The main conclusions are summarized as follows:

1.

The high temperature tensile properties and thermal stability of the alloy were effectively improved by adding trace Er.

2.

The Al₃Er particles will precipitate after homogenization treatment. It was observed that the Al₈Cu₄Er phase is uniformly distributed along the grain boundary in the Er-added alloy, which can effectively pin the grain boundary when the alloy is stretched at high temperatures.

3.

The precipitates in both alloys coarsen with the increase in thermal exposure time or temperature, which subsequently reduces the tensile properties of the alloys. Compared with the thermal exposure time, the thermal exposure temperature has a greater influence on the thermal stability. During long-period thermal exposure, the metastable η′ phase transforms to coherent η phase in both alloys.