Effect of Heating Rate on Microstructure, Texture and Tensile Properties of AA7085
2019 Volume 60 Issue 3 Pages 464-470
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2019 Volume 60 Issue 3 Pages 464-470
The microstructure, texture and tensile properties of a deformed AA7085 alloy during solution treatment were investigated in the temperature range of 350 to 500°C with heating rates of 1 and 600°C/min. The results show that the alloy sheet solution treated with rapid heating rate exhibits a finer grain structure and narrower grain size distribution compared to the alloy sheet solution treated with slow heating rate, indicating that driving force is more important than sub-boundaries coalescence and mobility during rapid heating process. Heating rate affects the qualitative recrystallization textures, whereas it has no significant changes on the quantitative texture of the individual components. After two-stage ageing, the samples with low heating have slightly higher tensile properties than those with high heating rate.
Fig. 3 Microstructure evolutions of a deformed AA7085 alloy which were heated to 350°C (a, e), 400°C (b, f), 450°C (c, g) and 500°C (d, h) with heating rates of 1 (a, b, c, d) and 600°C/min (e, f, g, h).
The Al–Zn–Mg–Cu alloys are key materials which have been widely used in the aircraft and automotive industries.1,2) The production processes of thick plate consist of casting, homogenization, preheat, deformation, solution treatment and ageing.3) Significant changes of microstructure are found during the solution treatment, including recrystallization of deformed structure, formation of recrystallization texture, dissolution and/or precipitation of second-phase particle. These changes have prominent influence on subsequent ageing stage and its final mechanical properties.4,5) Therefore, it is very necessary to consider the effect of solution treatment on recrystallization behavior and mechanical property.
Generally, recrystallization mechanism and properties of the solution treated alloy are well known to be controlled by temperature, holding time, strain level, initial microstructure and heating rate. Recent research in AA 5083 alloy has attributed Lüders elongation phenomenon and serration amplitude to the annealing temperature.6) O.N. Senkov showed that growth and coarsening of Al3(Sc,Zr) particle occurs during the isothermal holding at 480°C for 4 h.7) Liu et al. revealed that the cold-rolled Al–0.8Si alloy exhibit finer grain size than that of the cold-rolled Al–0.05Si alloy.8) However, less work has been done on the effect of heating rate on softening behavior of the alloy from a microstructure point of view, while recent reports have pointed out that heating rate plays a very important role on the recrystallization behavior in tantalum, molybdenum, steel and aluminium alloy.9–12)
Compared with the researches on high melting point materials, the extent and nature of the restoration behavior of Al–Zn–Mg–Cu alloys during the heating-up process of solution treatment with different heating rates are still not clearly understood. AA7085 alloy is widely used in the aeronautical and aerospace industries because of its high strength, good toughness and slow quench sensitivity.13) With the rapid development and wide application of large slabs, understanding the restoration behavior becomes important from an industrial perspective. In addition, the recrystallization behavior of a deformed AA7085 alloy is different from that of the other 7000 series aluminum alloys. The heat treatment of extruded AA7075 alloy and hot-rolled AA7050 alloy reveal that recovery and continuous recrystallization are the main restoration mechanisms.14,15) However, AA7085 alloy shows that initial deformation stored energy is essential factor for restoration behavior.16) The objective of this work is to reveal the microstructure, texture and mechanical properties evolution of a deformed AA7085 alloy during solution treatment with various heating rates, providing a reference for the optimization of heat treatment to obtain high-performance alloy materials.
A commercial AA7085 alloy is used in this study with the chemical composition (in mass%) of: 7.0–8.0%Zn, 1.2–1.8%Mg, 1.3–2.0%Cu, 0.08–0.15%Zr, Fe ≤ 0.08%, Si ≤ 0.06%. The casting ingot was homogenized at 400°C for 12 h and 460°C for 12 h, and then rapidly quenched into room-temperature water. After homogenization, the plate was subjected to 50% hot rolling and then cold-rolled to produce a 20% reduction in thickness.
The samples with nominal dimensions of 10 × 8 × 6 mm3 were machined and then divided into two groups for tests. One group of the samples were heated to the target temperatures between 350 and 500°C in air furnace with heating rate of 1°C/min and the others were conducted using a Gleeble-3500 thermal and mechanical simulator with heating rate of 600°C/min, followed by quenching in cold water immediately. After solution treatment and water quenching, the samples which were continuously heated up to 500°C with heating rate of 1°C/min and 600°C/min were subjected to two-stage ageing (110°C × 10 h + 180°C × 24 h) in the air furnace.
Microstructural analyses were performed using scanning electron microscope (SEM), electron backscatter diffraction (EBSD), transmission electron microscopy (TEM) and X-Ray diffraction (XRD) techniques. The preparations of the SEM and EBSD samples were electro-polished in a 10% HClO4 and 90% C2H5OH solution with current density of 1.5 A cm−2 for 30 s at −20°C. The EBSD test was conducted using a field emission gun-environmental scanning electron microscopy (FEG-SEM) FEI device equipped with an HKL Channel 5 EBSD System. Thin TEM foils were twin-jet electro-polished at −30°C in a 30% HNO3 and 70% CH4O solution. TEM examination was performed using a FEI TECNAL G2 F20 TEM microscope operated at 200 kV. XRD measurements were performed on a Rigaku X-ray diffractometer (D/MAX-3500) equipped with a Cu target at 6 kW. Texture measurements using standard X-ray diffraction were taken on the mid-thickness rolling plane of the samples. The (111), (200) and (220) pole figures were measured up to a maximum tilt angle of 70° by the Schulz back-reflection method using Cu Kα radiation. The orientation distribution functions (ODF) were calculated by using the conventional series expansion method. The tensile samples of 20 mm in length and 6 mm in width and 1 mm in thickness were machined from the rolling plane of 1/2 thickness along the normal direction (ND) surface. Tensile tests were taken on an AG-X testing machine with a strain speed of 0.006 s−1. Three to six samples per condition were tested.
Microstructures of a deformed AA7085 alloy are presented in Fig. 1. It can be seen that deformed grains show elongated shape. Coarse particles with the mean size of 3–18 µm are distributed along original grain boundaries, as marked with white arrows in Fig. 1(a). Using energy disperse spectroscopy (EDS) analysis, these particles are likely to be Al65Cu20Fe15, which are composed of Al, Cu, and Fe. From high magnification SEM micrograph, the relative fine particles with the average size of 0.25 µm are presented within deformed grains (Fig. 1(b)). A detailed TEM examination in Fig. 1(c)–(d) shows tangled dislocation accompanied by the majority of second phase particles within the deformed grains, illustrating that a deformed alloy sheet has high stored energy. The textures of the deformed sample display the typical rolling textures, which are composed of Brass-orientation {011}⟨211⟩, Copper-orientation {112}⟨111⟩, S-orientation {123}⟨634⟩, and Goss-orientation {011}⟨100⟩, as shown in Fig. 2.
Microstructures of a deformed AA7085 alloy: (a) low magnification SEM micrograph (b) high magnification SEM micrograph, (c) low magnification TEM micrograph and (d) high magnification TEM micrograph.
ODF map of a deformed AA7085 alloy.
The microstructure evolutions of a deformed AA7085 alloy during slow and rapid heating process are compared in details using EBSD technique, as shown in Fig. 3. The area fraction of recrystallization and recrystallized grain size are depicted in Fig. 4.
Microstructure evolutions of a deformed AA7085 alloy which were heated to 350°C (a, e), 400°C (b, f), 450°C (c, g) and 500°C (d, h) with heating rates of 1 (a, b, c, d) and 600°C/min (e, f, g, h).
Effect of heating rate on (a) fraction recrystallized and (b) recrystallized grain size and (c) grain size distribution of the samples which were heated to 500°C.
At 350°C, the presence of new recrystallized grains in the 1°C/min sample shows that recrystallization occurs predominantly along original grain boundaries, where coarse second-phase particles with the average size of 7–14 µm (>1 µm) (Fig. 3(a)), insinuating that recrystallization may be attributed to particle stimulated nuclei (PSN).17,18) However, the microstructure of the 600°C/min sample shows similar characteristic as the deformed sample (Fig. 3(e)), suggesting that the majority of the deformed grains are still at an early stage of recrystallization. By further heating to 400°C, some deformed grains of the 1°C/min sample are replaced by almost dislocation-free grains, which exhibit coarse elongated shape with the average size of 27 µm (Fig. 3(b) and Fig. 4(b)), whereas fine recrystallized grains with the average size of 10 µm are relatively homogeneous distributed with heating rate of 600°C/min (Fig. 3(f) and Fig. 4(b)). As the temperature increases (Fig. 3(c)–(d) and Fig. 3(g)–(h)), predominant recrystallization is developed. The samples with high heating rate show finer grain size and higher area fraction compared to the samples with low heating rate, as shown in Fig. 4. Furthermore, recrystallization grain size distributions of the 500°C samples suggest that rapid heating results in a fine grain size and narrow grain size distribution. The values of grain size variance are 25.1 and 6.6 with heating rates of 1 and 600°C/min, respectively (Fig. 4(c)). Such microstructure evolutions suggest that although the increase in heating rate has a profound effect on the recrystallization kinetics due to the absence of recovery, second-phase particles have a significant impact on recrystallization progress.
In general, recrystallization behavior is co-affected by recovery and second-phase particles in precipitation strengthened aluminum alloys. In this study, the presence of second phase particles in Fig. 1 suggests that precipitation occurs prior to recrystallization. During the heating-up process of solution treatment, effect of heating rate on second-phase particles is shown in Fig. 5. The diffraction intensities of MgZn2 particles become weak progressively, whereas the diffraction peaks of Al65Cu20Fe15 particles have no obvious change. This indicates that MgZn2 particles are dissolved into matrix with increasing temperature, while Al65Cu20Fe15 particles can’t dissolve. It is noted that the diffraction intensities of MgZn2 particles decrease obviously in the 1°C/min samples, whereas they are still found in the 600°C/min samples at 450°C, suggesting that high heating rate can retard the dissolution of MgZn2 particles. In the case of low heating rate, well-developed sub-grains and Al65Cu20Fe15 particles suggest that PSN may be responsible for initiation of recrystallization (Fig. 6 and Fig. 5). In addition, the dissolution of MgZn2 particles decrease grain boundary pinning effect, and enhances sub-grain boundary migration and growth (Fig. 6(b)–(c)). Under this case, recovery and precipitation before recrystallization can accelerate recrystallization.19,20) The early nucleation sites will consume gradually deformation matrix, resulting in a coarse elongated grain structure. Continuous recrystallization seem to become more significant due to sub-grain growth and coalescence (Fig. 6(b)–(c)). As the heating rate increases, high density tangled dislocations accompanied by the majority of the particles reveal that the release of stored energy and the dissolving of the MgZn2 particles are retard (Fig. 6(d)). These fine precipitates can retard motion of dislocations, and inhibit recovery at low temperatures (<400°C).21) As a consequence, stored energy is shifted to high temperature, and provides a large driving force for recrystallization. Under this case, discontinuous recrystallization can occur simultaneously in local regions of high stored energy (Fig. 6(e)). Recrystallized grains have the nearly identical growth opportunity. More prominent recrystallization and finer grain structure are found in the 600°C/min samples than that of the 1°C/min samples at high temperatures (≥450°C).
Effect of heating rate on second-phase particles: (a) 1°C/min and (b) 600°C/min.
TEM micrographs of a deformed AA7085 alloy which were heated to 350°C (a, d), 400°C (b, e) and 450°C (c, f) with heating rates of 1 (a, b, c) and 600°C/min (d, e, f).
The developments of texture in a deformed AA7085 alloy with heating rates of 1 and 600°C/min are shown in Fig. 7 and Fig. 8, respectively. The ODF sections clearly show that no new texture component is observed at 350°C (Fig. 7(a) and Fig. 8(a)). As the recrystallization processes, the presences of $\{ 001\} \langle 5\bar{1}0\rangle $ and $\{ 125\} \langle 5\bar{5}1\rangle $ texture components are confirmed from the respective ODF sections with heating rate of 1°C/min (Fig. 7(b)–(d)). In the case of high heating rate, no significant differences are observed (Fig. 8(b)–(c)). Recrystallization textures are composed of random textures and retaining rolling textures (Fig. 8(d)).
Recrystallization texture evolution of a deformed AA7085 alloy with heating rate of 1°C/min: (a) 350°C, (b) 400°C, (c) 450°C and (d) 500°C.
Recrystallization texture evolution of a deformed AA7085 alloy with heating rate of 600°C/min: (a) 350°C, (b) 400°C, (c) 450°C and (d) 500°C.
In order to further illustrate the effect of heating rate on texture evolution, the volume fractions of important texture components are shown in Fig. 9. By comparison of Fig. 9(a) and Fig. 9(b), the volume fractions of important texture components are not significantly affected by heating rate.
Effect of heating rate on the volume fractions of texture components with heating rates of: (a) 1°C/min and (b) 600°C/min.
Heating rate affects the qualitative recrystallization textures, which is attributed to nucleation sites and growth behavior.22–24) In the case of low heating rate, recrystallization undergoes a wide range of temperature. Oriented nucleation (ON) and oriented growth (OG) can affect the formation of recrystallization texture.25) The orientation map of the deformed alloy after annealing at 375°C shows profuse cube oriented grains (Fig. 10). ON developed preferentially from cube bands/sub-grains may be responsible for the formation of $\{ 001\} \langle 5\bar{1}0\rangle $ texture due to high mobility. As the temperature increases, the remaining deformed grains occur recrystallization, resulting in the formation of $\{ 125\} \langle 5\bar{5}1\rangle $ texture.
Orientation map of a deformed 7085 alloy which is heated to 375°C with heating rate of 1°C/min.
In the case of high heating rate, the distribution of recrystallized nucleation in the rapid heating samples is more homogeneous and random than that of the slow heating ones during the early stage of recrystallization (Fig. 3), resulting in a severe decrease in preferential nucleation and growth. Therefore, dominant recrystallization textures are characterized by random textures and weak rolling textures. However, the volume fractions of textures in the 1°C/min and 600°C/min samples appear very similar. It is apparent that the quantitative texture is related to the intensity and the orientation scattering of texture components. From the ODF sections, the intensities of cube and $\{ 125\} \langle 5\bar{5}1\rangle $ orientations in the 1°C/min samples are higher than that of the 600°C/min ones, whereas their scatterings are lower than that of the 600°C/min ones. Consequently, the intensities of the individual texture components can counteract their scatterings, resulting in the similar quantitative textures with different heating rates.
3.4 Effect of heating rate on the tensile propertiesThe tensile properties evolution of the 1 and the 600°C/min samples are presented in Fig. 11. It can be seen that the tensile strength decreases with increasing temperature, whereas the total elongation increases. 7085 alloy is a precipitation-hardenable aluminum alloy. The samples which are continuously heated up to 500°C with heating rate of 1°C/min and 600°C/min are subjected to two-stage ageing. The tensile tests reveal that the yield strength (YS) and ultimate tensile strength (UTS) of the 1°C/min sample are 493 MPa and 520 MPa, respectively, whereas they are decreased to 467 MPa and 495 MPa with heating rate of 600°C/min. This is not in accordance with the prediction of Hall-Petch relationship. In combination with Fig. 1 and Fig. 12, the coarsening of precipitation are observed with low heating rate. This result suggests that the coarsening of S precipitation during solution treatment can decrease the YS and UTS.26) However, the corresponding X-ray diffraction analysis shows that the majority of second phase particles are remained with heating rate of 600°C/min (Fig. 5(b)), resulting in low level of supersaturation. In order to further clarify the differences of precipitation under subsequent ageing stage, TEM micrographs exhibit the majority of second phase particles (Fig. 13). The 1°C/min treated sample shows discontinuous distribution of precipitations along the grain boundary, while the 600°C/min treated sample shows continuous distribution of precipitations and obvious precipitate free zone (PFZ). Consequently, the slow heating samples have slightly higher tensile properties than the rapid heating ones after two-stage ageing. Therefore, the final mechanical properties of AA7085 alloy are affected significantly by the solution extent and the subsequent ageing precipitation.
Effect of heating rate on the tensile properties: (a) 1°C/min and (b) 600°C/min.
The coarsening of precipitation at 350°C with heating rate of 1°C/min: (a) low magnification micrograph and (b) high magnification micrograph.
Typical TEM morphologies of grain boundary precipitates for the solution treated sheets with heating rates of: (a) 1°C/min and (b) 600°C/min.
This study is supported by The Electronic Microscopy Center of Chongqing University of China and Chongqing Key Laboratory of Extraordinary Bond Engineering and Advanced Materials Technologies.
