The 7xxx series aluminum alloys are very attractive materials to be used in the automotive and aerospace industries, due to their excellent combination of properties, such as high strength-to-density ratio, high fracture toughness and resistance to stress corrosion cracking [1]. A good understanding of the hot deformation behavior and microstructural evolution is of primary importance for the design of hot-forming processes, such as rolling, extrusion and forging. Thermomechanical factors, such as the degree of deformation, deformation temperature and strain rate, are the main factors that influence the flow stress and the associated microstructure [2,3]. Furthermore, the evolution of flow stress is correlated with different dynamic softening mechanisms during hot deformation at various deformation conditions [2,4].

Dynamic recovery (DRV) and dynamic recrystallization (DRX) are the typical softening mechanisms in metals and alloys during deformation at elevated temperatures [2,4]. Aluminum and its alloys with high stacking fault energy exhibit a high rate of DRV, which significantly inhibits DRX [2]. However, the formation of new grains during hot deformation of aluminum alloys was frequently reported, while several mechanisms of DRX for aluminum alloys have been proposed [5–13]. The discontinuous dynamic recrystallization (DDRX) was observed in aluminum and aluminum alloys by many researchers [5–9]. The DDRX involves the development of high angle grain boundaries via nucleation and growth of new grains, which typically initiates at high angle boundaries, such as original grain boundaries, boundaries of dynamically recrystallized grains and boundaries created during deformation [4]. The bulging of grain boundaries is frequently observed as a prelude to DDRX, and subsequent grain growth is processed by strain-induced boundary migration [4]. Furthermore, continuous dynamic recrystallization (CDRX) was also observed in aluminum alloys [10–13]. This mechanism differs from that of DDRX, in which new grains are formed progressively within the deformed original grains from a continuous increase of subgrain boundary misorientations, as a result of the accumulation of dislocations in low angle boundaries [2,10–13].

Moreover, during the plastic deformation process, most of the deformation work is converted to heat [14]. Plastic deformation performed at high strain rate conditions is essentially adiabatic, thus generating a significant deformation heat and leading to a noticeable increase in temperature [15]. The deformation heating may cause the thermal softening of materials, which offsets the work hardening effect and result in a decrease in flow stress [16–18]. In addition, the formation of adiabatic shear bands is usually observed at regions where deformation is localized [15]. Furthermore, precipitation hardening alloys may alter dynamic softening behavior as a result of changes in precipitate morphology, when the material is not thoroughly over-aged. Materials pre-aged at deformation temperature have a low peak stress and exhibit a gradual softening as precipitates coalesce and subgrains form [19,20]. Solution-treated aluminum alloys exhibit a high peak stress due to dynamic precipitation (DPN), followed by a rapid dynamic softening as precipitates coalesce, solute depletes and DRV progresses [17,21–24].

To date, the previous studies have been mainly on determining the Zener–Hollomon parameter (Z), the activation energy for hot deformation and various material constants in the constitutive equations for describing the hot deformation behavior of 7xxx aluminum alloys [25,26]. The softening mechanism of these alloys is generally reported to be DRV at low temperature with a high strain rate, while DRX occurs as deformation at a high temperature with a low strain rate [27–29]. Different dynamic recrystallization mechanisms of 7075-T6 aluminum alloy have been studied during deformation at a semi-solid temperature range (450–580 °C) [30]. Several researchers have studied the kinetics and rate of DRX and DRV using various models [31,32]. Jonas et al. [31] measured the kinetics of DRX using the Avrami formulation and predicted the DRX flow curves. Mostafaei and Kazeminezhad [32] proposed a mathematical method based on hot flow curves to predict the effects of temperature and strain rate on the kinetics of DRV in the Al-Mg alloy. However, quantitative characterization of the evolution of flow stress curves at different deformation conditions has rarely been reported. The relationship between the evolution of flow stress and various dynamic softening mechanisms during hot deformation needs to be clarified. During deformation at a medium temperature range (up to 450 °C), a systematical investigation of different dynamic softening mechanisms at various deformation conditions for 7150 aluminum alloy is limited.

In the present paper, the hot deformation behavior of an Al-Zn-Mg-Cu alloy (7150) is studied using uniaxial compression tests performed at various temperatures and strain rates. The microstructural evolution during the hot deformation is investigated in order to understand the various dynamic softening mechanisms at different deformation conditions. Based on the different decline levels in flow stresses with respect to peak stresses, a decline ratio map of flow stresses is proposed to correlate flow stress behavior with dynamic softening mechanisms at various deformation conditions.

The experiments were conducted on an Al-Zn-Mg-Cu alloy. The chemical composition of the experimental alloy analyzed by an optical emission spectrometer is given in Table 1. Most of the elements of the alloy are in the range of 7150 composition designation. However, no Zr is added in this study. Approximately 3 kg of materials were melted in an electrical resistance furnace and then cast into a rectangular permanent steel mold measuring 30 × 40 × 80 mm³. The cast ingots were homogenized at 465 °C for 24 h, followed by a direct water quenching. Cylindrical samples 10 mm in diameter and 15 mm long were machined from the homogenized ingots. Uniaxial compression tests were conducted on a Gleeble 3800 thermomechanical simulation unit at strain rates of 0.001, 0.01, 0.1, 1 and 10 s⁻¹ and deformation temperatures of 300, 350, 400 and 450 °C, respectively. During the tests on the Gleeble unit, the samples were heated to the desirable deformation temperature at a heating rate of 10 °C/s and held for 3 min to ensure a homogenous temperature distribution through the samples. The samples were deformed to a total true strain of 0.8 and then immediately water-quenched to maintain the microstructure at the deformation temperature. Some selected samples were also deformed to different true strains of 0.1, 0.3 and 0.5, followed by water quenching, in order to investigate the microstructural evolution during the deformation process.

The microstructures of the as-cast and as-homogenized materials were examined prior to hot deformation under an optical microscope. The intermetallic phases were identified using the electron backscattered diffraction (EBSD) technique and the energy dispersive spectroscope (EDS) with a scanning electron microscope (SEM, JEOL JSM-6480LV, JEOL Ltd., Tokyo, Japan). All deformed samples were sectioned parallel to the compression axis along the centerline and then polished and etched in Keller solution for optical microscope observation. Additionally, some deformed and as-homogenized samples were selected for EBSD analysis. In the EBSD analysis, the boundaries of both the grains and subgrains are defined as low angle boundaries (LABs), medium angle boundaries (MABs) and high angle boundaries (HABs) for which the misorientation angles of boundaries occur in the ranges of 1°−5°, 5°−15° and greater than 15°, respectively [33]. The step size between the scanning points was set to 1.0 μm for the grain structure of the samples. The dynamic recrystallized grains were separated from the deformed ones using the grain average misorientation (GAM) method [34,35]. A threshold value of GAM (2.5°) inside a grain was determined to distinguish a recrystallized grain from a deformed one. Grains with GAM less than 2.5° were considered to be recrystallized grains. Subsequently, the size of the recrystallized grains was measured by means of a grain reconstruction method [36]. For quantitative measurement of the misorientation distribution of boundaries, am EBSD line scan was carried out [36,37], and a sample surface area of about 1.2 mm² with a scanning step size of 3.0 μm was selected. In addition, EBSD analysis was performed to measure the grain size of the as-homogenized sample using the linear intercept method [36]; a surface area of approximately 10 mm² with a scanning step size of 5.0 μm was selected for the sample. Samples for TEM observation were mechanically ground to the thicknesses of 35–60 μm followed by electropolish in a twin-jet polishing unit, which was operated at 15 V and −20 °C using a 30% nitric acid and 70% methanol solution. The disks were observed in a transmission electron microscope (TEM, JEOL-JEM-2100), operated at 200 kV.

The decline ratio map of flow stress as a function of temperature and strain rate is divided into five domains, which can be utilized to study the relationship between flow stress behavior and various dynamic softening mechanisms in the Al-Zn-Mg-Cu alloy (7150) during hot deformation at different deformation temperatures and strain rates.

In Domain I, the microstructure observation confirms that DRV is the main dynamic softening mechanism during hot deformation (Figure 5). Although deformed at high strain rates, no cracking was observed in the deformed samples. The continuous decline of the flow stress after attaining the peak stress is likely associated with a severe deformation heat release. The variations of sample temperatures are shown in Figure 10, which were monitored during the compression tests at the target temperatures from 300 to 450 °C with different strain rates. When the deformation was performed at the high strain rate of 10 s⁻¹ (Domain I), the sample temperature rose continuously with increasing strain at each target deformation temperature. The maximum overheating temperature reached 22 °C at the target temperature of 450 °C and increased as high as 35 °C at the target temperature of 300 °C. Those results suggest that the heat release during deformation at the high strain rate of 10 s⁻¹ plays an important role in the decrease of flow stress, leading to a significant thermal softening, as confirmed by other researchers [16–18].

However, in other domains, when the deformation was carried out at low strain rates between 0.001 s⁻¹ and 0.1 s⁻¹, the sample temperatures kept close to the target temperatures (within 2 °C). Thus, it is reasonable to consider that the deformation was under an isothermal process. Even at the intermediate strain rate of 1 s⁻¹, the temperature increase was generally limited within several degrees (4 to 8 °C). Therefore, it is believed that the thermal softening effect in this case plays a minor role in the flow softening.

Domain II represents a typical recovered structure (Figure 6a,b), indicating that DRV is the sole dynamic softening mechanism. The flow stress curves showed a continuous increase of flow stress after the peak stress at a higher strain rate (1 s⁻¹) or a slight decrease at a lower strain rate (0.1 s⁻¹). It is evident that the effect of work hardening was dominant during hot deformation at the strain rate of 1 s⁻¹, whereas the dynamic softening slightly overcame the work hardening at the lower strain rate of 0.1 s⁻¹.

In Domain III, DRV still operated as the main dynamic softening mechanism (Figure 6c,d). Differentiated from Domain II, a significant decline of the flow stress after the peak stress occurred. It was observed that the dynamic precipitation and coarsening took place during the hot compression process. Figure 11a shows a precipitate-free aluminum matrix after the homogenization prior to the hot deformation. At the beginning of deformation (at a true strain of 0.1), as exemplified at the deformation condition of 350 °C and 0.001s⁻¹, a large number of spherical and rod-shaped precipitates with an average size of 60 nm appeared (Figure 11b). The TEM-EDS result (Table 2) shows that the precipitates contain Mg, Zn and Cu, with a composition approaching the stoichiometric Mg(Zn,Cu)₂ phase [47–49]. Those fine precipitates interacted with dislocations and exhibited a strong pinning effect on dislocation movement, leading to a high value of the peak stress at the initial stage of deformation. As the deformation progressed to a true strain of 0.8, the precipitates remarkably coarsened to an average size of 125 nm with a larger interparticle spacing (Figure 11c). The coalescence of precipitates, accompanied by enhanced DRV, resulted in a significant flow stress softening. Similar behaviors have been observed in solution treated 7012, 7075 and 7085, as well as homogenized 2026 aluminum alloys during hot deformation at low strain rates [17,22,24].

It should be mentioned that dynamic precipitation can also occur in Domain II. However, the coalescence of precipitates was not observed by TEM, due to a relatively short deformation time at high strain rates. Those fine precipitates significantly increased the multiplication rate of dislocations [2,4], which resulted in a strong work hardening effect and a continuous increase of flow stress at a strain rate of 1 s⁻¹, as shown in Figure 2.

Moreover, in Domain IV, the deformed microstructure reveals a partially recrystallized microstructure at the serrated grain boundaries (Figure 7), which probably involves the discontinuous dynamic recrystallization (DDRX) by strain-induced boundary migration (SIBM). Figure 12 illustrates an example of the formation of one dynamically recrystallized grain before pinching off from the original grain in the sample deformed at 450 °C and 0.01 s⁻¹. The scanning transmission electron microscope (STEM) micrograph gives an overview of the recrystallized grain and its surroundings, and four bright-field TEM images show the enlarged views in the corresponding areas. The misorientation angles of the boundary were determined by measuring the shift of the Kikuchi line intersections on the crossing boundary under the TEM mode [2,23].

It can be seen that the recrystallized grain C is comprised of a partially bulged original grain boundary (Figure 12a,c) and a bridging of the low angle boundary (3°–5°) with grain B (Figure 12a,b). This suggests a characteristic feature of SIBM, by which the grain C was gradually formed at the expense of the higher density of dislocations in grain A and displayed a similar orientation to its parent grain, B. Moreover, grain C contains substructures of dislocations and dislocation networks (Figure 12a,d,e). This provides direct evidence that the new grain was formed dynamically during deformation and suffered a significant strain [2,7]. Furthermore, it was reported that a high content of solute atoms, Mg, Zn and Cu in the aluminum matrix of the homogenized sample, could reduce the dislocation mobility and retard dynamic recovery, thus leading to a high stored energy [2,4,12]. Therefore, the driving force for dynamic recrystallization was increased and DDRX took place during the hot deformation process at this deformation condition, resulting in an appreciable flow softening [2,12].

During the deformation at a high temperature of 450 °C and at a low strain rate of 0.001 s⁻¹ in Domain V, the bulging of grain boundaries and the formation of recrystallized grains by DDRX were not observed in Figure 8. As a high level of DRV was induced during the slow straining process at the high temperature, the accumulation rate of the dislocations was considerably decreased. Hence, the driving force was not sufficient for nucleation and the growth of recrystallized grains by DDRX. However, a number of newly formed grains with high angle boundaries can be clearly seen in the microstructure (Figure 8), in which it most probably involved a continuous dynamic recrystallization (CDRX). Figure 13 shows the evolution of the substructure and misorientation distribution with increasing strain during the deformation at 450 °C and 0.001 s⁻¹. At a true strain of 0.3, subgrains were well developed inside the original grains (see the arrows in Figure 13a) and characterized mainly by low angle boundaries between one and 5° (Figure 13c). As the deformation was processed to a true strain of 0.5, subgrains were presented with boundaries of increased misorientations (see the arrows in Figure 13b). Besides, the fraction of subgrain boundaries with low misorientations (1°–5°) was significantly reduced, while the number of boundaries with higher misorientations increased (Figure 13d). When the true strain approached 0.8, recrystallized grains with high angle boundaries could be observed (Figure 8), and the fraction of high angle boundaries continuously increased, reaching more than 45% (Figure 13e). This result indicates that the low angle boundaries (1°–5°) were progressively converted into medium angle boundaries (5°–15°), which, in turn, transformed into high angle boundaries during the slow deformation process (Domain V). Such an evolution can be considered as a process of continuous dynamic recrystallization (CDRX) [10–12]. In this case, the flow stress curve attained a peak stress and then showed a slow decline upon further straining [50].

In brief, the decline ratio map of flow stress was utilized to study the relationship between flow stress behavior and various dynamic softening mechanisms in the Al-Zn-Mg-Cu alloy (7150) during hot deformation. Recovered structures are generally observed in the samples deformed in Domains I, II and III. In these cases, DRV is the sole dynamic softening mechanism during hot deformation, which is commonly reported in 7xxx aluminum alloys [2,22,46,51]. However, a partially recrystallized microstructure is presented in the deformed samples in Domains IV and V. It is proved using EBSD and TEM that DRX does take place during hot deformation at certain conditions in this alloy. Therefore, besides DRV, DRX plays an important role in flow stress softening at the high temperature. Furthermore, results reveal that two possible mechanisms of DRX operate during hot forming, in which DDRX is involved in Domain IV and CDRX is implied in Domain V. Finally, it should be emphasized that the generation of a decline ratio map of flow stress, accompanied by investigation on typical microstructures in each domain, could provide a convenient approach to study the dynamic softening mechanisms of materials through hot processing.

Recently, several works have been done to optimize the hot working processes of 7xxx aluminum alloys using the processing maps [52–55]. Correlating various deformation mechanisms suggested from the decline ratio map with those processing maps, it is possible to predict the hot workability of the Al-Zn-Mg-Cu aluminum alloy studied. The results obtained in this study show that Domains II, IV and V correspond to the safe processing domains, due to DRV and DRX operating as the dynamic softening mechanisms [42,52–55]. Furthermore, the optimum hot working parameters for the alloy are considered to be in Domains IV and V, i.e., at the temperature of 450 °C and strain rates of 0.001–0.1 s⁻¹, where DRX occurs during hot deformation, leading to softening and the reconstituting of the microstructure [42]. This is in agreement with the results obtained from 7xxx aluminum alloys by other researchers [52–55]. On the other hand, the deformation in Domain I should be avoided, due to the flow instability and the formation of adiabatic shear bands [42,54,55]. The coarse dynamic precipitates generated in Domain III may lead to the deep inter-granular corrosion and large areas of the denudation layer, which are detrimental to the properties of the deformed material [54].