2019 Volume 60 Issue 9 Pages 2041-2047
The flow softening behavior and deformation mechanism of AA7050 aluminium alloy are investigated though hot compressive tests using a Gleebe-1500 thermal simulator at strain rates of 0.01∼10 s−1 and temperatures of 300∼450°C. The results show that the maximal relative softening transforms from 350°C to 450°C with increasing strain rate. Based on the processing maps and microstructure characterizations, dynamic recovery, dynamic precipitation and coarsening may be responsible for the softening behavior of AA7050 aluminium alloy at 350°C and 0.01 s−1. With increasing strain rate and temperature, high level of the flow softening is associated with deformation heating, dynamic recovery and dynamic recrystallization. The superimposed processing map with various strains is established, indicating that the optimum hot workability domain is the temperature range of 390∼450°C and strain rate range of 0.1∼0.01 s−1.
Fig. 6 Microstructures of AA7050 aluminum alloy under different deformation conditions: (a) 350°C, 10 s−1, (b) 400°C, 10 s−1, (c) 350°C, 0.01 s−1, (d) 450°C, 0.01 s−1.
AA7050 aluminum alloy belonged to 7000 (Al–Zn–Mg–Cu) series alloys was firstly introduced by Aluminum Company of America (ALCOA) in twenty century as the high strength thick plate alloy. It is widely used in the aircraft due to excellent service performance and good processing properties.1) AA7050 aluminum alloy is highly alloyed, and hot deformation is one of necessary procedures. In the past several decades, the hot forming behaviors based on flow softening, microstructure evolution and processing map have been researched for optimizing work parameters.2–5) It is generally known that the flow softening behavior is related to the thermal softening and microstructural softening, which are considered to be originated by many factors: deformation heating, strain rate, strain localization, the extent of dynamic recrystallization.6–8) Researches in 6063, 7020 and Al–Mg alloy have found the deformation heating can give rise to flow softening due to the increase in temperature.9–11) Luo12) showed that flow softening behavior is directly related to the occurrence of flow instability or cracking. Yang13) considered that deformation mechanism caused different flow softening behavior during hot deformation in AA7085 aluminum alloy. However, very few study have considered the extent and nature of the flow softening behavior from a microstructure point of view.
In the present work, the comprehensive effects of deformation temperature, strain rate and strain on flow softening behavior and deformation mechanism of AA7050 aluminum alloy during isothermal compression are investigated by the relative softening, processing map and microstructure evolution. The aim is to better understand the extent and nature of the dynamic softening, which is relevant to the industrial hot rolling processes.
The experimental material in this study is AA7050 aluminum alloy with chemical composition of 6.00%Zn, 2.20%Mg, 2.24%Cu, 0.10%Zr, 0.25%Fe, 0.15%Si, 0.10%Ti, 0.20%Ni (in mass%). The ingot was homogenized at 475°C for 48 h followed by water quenching. Cylindrical samples with a size of Φ8 × 12 mm were machined from the homogenized plate. The compression tests were carried out on a Gleeble-1500 thermal simulation machine at temperatures range from 300°C to 450°C and strain rates from 0.01 s−1 to 10 s−1. The samples were heated to the preset temperature at a heating rate of 5°C/s and held for 3 min prior to compression. Then, the samples were compressed to the required strain levels, and instantly quenched into cold water to preserve deformation microstructure.
The deformed samples were sectioned parallel to the loading axis. Microstructure observations were carried out by electron backscatter diffraction (EBSD) and transmission electron microscopy (TEM). The preparation of the EBSD samples was mechanically polished, followed by electrolytic polishing with a solution of 10% HClO4+90% C2H5OH. 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. The preparation of TEM was 50∼60 µm foil and then twin-jet polished in a 30% HNO3 and 70% CH4O solution. TEM examination was performed using a FEI TECNAL G2 F20 TEM microscope operated at 200 kV.
The typical true stress-true strain curves from hot compression tests are presented in Fig. 1. It can be seen that the flow stress increases with increasing strain rate and decreases with increasing temperature. At low strain rates ($ \leq 1$ s−1), the flow stress increases sharply and reaches the peak value, and then decreases continuously to a steady state. This suggest that dynamic softening could counteract the work hardening. Dynamic recovery is the main the softening mechanism.14) With increasing strain rate ($ \geq 10$ s−1), the flow stress after the peak value demonstrates a continuous softening behavior. It is noted that all the stress-stain curves at 450°C reveal a single peak at low strain level, implying the occurrence of the dynamic recrystallization originated preferentially from triple junctions and boundaries.13) As the strain increases, the dynamic softening is sufficient to counteract or surpass the work hardening during hot deformation, resulting in a steady or a continuous flow softening behavior.
True stress-strain curves of AA7050 aluminum alloy under different deformation conditions: (a) 0.01 s−1, (b) 0.1 s−1, (c) 1 s−1 and (d) 10 s−1.
In order to investigate the effect of deformation condition on the softening mechanism, the relative softening can be expressed:2)
| \begin{equation} S_{r} = \frac{\sigma_{p} - \sigma_{i}}{\sigma_{p}}\times 100\% \end{equation} | (1) |
Relative softening of AA7050 aluminum alloy under different conditions: (a) 0.01 s−1, (b) 0.1 s−1, (c) 1 s−1 and (d) 10 s−1.
In general, the relationship between flow stress and deformation parameter can be expressed as:15,16)
| \begin{equation} Z = \dot{\varepsilon}\exp[Q/RT] = A[\sinh(\alpha\sigma)]^{n} \end{equation} | (2) |
| \begin{equation} \left\{ \begin{array}{l} a = 0.9591\varepsilon^{5} - 1.7087\varepsilon^{4} + 0.9879\varepsilon^{3} - 0.1835\varepsilon^{2} \\ \qquad - 0.0073\varepsilon + 0.0230\\ n = -263.9033\varepsilon^{5} + 519.7577\varepsilon^{4} - 383.5277\varepsilon^{3} \\ \qquad + 136.5046\varepsilon^{2} - 23.8418\varepsilon + 6.1491\\ Q = -27751.8183\varepsilon^{5} + 56761.0175\varepsilon^{4} - 44585.0667\varepsilon^{3}\\ \qquad + 16862.8599\varepsilon^{2} - 3089.9409\varepsilon + 440.4428\\ \mathit{ln}\,A = -4763.1857\varepsilon^{5} + 9721.0974\varepsilon^{4} - 7571.9380\varepsilon^{3}\\ \qquad\quad + 2823.9367\varepsilon^{2} - 507.6300\varepsilon + 69.4120 \end{array} \right. \end{equation} | (3) |
Relationships between (a) $\ln \dot{\varepsilon }$ and σ, (b) $\ln \dot{\varepsilon }$ and ln σ, (c) $\ln \dot{\varepsilon }$ and ln sinh(ασ) and (d) ln sinh(ασ) and 1/T.
Variation of (a) α, (b) n, (c) Q and (d) ln A with strain.
As can be seen in Fig. 4, the value of α decreases at low strain levels (ε ≤ 0.3), whereas it increases firstly and then decreases at high strain levels ($\varepsilon \geq 0.3$). This suggests that the value of α is altered for strain. Figure 4(b) shows that the value of n decreases firstly and then increases after reaching its minimum value (4.50) at a strain of 0.4, indicating better workability at the intermediate strain level. Q and ln A exhibit the same tendency and they reach the minimum values of 217.33 kJ/mol and 33.50409 at a strain of 0.4, which reveal a decrease to deformation resistance. The value of Q varies in the range of 217.33∼260.92 kJ/mol, which are close to those reported aluminium alloys such as 229.75 kJ/mol in AA7150 aluminum alloy, 246.16 kJ/mol in 7075 aluminum alloy and 223 kJ/mol in furnace-cooled 7050 aluminum alloy.2,17,18)
3.3 Processing map and microstructure examinationProcessing map, which were based on the dynamic materials model (DMM), has been proven to be a useful tool for optimizing hot working parameters and predicting deformation mechanisms for metals and their alloys.19–21) The processing map can be described by an instability map combining with a power dissipation map. In the DMM, the work-piece was regarded as an energy dissipation system. The total input power P can be described as follows:
| \begin{equation} \mathrm{P} = \mathrm{G} + \mathrm{J} = \sigma\dot{\varepsilon} = \int_{0}^{\sigma}{\dot{\varepsilon}}\mathrm{d}\sigma + \int_{0}^{\dot{\varepsilon}}\sigma d\dot{\varepsilon} \end{equation} | (4) |
The strain rate sensitivity (m) of flow stress can be expressed as:
| \begin{equation} m = \frac{\partial\mathrm{J}}{\partial\mathrm{G}} = \frac{\dot{\varepsilon}\partial\sigma}{\sigma\partial\dot{\varepsilon}} = \frac{\partial\ln\sigma}{\partial\ln\dot{\varepsilon}} \end{equation} | (5) |
For an ideal linear dissipater, m = 1. J reaches its maximum of Jmax = P/2. Under this case, the efficiency of power dissipation, η, which is a function of temperature and strain rate, and defined as following:
| \begin{equation} \eta = \frac{\mathrm{2m}}{\mathrm{m} + 1} \end{equation} | (6) |
The variation of the efficiency of power dissipation represents a power dissipation map, which can be described by iso-efficiency contour map. The instability map is developed based on the dimensionless parameter, which is described by the highlighted region. The instability parameter can be expressed as:
| \begin{equation} \xi(\dot{\varepsilon}) = \frac{\partial\ln[\mathrm{m}/(\mathrm{m} + 1)]}{\partial\ln\dot{\varepsilon}} + m\leq 0 \end{equation} | (7) |
According to the eqs. (6) and (7), the processing maps of AA7050 aluminium alloy with various strains are presented in Fig. 5. It can be seen that an unsafe safe region is obtained in the strain rate range of 1∼10 s−1 and the temperature range of 300∼450°C (Fig. 5(a)). With increasing strain, another instability region appears at low strain rate and low temperature (Fig. 5(b)). At a strain of ε = 0.7, the instability regions invade progressively into each other, leading to the formation of new hot-working region. Such a evolution suggests that the degree of deformation has a significant impact on the workability of AA7050 aluminium alloy. However, the maximum value of η is 0.42, which lies in the stability region under the condition of 450°C/0.01 s−1 (Fig. 5(a)–(c)). Moreover, the peak value of η and it’s position have no obvious change during hot compression. Due to the instability of hot workability in AA7050 aluminium alloy, the superimposed processing maps at different strains are considered to choose the working parameters (Fig. 5(d)). As a consequence, the deformation parameters of the temperature at 390∼450°C and strain rate of 0.01∼0.1 s−1 are determined for optimum processing.
Hot processing maps of AA7050 aluminium alloy with various strains: (a) ε = 0.3, (b) ε = 0.5, (c) ε = 0.7 and (d) the superposition process maps at various strains.
The corresponding microstructures under different deformation conditions are presented in Fig. 6. Loading direction are indicated with green arrows. In these EBSD maps, boundaries with misorientations of 2∼5°, 5∼10°, 10∼15° and >15° are described by green lines, red lines, blue lines and black lines, respectively. Figure 6 shows that the original grains are elongated. At strain rate of 10 s−1, the majority of the low angle boundaries accompanied by deformation bands and shear bands marked with arrows are uniformly distributed in the interior of deformed grains (Fig. 6(a)). The width of adiabatic shear bands and deformation bands increases with increasing temperature. Deformation bands and adiabatic shear bands are delineated by banded-type structure aligned at approximately 0° and 20∼30° perpendicular to the compression direction, respectively (Fig. 6(b)). With decreasing strain rate, the adiabatic shear bands disappear, whereas deformation bands are still presented (Fig. 6(c)). At 450°C, new strain-free grains along original grain boundaries and the fragments of high-angle boundaries within deformation grains are observed, initiating the occurrence of dynamic recrystallization (Fig. 6(d)).
Microstructures of AA7050 aluminum alloy under different deformation conditions: (a) 350°C, 10 s−1, (b) 400°C, 10 s−1, (c) 350°C, 0.01 s−1, (d) 450°C, 0.01 s−1.
For Al–Zn–Mg–Cu alloy, the flow softening behavior during the high temperature deformation is established by the relative softening. In the present investigation, the peak relative softening of AA7050 aluminum alloy transforms from 350°C to 450°C when strain rate increases from 0.01 s−1 to 10 s−1. In an attempt to investigate the flow softening for the studied alloy in detail, the plots of the discrepancy between the preset temperature and the actual deformation temperature are obtained in Fig. 7. The increment of temperature reaches nearly 35°C compared to the pre-set temperature at the strain rate of 10 s−1, whereas it is neglected at low strain rates ($ \leq 1$ s−1). This suggests that the thermal softening occurs mainly due to deformation heating at high strain rate. Combining with the processing maps and the typical microstructures in the Fig. 5 and Fig. 6, adiabatic shear bands/deformation bands marked with arrows are observed in the instability domains. As investigated and analyzed by Luo and Cheng,12,22) deformation heating results in multiple slip systems are activated aligned along the lines of velocity discontinuity of the slip planes. Under the control of maximum shear stress, the directions of shear bands occur rotate away from the compression axis. In addition, the original grain orientation may be responsible for the microstructure characteristics within deformation grains. Taylor23) have shown that the S grain and Copper grain reveal microbands in the form of larger-angle boundaries aligned at 20∼50° to the rolling direction. However, due to the absence of the misorientation relationship during grains before deformation, microstructure depended on the grain orientation is difficult to explain the formation of deformation bands/shear bands in this study. Using TEM, microstructure evolution of the 450°C/10 s−1 sample exhibits high dense dislocations are accumulated in the vicinity of triple junction due to the specific geometric feature at a strain of ε = 0.3 (Fig. 8(a)). With increasing strain, dislocations are ejected from a grain boundary and enter dense dislocation walls (DDWs), resulting in a decrease in width and annihilation of DDWs (Fig. 8(b)).24) Thus, the high angle grain boundary is formed, as shown in Fig. 8(c). Such a feature implies the occurrence of continuous dynamic recrystallization. Therefore, complicated microstructural softening accompanied by thermal softening may be responsible for continuous softening at 450°C and 10 s−1.
The actual temperatures of AA7050 aluminium alloy during isothermal compression; (a) 300°C, (b) 350°C, (c) 400°C and (d) 450°C.
TEM micrographs of the samples are deformed to various strains at 450°C, 10 s−1: (a) ε = 0.3, (b) ε = 0.5, (c) ε = 0.7.
In contrast to the high strain rate, deformation heating is neglected at low strain rate. The flow softening is attributed to microstructural softening, consisting of dynamic recovery, dynamic recrystallization, dynamic precipitation.24,25) According to the results of Fig. 6 and Fig. 7, dynamic recovery is considered as a main softening mechanism in the 350°C/0.01 s−1 sample (Fig. 6(c)). However, the presence of deformation bands is benefit to work hardening due to intense dislocation activities.26) The relative softening of the 350°C/0.01 s−1 sample may be related to the fastest dynamic precipitation and coarsening. According to the TTP diagrams of 7050 alloy, the nose temperature is about 330°C.27) Therefore, the fastest transformation can occur around this temperature. The further illustration is identified by TEM. Figure 9(a) shows the majority of the fine particles accompanied by high density of dislocations. With increasing strain, the dislocation density decrease. Well-developed sub-grain structure is formed, implying that dynamic recovery is the main deformation mechanism. The number density of second phase particles increases from 2.1 × 107 to 4.3 × 107 (mm−2). The average size increases from 48 nm to 60 nm (Fig. 9(a)–(c)). As a consequence, the effect of second phase particles on dislocation become less and less. Therefore, the flow softening of the 350°C sample caused by dynamic recovery, dynamic precipitation and coarsening is higher than that of other ones at strain of 0.01 s−1.
TEM micrographs of the samples are deformed to various strains at 350°C, 0.01 s−1: (a) ε = 0.3, (b) ε = 0.5, (c) ε = 0.7.
The flow softening behaviors during hot compression are attribute to many factors. In order to further clarify the differences of microstructures evolution under different deformation conditions, further studies by orientation dependence of the substructure would be necessary.
This study is supported by The Electronic Microscopy Center of Chongqing University of China and Chongqing Key of Extraordinary Bond Engineering and Laboratory Advanced Materials Technologies, Yangtze Normal University.
