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Flow behavior and microstructure evolution of 7055 aluminum alloy impacted at high strain rates
Materials Science & Engineering A 677 (2016) 203–210
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Flow behavior and microstructure evolution of 7055 aluminum alloy impacted at high strain rates Shengdan Liu a,b,n, Shaoling Wang a,b, Lingying Ye a,b, Yunlai Deng a,b, Xinming Zhang a,b a b
School of Materials Science and Engineering, Central South University, Changsha 410083, PR China Key Laboratory of Nonferrous Metal Materials Science and Engineering, Ministry of Education, Changsha 410083, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 21 July 2016 Received in revised form 10 September 2016 Accepted 13 September 2016 Available online 14 September 2016
The flow behavior and microstructure evolution of 7055 aluminum alloy impacted at the strain rates from 1000 s 1 to 6000 s 1 were investigated by a compressive split-Hopkinson pressure bar, optical microscopy, transmission electron microscopy and scanning transmission electron microscopy. The maximum stress tends to increase first with the increase of strain rate from 1000 s 1 to 4000 s 1, and then decrease at higher strain rate. The strain rate sensitivity is positive in the strain rate range of 1000– 4000 s 1, but became negative above 4000 s 1 because of the presence of shear bands and shear band induced cracks due to intensive strain localization. Impact deformation prompts transformation of η′ phase into η2 phase, and changes the morphology of precipitates from a platelet-like shape to a globelike shape. & 2016 Elsevier B.V. All rights reserved.
Keywords: 7055 aluminum alloy High strain rate Microstructure
1. Introduction 7000 series aluminum alloys are well known precipitation hardenable alloys, and high strength and good ductility can be achieved after aging. Therefore, some alloys such as 7039 and 7020 have been used as armor materials for combat vehicles. In order to enhance the performance of combat vehicles, it is always desirable to use an alloy with higher specific strength and improved ballistic properties. 7055 aluminum alloy can have very high strength more than 600 MPa and simultaneously good ductility and fracture toughness after heat treatment [1–3]. Mondal and Mishra et al. [4] shown that peak-aged 7055 aluminum alloy exhibits better ballistic properties than 7017 aluminum alloy. This implies that this alloy is a good candidate for light armor applications. For this purpose, it is required to fully understand high strain rate behavior of this alloy. To date, there are few literature on this topic. Zhu and Pang [5] reported shear localization of 7055-T77 aluminum alloy under dynamic compression by split Hopkinson pressure bar (SHPB) test method. Mishra and Mondal [6] investigated plastic flow behavior of 7055 aluminum alloy under different high strain rate test methods. In these investigations, however, no attention was paid to precipitation structure in the impacted alloys. In the present study, the dynamic response of 7055 aluminum alloy plate was investigated using SHPB test method, and n Corresponding author at: School of Materials Science and Engineering, Central South University, Changsha 410083, PR China. E-mail addresses: [email protected], [email protected] (S. Liu).
http://dx.doi.org/10.1016/j.msea.2016.09.047 0921-5093/& 2016 Elsevier B.V. All rights reserved.
microstructural evolution was investigated by transmission electron microscopy (TEM) and scanning transmission electron microscopy (STEM).
2. Experimental Cylindrical specimens with a length of 4 mm in thickness and 6 mm in diameter were cut from 7055-T7 aluminum alloy plate for dynamic impact tests, which were performed at room temperature on a split Hopkinson pressure bar machine. The nominal strain rates were selected from 1000 s 1 to 6000 s 1. For each strain rate, three specimens were tested to obtain reliable results. The impact direction was parallel to the rolling direction of the plate. Uniaxial compression tests were also performed on a CRIMS-DDL 100 testing machine with a nominal strain rate of 8 10 4 s 1 at room temperature to obtain quasi-static true stress-strain curves for comparison. The chemical compositions of the alloy were (wt%): Al-7.87 Zn-2.16 Mg-2.05 Cu-0.12 Zr, the contents of Fe and Si were lower than 0.10%; and the tensile properties at room temperature are given in Table 1. Microstructure examination was performed on the specimens before and after impact tests. Metallographic samples were ground, polished and etched in a solution made up of 1 mL HFþ16 mL HNO3 þ3 g CrO3 þ83 mL H2O to reveal grain structure, and then examined on a MX3000 optical microscopy. Moreover, disc samples with a size of 3 mm in diameter and 0.1 mm in thickness were prepared and electropolished in a solution containing 80% ethanol and 20% nitric acid below 20 °C in liquid
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Table 1 Tensile properties of the as-received 7055-T7 aluminum alloy plate. Material
s0.2 (MPa)
sb (MPa)
δ (%)
7055-T7 aluminum alloy plate
586
605
11.6
nitrogen; they were examined on a Tecnai G2 F20 S-Twin scanning transmission electron microscopy (STEM) with a high angle annular dark field (HAADF) detector and probe corrector operated at 200 kV.
3. Results 3.1. Stress-strain curves Fig. 1 shows typical true stress-strain curves of 7055-T7 aluminum alloy under quasi-static compression at a nominal strain rate of 8 10 4 s 1. As seen from the figure, flow stress tends to increase rapidly with the increase of strain and then very slowly with strain larger than about 0.1; stress is about 700 MPa at the strain of 0.4. No softening was observed during quasi-static compression. Fig. 2 shows typical stress-strain curves of 7055-T7 Al alloy impacted at strain rates from 1000 s 1 to 6000 s 1. Strain rate has significant effect on the shape of stress-strain curves. At 1000 s 1, stress increases rapidly in the initial stage and then slowly with the further increase of strain; this trend is similar to that under quasi-static compression as shown in Fig. 1. At higher strain rates, 2000–6000 s 1, the shape of stress-strain curves is quite different; stress tends to increase rapidly to a yield point and then exhibit a significant drop with the increase of strain. This phenomenon is similar to that of some aluminum alloys, for instance, 2219 alloy impacted at high strain rates from 2000 s 1 to 3000 s 1 [7]. Most of impact energy (about 90%) may convert to thermal energy, and therefore increases temperature; consequently, flow softening occurs. With the further increase of strain, stress starts to rise, fluctuate and finally tends to be a constant value. This may indicate that strain hardening and thermal softening occurred during impacted deformation, and finally a balance was achieved. From Fig. 2, it can be seen that the maximum stress tends to increase with the increase of strain rate below 4000 s 1, but decrease at 6000 s 1. The maximum stress was about 650 MPa at the strain rate of 1000 s 1, 690 MPa at 2000 s 1, 790 MPa at 4000 s 1 and 710 MPa at 6000 s 1. In order to have better understanding of the effect of strain rate on the flow behavior during impact
Fig. 2. True stress-strain curves of 7055-T7 aluminum alloy impacted at different strain rates.
⋅
Fig. 3. Relationship between logs and log ε .
deformation, strain rate sensitivity, m, was calculated by Eq. (1) [8] ⋅
m = d log σ /d log ε
(1) .
Where σ is the stress and ε is the strain rate. The results are shown in Fig. 3. From 1000 s 1 to 4000 s 1, there is a positive strain rate sensitivity, which indicates that a higher strain rate would give rise to higher stress due to strain hardening; m was estimated to be 0.088. However, above 4000 s 1, the value of m tends to be negative, about 0.259. This phenomenon was also observed in some previous investigations. For instance, Oosterkamp et al. [8] found for a 7108 aluminum alloy at room temperature strain rate sensitivity was negative at strain rates larger than 2000 s 1. This likely results from the formation of shear bands due to intensive strain localization as shown in following section. 3.2. Microstructure
Fig. 1. True stress-strain curves of 7055-T7 aluminum alloy under quasi-static compression at a nominal strain rate of 8 10 4 s 1.
3.2.1. Optical micrographs Fig. 4 shows optical micrographs of the specimens before and after impact deformation. In the as-received specimen, there is a partially-recrystallized microstructure (Fig. 4(a)); the white regions are recrystallized grains, while the dark regions are unrecrystallized grains containing a number of subgrains. This kind of grain structure is typical in 7000 series aluminum alloy plate or sheet [9,10], and is primarily attributed to the presence and band distribution of Al3Zr dispersoids, which can inhibit recrystallization during solution heat treatment at high temperature [11,12]. In the specimen impacted at 1000 s 1, there is no noticeable change
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Fig. 4. Optical micrographs of 7055 aluminum alloy specimen (a) before impact and impacted at (b) 1000 s 1 (c) 2000 s 1 (d) 4000 s 1 (RD: rolling direction, ND: normal direction, ID: impact direction).
in the grain structure, see Fig. 4(b); this may imply that deformation was quite uniform in the specimens. At the strain rate of 2000 s 1, the grains in the impacted specimen started to be distorted and some deformed shear bands can be found (Fig. 4(c)), which indicates the occurrence of inhomogeneous deformation. In the specimens impacted at 4000 s 1, the grains were distorted more seriously, see Fig. 4(d), and deformed shear bands are clearly visible; this shows that inhomogeneous deformation was serious, leading to strain localization. In the specimen impacted at the highest strain rate of 6000 s 1, adiabatic shear bands (ASB) and ASB induced cracks are clearly visible (Fig. 5), and they are at an angle of about 45°to the impacted direction. Similar phenomenon was also found in a peak-aged 7055 Al alloy impacted at 6500 s 1 [6]. The formation of ASB is generally an indication of highly localized and inhomogeneous deformation, and primarily responsible for the drop in strength during deformation [8]. Therefore, the results in Figs. 4 and 5 may explain the decrease of the maximum stress with strain rate increasing from 4000 s 1 to 6000 s 1 shown in Fig. 2 and the negative strain rate sensitivity shown in Fig. 3.
Fig. 5. Optical micrographs of 7055 aluminum alloy specimen impacted at a strain rate of 6000 s 1 (RD: rolling direction, ND: normal direction).
3.2.2. TEM and STEM examination results The high strength of 7000 series aluminum alloy primarily results from a high density of fine η0 phase strengthening precipitates in the matrix [1,2,13,14]. Though some investigations have been carried out on dynamic impact properties of these alloys, little attention was paid to the change of precipitation structure. Figs. 6–11 show some typical TEM and STEM images of the specimens before and after impact deformation.
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Fig. 6. TEM bright field images of (a) the as-received specimen and (b)(c) the specimen impacted at 6000 s 1.
In the as-received specimen, there is a low density of dislocations inside grains, as shown by the TEM bright field image in Fig. 6(a); (sub) grain boundaries can be identified clearly. After impact deformation at 6000 s 1, a large number of dislocations were introduced, see Fig. 6(b), and it is not easy to see (sub) grain boundaries clearly. Higher magnification image shown in the interior of grains there are a number of dislocation walls, which contained some parallel dislocation lines (Fig. 6(c)); this can enhance strain hardening [15]. Fig. 7 shows TEM bright field images taken along o011 4 Al and o001 4 Al direction and the corresponding selected area diffraction pattern (SADP). From Fig. 7(a), the matrix is covered by a number of needle-like and ellipsoidal strengthening precipitates; the clear diffraction spots near 1/3 and 2/3 {022}Al direction and streaks along {111}Al directions (Fig. 7(b)) indicate they are η0 metastable phase. It is generally supposed that η0 phase has a platelet morphology lying on (111)Al plane [16–18], and therefore, viewed along [011]Al direction, two of the four different {111} variants are “edge” on and exhibit a needle-like shape, while the other two exhibit an ellipsoidal shape; this is shown schematically in Fig. 8(a)(b). The platelet-like η0 phase precipitates have a diameter from about 4 nm–15 nm and a thickness from about 0.6 nm–2.5 nm; the aspect ratios of diameter to thickness are from about 4:1 to 10:1. Viewed from o 001 4 Al direction as shown in Fig. 7(c)(d), most strengthening precipitates exhibit a near-round shape, and the diffraction spots due to η0 phase are visible at 1/3 and 2/3 of {220}Al direction. According to these results, η0 metastable phase precipitates are dominant in the specimen; they have very good strengthening effect and therefore results in high strength (Table 1). Moreover, the diffraction spots from Al3Zr dispersoids can also be identified clearly from Fig. 7(d). It was found that in the specimens impacted with strain rate lower than 4000 s 1, the morphology of hardening precipitates in the matrix was similar to that in the as-received specimen; however, the precipitates tend to coarsen with the increase of strain rates. Under bright field mode of TEM, the precipitates could not be seen very clearly due to the presence of a high density of dislocations in the matrix; therefore images were taken under the HAADF mode, and as an example, Fig. 9(a) shows a HAADF image of the specimen impacted at 4000 s 1. Seen along o0114 Al direction, precipitates exhibit a needle-like or an ellipsoidal shape in this specimen, which is similar to that in the specimen before impact; however, their density is lower and some precipitates coarsened. The large platelet-like precipitates have a diameter from about 8 nm–30 nm and a thickness from about 2 nm–6 nm; the aspect ratios of diameter to thickness are from 3:1 to 8:1. Fig. 10 shows o 011 4 Al and o001 4 Al SADPs of the specimen impacted at 4000 s 1. Quite strong diffraction spots from η0 metastable phase and weaker spots from η phase can be seen; this indicates the presence of both η0 and η phase precipitates in the matrix but η0 metastable phase is dominant. There are generally 11 variants of η phase, designated η1-η11, in 7000 series aluminum alloy [19]. According to the feature of diffraction spots of η phase (Fig. 10), it is likely that most η precipitates belong to η2 phase. Moreover, the reflections from η0 metastable phase and η phase precipitates tend to be arced slightly, which indicates a rotation of these precipitates. In the specimen impacted at 6000 s 1, the precipitates density was lower and their morphology was quite different. Seen along o0114 Al direction, as shown by the HAADF image in Fig. 9(b), almost all precipitates exhibit a near-round shape. The size of precipitates is not uniform; for small precipitates, the average size is about 6 nm; for large ones, it is about 14 nm. Fig. 11 shows o0114 Al and o001 4 Al SADPs of this specimen. The diffraction spots from both η0 phase and η phase are visible, which indicates the presence of η0 phase and η phase precipitates. The small
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Fig. 7. TEM images and the corresponding SADPs of the as-received specimen (a)(c) bright field images (b)o 011 4 Al SADP (d) o 0014 Al SADP.
precipitates are likely η0 phase and those large ones are likely η phase. Moreover, the reflections from η0 and η phase precipitates are arced more seriously, which means a significant rotation of these precipitates. Most η phase precipitates may belong to η2 as indicated by the feature of diffraction spots of η phase in Fig. 11(a).
4. Discussion According to above microstructural results, impact deformation exerted significant influence on the precipitation structure, and the change in the type, morphology and size of precipitates is dependent on strain rates. In order to fully understand this phenomenon, it is essential to know the temperature rise during impact deformation. By integrating the true stress-strain curves in Fig. 2, the absorption energy per unit volume of the specimen can be calculated by Eq. (2) [20],
W=
∫0
εf
σ (ε)dε
(2)
Where W is the absorption energy per unit volume of specimen, σ and ε are the true stress and true strain, respectively. And the temperature rise caused by the absorption energy is estimated by Eq. (3) [20],
ΔT =
β ρ Cv
∫0
εf
σ (ε)dε
(3)
Where ΔT is the temperature rise, β is an inverting factor, 0.9, for aluminum, ρ is the density of 7055 aluminum alloy, 2850 kg m 3 and Cv is the specific heat of 7055 aluminum alloy, 860 J kg 1 K 1. The results are given in Fig. 12. It is seen that both W and ΔT tend to increase with the increase of strain rates from 1000 s 1 to 4000 s 1, and then decrease at higher strain rate of 6000 s 1. The highest temperature rise of about 120 K is observed at the strain rate of 4000 s 1. The stressstrain curves in Fig. 2 are actually not representative for room temperature although dynamic impact tests were performed at
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Fig. 8. Schematic of precipitates morphology in the specimens (a)three dimen−− sional view and (b) projection −along [ 110]Al of precipitates in the as-received − specimen, (c) projection along [ 110]Al of precipitates in the specimen impacted at 6000 s 1.
room temperature. The increase of temperature can accelerate dislocation annihilation, and therefore softening occurs, making contribution to the reduction of flow stress during deformation. The temperature rise may also have effect on precipitation structure in the material, however, considering the exposure time was very short (about 100–200 μs), the temperature rise alone could not give rise to significant growth or transformation of precipitates during impact process, and therefore the severe plastic deformation plays a critical role for the change of precipitation structure in the impacted specimen. It has been shown that severe plastic deformation can remarkably alter precipitation behavior in aluminum alloys by introducing a high density of dislocations [17,21,22]. During deformation, dislocations may cut small metastable precipitates and therefore prompt their dissolution; moreover, dislocations can act as hyper channel for rapid diffusion of solutes in the matrix, thereby contributing to fast growth of large precipitates. It was
Fig. 9. HAADF images taken along o 0114 Al direction of the specimen impacted at (a) 4000 s 1 (b) 6000 s 1.
supposed that strain rate of plastic deformation has great influence on diffusion coefficient, and can be expressed by [23,24], ⋅ D′ = 1 + βε DL
(4)
Where D′ is the dynamic diffusion coefficient, DL the static diffu⋅ sion coefficient, β the coefficient of strain and strain rate and ε the strain rate. From this relationship, diffusion coefficient increases greatly with the increase of strain rate during impact deformation. This may explain the increased diameter and thickness of precipitates in the specimen impacted at 4000 s 1. The aspect ratios of diameter to thickness decrease from 4-10 to 3–8, indicating that the thickness increases more rapidly than the diameter, i.e., anisotropic growth occurs. Moreover, some η2 phase precipitates are present, which may indicate that transformation of ηʹ into η2 phase occurred during impact deformation, because η2 phase was supposed to form by direction transformation of ηʹ [19].
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Fig. 10. SADPs of the specimen impacted at 4000 s 1 (a)o 011 4 Al (b) o 0014 Al.
This phenomenon was more significant in the specimen impacted at a higher strain rate of 6000 s 1, as the amount of η2 phase precipitates increases. Therefore, it is reasonable to conclude that severe plastic deformation can prompt transformation of η0 into η phase. However, as shown in Fig. 9(b), most η phase precipitates exhibit a round shape rather than similar morphologies to η0 precipitates. This is similar to previous results reported by Sha et al. [17] that equiaxed η phase precipitates were present in 7136 Al alloy subjected to multiple passes of equal-channel angular pressing (ECAP) at 200 °C. They thought ECAP alters the orientation of η phase precipitates and thus influences the atomic configuration and the interfacial energy at the η/α-Al interfaces; consequently, isotropic growth of η phase precipitates occur, leading to an equiaxed shape. The SADPs in Fig. 11 shows that reflections from η phase precipitates are arced seriously, indicating that impact deformation changed the orientation of η phase greatly; this may make contribution to the change of morphology of η phase precipitates, i.e., from a platelet shape to a global shape (Figs. 7(a), 8(c) & 9(b)). However, the size increased with a higher rate along the thickness direction (perpendicular to {111}Al plane) than along the diameter direction of the platelet precipitate. This is different from the isotropic growth mode of η phase during ECAP
Fig. 11. SADPs of the specimen impacted at 6000 s 1 (a)o0114 Al (b)o 0014 Al.
severe plastic deformation [17]. In order to understand this better, the sizes of one hundred precipitates were measured and their size distribution is shown in Fig. 13. In the as-received specimen, only the diameter of platelet η0 phase precipitates is given; it can be seen that the precipitates with a diameter from about 5 nm–10 nm are dominant. While in the specimen impacted at 6000 s 1, the precipitates with a diameter of about 5–8 nm and about 10–15 nm are dominant. After impact deformation, the number density of precipitates decreases, and the proportion of precipitates with a size larger than about 10 nm increases greatly, while that of smaller ones decreases; a bimodal distribution is visible from Fig. 13. It is known that during severe plastic deformation small precipitates are generally sheared
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highest strain rate of 6000 s 1 gives rise to transformation of η0 phase with a platelet shape into η2 phase with a globe shape; this is likely attributed to the high density of dislocations and the change of orientation of η phase caused by severe plastic deformation during impact deformation.
Acknowledgments This work is supported by the Shenghua Yuying Project (20130603) and the Open-End Fund for the Valuable and Precision Instrument of Central South University (CSUZC201615).
References Fig. 12. Absorption energy every unit volume of the specimen and temperature rise as a function of strain rate.
Fig. 13. Diameter distribution of the precipitates in the as-received specimen and the specimen impacted at 6000 s 1.
by dislocations and can dissolve rapidly [25,26], and this increases supersaturation of the solid solution. Large precipitates are likely to be bypassed by dislocations, and consequently they are surrounded by a high density of dislocations. The pipe diffusion of dislocations can enhance the diffusivity of solutes, and therefore these large precipitates can absorb solutes around them and grow. It is probably that those platelet precipitates with a diameter larger than about 9 nm were more stable and therefore spheroidized and grew at the expense of dissolution of small platelet precipitates around them. In the regions without large precipitates, those small precipitates may dissolve partially or completely due to dislocation shearing and re-precipitate with a globe shape due to the significant change of orientation. As a result, the bimodal distribution of precipitate size shown in Fig. 13 was obtained.
5. Conclusions During impact deformation of 7055 aluminum alloy, the maximum stress and strain rate sensitivity increase first with the increase of strain rate, reach the peak value at about 4000 s 1, and then decrease at the highest strain rate of 6000 s 1. Impact deformation at 6000 s 1 results in serious strain localization and therefore ASB and ASB induced cracks; consequently, the maximum stress and strain rate sensitivity decrease. Impact deformation with strain rates r 4000 s 1 leads to larger η0 phase precipitates and some η2 phase precipitates. The
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Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Flow behavior and microstructure evolution of 7055 aluminum alloy impacted at high strain rates Shengdan Liu a,b,n, Shaoling Wang a,b, Lingying Ye a,b, Yunlai Deng a,b, Xinming Zhang a,b a b
School of Materials Science and Engineering, Central South University, Changsha 410083, PR China Key Laboratory of Nonferrous Metal Materials Science and Engineering, Ministry of Education, Changsha 410083, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 21 July 2016 Received in revised form 10 September 2016 Accepted 13 September 2016 Available online 14 September 2016
The flow behavior and microstructure evolution of 7055 aluminum alloy impacted at the strain rates from 1000 s 1 to 6000 s 1 were investigated by a compressive split-Hopkinson pressure bar, optical microscopy, transmission electron microscopy and scanning transmission electron microscopy. The maximum stress tends to increase first with the increase of strain rate from 1000 s 1 to 4000 s 1, and then decrease at higher strain rate. The strain rate sensitivity is positive in the strain rate range of 1000– 4000 s 1, but became negative above 4000 s 1 because of the presence of shear bands and shear band induced cracks due to intensive strain localization. Impact deformation prompts transformation of η′ phase into η2 phase, and changes the morphology of precipitates from a platelet-like shape to a globelike shape. & 2016 Elsevier B.V. All rights reserved.
Keywords: 7055 aluminum alloy High strain rate Microstructure
1. Introduction 7000 series aluminum alloys are well known precipitation hardenable alloys, and high strength and good ductility can be achieved after aging. Therefore, some alloys such as 7039 and 7020 have been used as armor materials for combat vehicles. In order to enhance the performance of combat vehicles, it is always desirable to use an alloy with higher specific strength and improved ballistic properties. 7055 aluminum alloy can have very high strength more than 600 MPa and simultaneously good ductility and fracture toughness after heat treatment [1–3]. Mondal and Mishra et al. [4] shown that peak-aged 7055 aluminum alloy exhibits better ballistic properties than 7017 aluminum alloy. This implies that this alloy is a good candidate for light armor applications. For this purpose, it is required to fully understand high strain rate behavior of this alloy. To date, there are few literature on this topic. Zhu and Pang [5] reported shear localization of 7055-T77 aluminum alloy under dynamic compression by split Hopkinson pressure bar (SHPB) test method. Mishra and Mondal [6] investigated plastic flow behavior of 7055 aluminum alloy under different high strain rate test methods. In these investigations, however, no attention was paid to precipitation structure in the impacted alloys. In the present study, the dynamic response of 7055 aluminum alloy plate was investigated using SHPB test method, and n Corresponding author at: School of Materials Science and Engineering, Central South University, Changsha 410083, PR China. E-mail addresses: [email protected], [email protected] (S. Liu).
http://dx.doi.org/10.1016/j.msea.2016.09.047 0921-5093/& 2016 Elsevier B.V. All rights reserved.
microstructural evolution was investigated by transmission electron microscopy (TEM) and scanning transmission electron microscopy (STEM).
2. Experimental Cylindrical specimens with a length of 4 mm in thickness and 6 mm in diameter were cut from 7055-T7 aluminum alloy plate for dynamic impact tests, which were performed at room temperature on a split Hopkinson pressure bar machine. The nominal strain rates were selected from 1000 s 1 to 6000 s 1. For each strain rate, three specimens were tested to obtain reliable results. The impact direction was parallel to the rolling direction of the plate. Uniaxial compression tests were also performed on a CRIMS-DDL 100 testing machine with a nominal strain rate of 8 10 4 s 1 at room temperature to obtain quasi-static true stress-strain curves for comparison. The chemical compositions of the alloy were (wt%): Al-7.87 Zn-2.16 Mg-2.05 Cu-0.12 Zr, the contents of Fe and Si were lower than 0.10%; and the tensile properties at room temperature are given in Table 1. Microstructure examination was performed on the specimens before and after impact tests. Metallographic samples were ground, polished and etched in a solution made up of 1 mL HFþ16 mL HNO3 þ3 g CrO3 þ83 mL H2O to reveal grain structure, and then examined on a MX3000 optical microscopy. Moreover, disc samples with a size of 3 mm in diameter and 0.1 mm in thickness were prepared and electropolished in a solution containing 80% ethanol and 20% nitric acid below 20 °C in liquid
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Table 1 Tensile properties of the as-received 7055-T7 aluminum alloy plate. Material
s0.2 (MPa)
sb (MPa)
δ (%)
7055-T7 aluminum alloy plate
586
605
11.6
nitrogen; they were examined on a Tecnai G2 F20 S-Twin scanning transmission electron microscopy (STEM) with a high angle annular dark field (HAADF) detector and probe corrector operated at 200 kV.
3. Results 3.1. Stress-strain curves Fig. 1 shows typical true stress-strain curves of 7055-T7 aluminum alloy under quasi-static compression at a nominal strain rate of 8 10 4 s 1. As seen from the figure, flow stress tends to increase rapidly with the increase of strain and then very slowly with strain larger than about 0.1; stress is about 700 MPa at the strain of 0.4. No softening was observed during quasi-static compression. Fig. 2 shows typical stress-strain curves of 7055-T7 Al alloy impacted at strain rates from 1000 s 1 to 6000 s 1. Strain rate has significant effect on the shape of stress-strain curves. At 1000 s 1, stress increases rapidly in the initial stage and then slowly with the further increase of strain; this trend is similar to that under quasi-static compression as shown in Fig. 1. At higher strain rates, 2000–6000 s 1, the shape of stress-strain curves is quite different; stress tends to increase rapidly to a yield point and then exhibit a significant drop with the increase of strain. This phenomenon is similar to that of some aluminum alloys, for instance, 2219 alloy impacted at high strain rates from 2000 s 1 to 3000 s 1 [7]. Most of impact energy (about 90%) may convert to thermal energy, and therefore increases temperature; consequently, flow softening occurs. With the further increase of strain, stress starts to rise, fluctuate and finally tends to be a constant value. This may indicate that strain hardening and thermal softening occurred during impacted deformation, and finally a balance was achieved. From Fig. 2, it can be seen that the maximum stress tends to increase with the increase of strain rate below 4000 s 1, but decrease at 6000 s 1. The maximum stress was about 650 MPa at the strain rate of 1000 s 1, 690 MPa at 2000 s 1, 790 MPa at 4000 s 1 and 710 MPa at 6000 s 1. In order to have better understanding of the effect of strain rate on the flow behavior during impact
Fig. 2. True stress-strain curves of 7055-T7 aluminum alloy impacted at different strain rates.
⋅
Fig. 3. Relationship between logs and log ε .
deformation, strain rate sensitivity, m, was calculated by Eq. (1) [8] ⋅
m = d log σ /d log ε
(1) .
Where σ is the stress and ε is the strain rate. The results are shown in Fig. 3. From 1000 s 1 to 4000 s 1, there is a positive strain rate sensitivity, which indicates that a higher strain rate would give rise to higher stress due to strain hardening; m was estimated to be 0.088. However, above 4000 s 1, the value of m tends to be negative, about 0.259. This phenomenon was also observed in some previous investigations. For instance, Oosterkamp et al. [8] found for a 7108 aluminum alloy at room temperature strain rate sensitivity was negative at strain rates larger than 2000 s 1. This likely results from the formation of shear bands due to intensive strain localization as shown in following section. 3.2. Microstructure
Fig. 1. True stress-strain curves of 7055-T7 aluminum alloy under quasi-static compression at a nominal strain rate of 8 10 4 s 1.
3.2.1. Optical micrographs Fig. 4 shows optical micrographs of the specimens before and after impact deformation. In the as-received specimen, there is a partially-recrystallized microstructure (Fig. 4(a)); the white regions are recrystallized grains, while the dark regions are unrecrystallized grains containing a number of subgrains. This kind of grain structure is typical in 7000 series aluminum alloy plate or sheet [9,10], and is primarily attributed to the presence and band distribution of Al3Zr dispersoids, which can inhibit recrystallization during solution heat treatment at high temperature [11,12]. In the specimen impacted at 1000 s 1, there is no noticeable change
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Fig. 4. Optical micrographs of 7055 aluminum alloy specimen (a) before impact and impacted at (b) 1000 s 1 (c) 2000 s 1 (d) 4000 s 1 (RD: rolling direction, ND: normal direction, ID: impact direction).
in the grain structure, see Fig. 4(b); this may imply that deformation was quite uniform in the specimens. At the strain rate of 2000 s 1, the grains in the impacted specimen started to be distorted and some deformed shear bands can be found (Fig. 4(c)), which indicates the occurrence of inhomogeneous deformation. In the specimens impacted at 4000 s 1, the grains were distorted more seriously, see Fig. 4(d), and deformed shear bands are clearly visible; this shows that inhomogeneous deformation was serious, leading to strain localization. In the specimen impacted at the highest strain rate of 6000 s 1, adiabatic shear bands (ASB) and ASB induced cracks are clearly visible (Fig. 5), and they are at an angle of about 45°to the impacted direction. Similar phenomenon was also found in a peak-aged 7055 Al alloy impacted at 6500 s 1 [6]. The formation of ASB is generally an indication of highly localized and inhomogeneous deformation, and primarily responsible for the drop in strength during deformation [8]. Therefore, the results in Figs. 4 and 5 may explain the decrease of the maximum stress with strain rate increasing from 4000 s 1 to 6000 s 1 shown in Fig. 2 and the negative strain rate sensitivity shown in Fig. 3.
Fig. 5. Optical micrographs of 7055 aluminum alloy specimen impacted at a strain rate of 6000 s 1 (RD: rolling direction, ND: normal direction).
3.2.2. TEM and STEM examination results The high strength of 7000 series aluminum alloy primarily results from a high density of fine η0 phase strengthening precipitates in the matrix [1,2,13,14]. Though some investigations have been carried out on dynamic impact properties of these alloys, little attention was paid to the change of precipitation structure. Figs. 6–11 show some typical TEM and STEM images of the specimens before and after impact deformation.
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Fig. 6. TEM bright field images of (a) the as-received specimen and (b)(c) the specimen impacted at 6000 s 1.
In the as-received specimen, there is a low density of dislocations inside grains, as shown by the TEM bright field image in Fig. 6(a); (sub) grain boundaries can be identified clearly. After impact deformation at 6000 s 1, a large number of dislocations were introduced, see Fig. 6(b), and it is not easy to see (sub) grain boundaries clearly. Higher magnification image shown in the interior of grains there are a number of dislocation walls, which contained some parallel dislocation lines (Fig. 6(c)); this can enhance strain hardening [15]. Fig. 7 shows TEM bright field images taken along o011 4 Al and o001 4 Al direction and the corresponding selected area diffraction pattern (SADP). From Fig. 7(a), the matrix is covered by a number of needle-like and ellipsoidal strengthening precipitates; the clear diffraction spots near 1/3 and 2/3 {022}Al direction and streaks along {111}Al directions (Fig. 7(b)) indicate they are η0 metastable phase. It is generally supposed that η0 phase has a platelet morphology lying on (111)Al plane [16–18], and therefore, viewed along [011]Al direction, two of the four different {111} variants are “edge” on and exhibit a needle-like shape, while the other two exhibit an ellipsoidal shape; this is shown schematically in Fig. 8(a)(b). The platelet-like η0 phase precipitates have a diameter from about 4 nm–15 nm and a thickness from about 0.6 nm–2.5 nm; the aspect ratios of diameter to thickness are from about 4:1 to 10:1. Viewed from o 001 4 Al direction as shown in Fig. 7(c)(d), most strengthening precipitates exhibit a near-round shape, and the diffraction spots due to η0 phase are visible at 1/3 and 2/3 of {220}Al direction. According to these results, η0 metastable phase precipitates are dominant in the specimen; they have very good strengthening effect and therefore results in high strength (Table 1). Moreover, the diffraction spots from Al3Zr dispersoids can also be identified clearly from Fig. 7(d). It was found that in the specimens impacted with strain rate lower than 4000 s 1, the morphology of hardening precipitates in the matrix was similar to that in the as-received specimen; however, the precipitates tend to coarsen with the increase of strain rates. Under bright field mode of TEM, the precipitates could not be seen very clearly due to the presence of a high density of dislocations in the matrix; therefore images were taken under the HAADF mode, and as an example, Fig. 9(a) shows a HAADF image of the specimen impacted at 4000 s 1. Seen along o0114 Al direction, precipitates exhibit a needle-like or an ellipsoidal shape in this specimen, which is similar to that in the specimen before impact; however, their density is lower and some precipitates coarsened. The large platelet-like precipitates have a diameter from about 8 nm–30 nm and a thickness from about 2 nm–6 nm; the aspect ratios of diameter to thickness are from 3:1 to 8:1. Fig. 10 shows o 011 4 Al and o001 4 Al SADPs of the specimen impacted at 4000 s 1. Quite strong diffraction spots from η0 metastable phase and weaker spots from η phase can be seen; this indicates the presence of both η0 and η phase precipitates in the matrix but η0 metastable phase is dominant. There are generally 11 variants of η phase, designated η1-η11, in 7000 series aluminum alloy [19]. According to the feature of diffraction spots of η phase (Fig. 10), it is likely that most η precipitates belong to η2 phase. Moreover, the reflections from η0 metastable phase and η phase precipitates tend to be arced slightly, which indicates a rotation of these precipitates. In the specimen impacted at 6000 s 1, the precipitates density was lower and their morphology was quite different. Seen along o0114 Al direction, as shown by the HAADF image in Fig. 9(b), almost all precipitates exhibit a near-round shape. The size of precipitates is not uniform; for small precipitates, the average size is about 6 nm; for large ones, it is about 14 nm. Fig. 11 shows o0114 Al and o001 4 Al SADPs of this specimen. The diffraction spots from both η0 phase and η phase are visible, which indicates the presence of η0 phase and η phase precipitates. The small
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Fig. 7. TEM images and the corresponding SADPs of the as-received specimen (a)(c) bright field images (b)o 011 4 Al SADP (d) o 0014 Al SADP.
precipitates are likely η0 phase and those large ones are likely η phase. Moreover, the reflections from η0 and η phase precipitates are arced more seriously, which means a significant rotation of these precipitates. Most η phase precipitates may belong to η2 as indicated by the feature of diffraction spots of η phase in Fig. 11(a).
4. Discussion According to above microstructural results, impact deformation exerted significant influence on the precipitation structure, and the change in the type, morphology and size of precipitates is dependent on strain rates. In order to fully understand this phenomenon, it is essential to know the temperature rise during impact deformation. By integrating the true stress-strain curves in Fig. 2, the absorption energy per unit volume of the specimen can be calculated by Eq. (2) [20],
W=
∫0
εf
σ (ε)dε
(2)
Where W is the absorption energy per unit volume of specimen, σ and ε are the true stress and true strain, respectively. And the temperature rise caused by the absorption energy is estimated by Eq. (3) [20],
ΔT =
β ρ Cv
∫0
εf
σ (ε)dε
(3)
Where ΔT is the temperature rise, β is an inverting factor, 0.9, for aluminum, ρ is the density of 7055 aluminum alloy, 2850 kg m 3 and Cv is the specific heat of 7055 aluminum alloy, 860 J kg 1 K 1. The results are given in Fig. 12. It is seen that both W and ΔT tend to increase with the increase of strain rates from 1000 s 1 to 4000 s 1, and then decrease at higher strain rate of 6000 s 1. The highest temperature rise of about 120 K is observed at the strain rate of 4000 s 1. The stressstrain curves in Fig. 2 are actually not representative for room temperature although dynamic impact tests were performed at
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Fig. 8. Schematic of precipitates morphology in the specimens (a)three dimen−− sional view and (b) projection −along [ 110]Al of precipitates in the as-received − specimen, (c) projection along [ 110]Al of precipitates in the specimen impacted at 6000 s 1.
room temperature. The increase of temperature can accelerate dislocation annihilation, and therefore softening occurs, making contribution to the reduction of flow stress during deformation. The temperature rise may also have effect on precipitation structure in the material, however, considering the exposure time was very short (about 100–200 μs), the temperature rise alone could not give rise to significant growth or transformation of precipitates during impact process, and therefore the severe plastic deformation plays a critical role for the change of precipitation structure in the impacted specimen. It has been shown that severe plastic deformation can remarkably alter precipitation behavior in aluminum alloys by introducing a high density of dislocations [17,21,22]. During deformation, dislocations may cut small metastable precipitates and therefore prompt their dissolution; moreover, dislocations can act as hyper channel for rapid diffusion of solutes in the matrix, thereby contributing to fast growth of large precipitates. It was
Fig. 9. HAADF images taken along o 0114 Al direction of the specimen impacted at (a) 4000 s 1 (b) 6000 s 1.
supposed that strain rate of plastic deformation has great influence on diffusion coefficient, and can be expressed by [23,24], ⋅ D′ = 1 + βε DL
(4)
Where D′ is the dynamic diffusion coefficient, DL the static diffu⋅ sion coefficient, β the coefficient of strain and strain rate and ε the strain rate. From this relationship, diffusion coefficient increases greatly with the increase of strain rate during impact deformation. This may explain the increased diameter and thickness of precipitates in the specimen impacted at 4000 s 1. The aspect ratios of diameter to thickness decrease from 4-10 to 3–8, indicating that the thickness increases more rapidly than the diameter, i.e., anisotropic growth occurs. Moreover, some η2 phase precipitates are present, which may indicate that transformation of ηʹ into η2 phase occurred during impact deformation, because η2 phase was supposed to form by direction transformation of ηʹ [19].
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Fig. 10. SADPs of the specimen impacted at 4000 s 1 (a)o 011 4 Al (b) o 0014 Al.
This phenomenon was more significant in the specimen impacted at a higher strain rate of 6000 s 1, as the amount of η2 phase precipitates increases. Therefore, it is reasonable to conclude that severe plastic deformation can prompt transformation of η0 into η phase. However, as shown in Fig. 9(b), most η phase precipitates exhibit a round shape rather than similar morphologies to η0 precipitates. This is similar to previous results reported by Sha et al. [17] that equiaxed η phase precipitates were present in 7136 Al alloy subjected to multiple passes of equal-channel angular pressing (ECAP) at 200 °C. They thought ECAP alters the orientation of η phase precipitates and thus influences the atomic configuration and the interfacial energy at the η/α-Al interfaces; consequently, isotropic growth of η phase precipitates occur, leading to an equiaxed shape. The SADPs in Fig. 11 shows that reflections from η phase precipitates are arced seriously, indicating that impact deformation changed the orientation of η phase greatly; this may make contribution to the change of morphology of η phase precipitates, i.e., from a platelet shape to a global shape (Figs. 7(a), 8(c) & 9(b)). However, the size increased with a higher rate along the thickness direction (perpendicular to {111}Al plane) than along the diameter direction of the platelet precipitate. This is different from the isotropic growth mode of η phase during ECAP
Fig. 11. SADPs of the specimen impacted at 6000 s 1 (a)o0114 Al (b)o 0014 Al.
severe plastic deformation [17]. In order to understand this better, the sizes of one hundred precipitates were measured and their size distribution is shown in Fig. 13. In the as-received specimen, only the diameter of platelet η0 phase precipitates is given; it can be seen that the precipitates with a diameter from about 5 nm–10 nm are dominant. While in the specimen impacted at 6000 s 1, the precipitates with a diameter of about 5–8 nm and about 10–15 nm are dominant. After impact deformation, the number density of precipitates decreases, and the proportion of precipitates with a size larger than about 10 nm increases greatly, while that of smaller ones decreases; a bimodal distribution is visible from Fig. 13. It is known that during severe plastic deformation small precipitates are generally sheared
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highest strain rate of 6000 s 1 gives rise to transformation of η0 phase with a platelet shape into η2 phase with a globe shape; this is likely attributed to the high density of dislocations and the change of orientation of η phase caused by severe plastic deformation during impact deformation.
Acknowledgments This work is supported by the Shenghua Yuying Project (20130603) and the Open-End Fund for the Valuable and Precision Instrument of Central South University (CSUZC201615).
References Fig. 12. Absorption energy every unit volume of the specimen and temperature rise as a function of strain rate.
Fig. 13. Diameter distribution of the precipitates in the as-received specimen and the specimen impacted at 6000 s 1.
by dislocations and can dissolve rapidly [25,26], and this increases supersaturation of the solid solution. Large precipitates are likely to be bypassed by dislocations, and consequently they are surrounded by a high density of dislocations. The pipe diffusion of dislocations can enhance the diffusivity of solutes, and therefore these large precipitates can absorb solutes around them and grow. It is probably that those platelet precipitates with a diameter larger than about 9 nm were more stable and therefore spheroidized and grew at the expense of dissolution of small platelet precipitates around them. In the regions without large precipitates, those small precipitates may dissolve partially or completely due to dislocation shearing and re-precipitate with a globe shape due to the significant change of orientation. As a result, the bimodal distribution of precipitate size shown in Fig. 13 was obtained.
5. Conclusions During impact deformation of 7055 aluminum alloy, the maximum stress and strain rate sensitivity increase first with the increase of strain rate, reach the peak value at about 4000 s 1, and then decrease at the highest strain rate of 6000 s 1. Impact deformation at 6000 s 1 results in serious strain localization and therefore ASB and ASB induced cracks; consequently, the maximum stress and strain rate sensitivity decrease. Impact deformation with strain rates r 4000 s 1 leads to larger η0 phase precipitates and some η2 phase precipitates. The
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