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An initial report on achieving high comprehensive performance in an Al-Mg-Si alloy via novel thermomechanical processing
Journal of Alloys and Compounds 764 (2018) 679e683
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
An initial report on achieving high comprehensive performance in an Al-Mg-Si alloy via novel thermomechanical processing Jieke Ren a, Zhiguo Chen a, b, *, Jing Peng a, Wenjing Ma a, Simon P. Ringer c, ** a
School of Materials Science and Engineering, Central South University, Changsha 410083, PR China Department of Materials Engineering, Hunan University of Humanities, Science and Technology, Loudi 417000, PR China c Australian Centre for Microscopy & Microanalysis, School of Aerospace, Mechanical & Mechatronic Engineering, The University of Sydney, NSW 2006, Australia b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 January 2018 Received in revised form 2 May 2018 Accepted 6 June 2018 Available online 18 June 2018
A novel thermomechanical processing that significantly augments the tensile mechanical strength of an Al-Mg-Si alloy while maintaining excellent tensile ductility is reported. Our thermomechanical processing involved solution treatment followed by low-temperature hold and asymmetric rolling with thickness reduction of 5e10%, and speed ratios between 1.1 and 1.5. The resultant strength and ductility is striking: the tensile strength levels were similar to or higher than the T6 and T8 tempers, while the tensile ductility was similar to that of the T3 temper. The mechanism of the novel thermomechanical processing is thought to be the result of solute clusters, dislocation substructures, and textural microstructural influences. © 2018 Elsevier B.V. All rights reserved.
Keywords: Thermomechanical processing Asymmetrical rolling Mechanical properties Aluminum alloys
1. Introduction Improvements in the characteristics of structural alloys provide new degrees of freedom in mechanical design and enable thinner structural sections that can result in significant weight reduction [1]. The 6xxx system based on Al-Mg-Si alloys contains the most widely used aluminum alloys for automotive body applications, as they possess good corrosion resistance, welding performance, and formability [2]. However, there is still a large gap in overall performance, with the combination of strength and plasticity less than satisfactory in Al-Mg-Si alloys. The main strengthening precipitate phases in Al-Mg-Si alloys are the b'' phase and b0 phase, which are intermediate metastable precursors of the equilibrium b phase (Mg2Si) [3e6]. Like other heat-treatable aluminum alloys [7,8], the role of pre-precipitate solute clusters rich in Mg and Si are thought to play a significant role in both the hardening and evolution of the
* Corresponding author. Department of Materials Engineering, Hunan University of Humanities, Science and Technology, Loudi 417000, PR China. ** Corresponding author. Australian Centre for Microscopy & Microanalysis, School of Aerospace, Mechanical & Mechatronic Engineering, The University of Sydney, NSW 2006, Australia. E-mail addresses: [email protected] (Z. Chen), [email protected]. au (S.P. Ringer). https://doi.org/10.1016/j.jallcom.2018.06.070 0925-8388/© 2018 Elsevier B.V. All rights reserved.
microstructure in 6xxx alloys [9]. While nonequilibrium secondphase precipitate particles can increase the strength of the alloy by hindering the movement of dislocations, stress concentrations can be introduced in various ways: at grain boundaries in the case of shearable precipitates and at the precipitate/matrix interface in the case of nonshearable precipitates [10]. As a result, conventional heat treatments (such as T6 and T8 heat treatment) for Al-Mg-Si alloys usually increase strength at the cost of ductility [11]. Stress concentration appears to be less of an issue when solute atom clusters are engineered into the microstructure [7e9,12]. Thermomechanical treatments can open up new opportunities to improve the properties of aluminum alloys. However, excellent combinations of strength and ductility are difficult to achieve [13,14]. Here, we provide an initial report on the effect of a novel thermomechanical processing (NTMP) approach on the microstructure and properties of an Al-1.0 Mg-1.0 Si alloy, the content of which falls into the category of many typical aluminum alloys for automobiles.
2. Materials and methods A cold-rolled (CR) Al-1.0 Mg-1.0 Si-0.9 Cu-0.2 Zn (wt. %) alloy 2 mm sheet was used in this research. All samples were solution treated (ST) at 530 C for 1 h and then quenched in water to room
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temperature. The NTMP samples were subsequently aged at 100 C for different periods of time before being subjected to the asymmetrical rolling (ASR) process at room temperature for a total thickness reduction of 5% or 10% with one rolling pass. All NTMP samples were subjected to a final natural aging (NA) period of 15 days (d) prior to mechanical testing and characterization. In this study, the speed ratios (SR) between the rolls during the ASR were 1.1, 1.3 and 1.5. The T3 (ST, 10% strain, NA for 15 d), T4 (ST, NA for 15 d), T6 peak age (aged at 175 C for 8 h), T8 peak age (10% strain, aged at 175 C for 6 h) states are also measured for comparison. Our tensile tests were performed on an MTS 858 instrument. A constant strain rate of 8 104 s1 was used. Three parallel samples were tested for each data point with the magnitudes of the error bars being the standard deviations. The orientation distribution functions (ODF) were measured on a D/Max 2500 X-ray diffractometer. The SEM observations were performed on a Quanta 200 field emission scanning electron microscope with an operating voltage of 20 kV. The specimens for TEM observation were prepared by the standard twin-jet electropolishing method using a solution of methanol and nitric acid (3:1 by volume). The TEM observations were performed on a TECNAI G220 transmission electron microscopy, with an operating voltage of 200 kV.
observed in Fig. 2-(a) that the dominant orientations in the conventional T6 temper are the Goss texture ({011}〈100〉), R-Goss texture ({011}〈110〉), and {001}〈610〉, which deviates slightly from the cube texture ({001}〈100〉). For samples in the conventional T8 temper (Fig. 2-(b)), there is more of a Brass texture ({011}〈211〉) component, due to the rolling deformation applied following ST. The former {001}〈610〉 orientation also continues to rotate towards an R-cube texture ({001}〈110〉), forming {001}〈310〉 as a result of the minor shear stress in conventional rolling. For the sample from Route 2 with SR ¼ 1.3 (Fig. 2-(d)), the volume of the R-cube texture is significantly higher. We attribute this to the fact that the shear stress is greatly increased during the ASR process [17e19]. The intensity of the rolling textures (mainly the Brass texture) is also higher, indicating a larger equivalent strain despite the identical thickness reduction (10%) [20]. The texture decomposition was performed by means of the particle swarm optimization method [21], and the Taylor theory of crystal plasticity was used to calculate the Taylor factor (M), in which the average Taylor factor is the superposition of each textural component (Mi) weighted by its volume fraction (fi) [22], as in Equation (1):
M ¼ Sf i Mi ; i2ð1; nÞ
3. Results and discussion The tensile properties of the Al-Mg-Si alloy after NTMP using different routes are provided in Table 1. It is clear that the strength of the Al-Mg-Si alloy significantly increased by using the NTMP featuring the initial low-temperature hold. In the case of NTMP (Route 2 with SR ¼ 1.5), the ultimate tensile strength increased by over 60 MPa compared with that of the T3 alloy. Generally, the strength of aluminum alloys increases at the expense of their ductility. The T4 alloy exhibits excellent ductility but also records the lowest strength, while the T8 alloy possesses the highest strength, but records the lowest ductility. On the other hand, the NTMP techniques increase the strength of the Al-Mg-Si alloy studied here while maintaining excellent ductility. Specifically, the NTMP samples (Route 2 with SR ¼ 1.3) demonstrate an increase in tensile yield strength (YS) to 289 MPa, ultimate tensile strength (UTS) to 371 MPa, and tensile elongation to 21.4%, which is similar to that of the T3 alloy (22.2%). The relationship between the SR and tensile properties is provided in Fig. 1. Higher strength levels were achieved with higher SR values of 1.3 or 1.5. However, for both 5% and 10% ASR reduction, the elongation was the highest when the SR was 1.3. Following previous research that correlated variations in the tensile properties in materials processed with different SRs with changes in texture [15,16], we have also aimed to investigate changes in the microstructural texture. The ODFs of the Al-Mg-Si alloy are shown in Fig. 2. It can be
(1)
The major decomposed components of the ODFs as well as the Taylor factor (M) are listed in Table 2. The b-fiber components, which are textures with high M values (the Brass, S ({123}〈634〉) and Copper ({112}〈111〉) textures), are greater in Route 2 with SR ¼ 1.3, however, the cube and R-cube components are also high; thus the overall M is only slightly increased. The change in texture components is partly in agreement with a previous work, in which the texture components are significantly affected by SR, and the b-fiber textural components reach a maximum with an SR approximately 1.3 [16]. Fracture surface morphologies of the Al-Mg-Si alloy sample post-tensile testing were investigated using SEM. The results comparing NTMP (both Route 2 with SR ¼ 1.3 and Route 3), T6 and T8 tempers are provided in Fig. 3. As shown in Fig. 3-(a, b), the NTMP sample fractures are mainly composed of tough dimples, which is typical for ductile fracture. In Fig. 3-(c, d), both shear surfaces and tearing ridges were observed, indicating that the fracture mechanism is likely the combination of quasi-cleavage fracture and ductile fracture; in Fig. 3-(d), the features of quasi-cleavage fracture are more pronounced. The ductility of the conventionally treated samples is lower because the quasicleavage fracture is usually considered as a form of brittle fracture [23]. TEM images of the Al-Mg-Si alloy under NTMP (Route 2 with SR ¼ 1.3 and Route 3), are compared with overview microstructures of the T4 and T6 at conventional conditions in Fig. 4.
Table 1 Tensile mechanical properties of the Al-Mg-Si alloy after different treatments. Process
Preaging
Deformation
Final aging
YS [MPa]
UTS [MPa]
Elongation [%]
T3 T4 T6 T8 Route 1
e
100 C 16 h
NA15 d NA15 d 175 C 8 h 175 C 6 h NA15 d
Route 2
100 C 16 h
Route 3
100 C 16 h
CR 10% e e CR 10% ASR5%,1.1SR ASR5%,1.3SR ASR5%,1.5SR ASR10%,1.1SR ASR10%,1.3SR ASR10%,1.5SR CR 10%
206 ± 3 194 ± 3 212 ± 1 313 ± 3 241 ± 9 259 ± 3 249 ± 6 273 ± 5 289 ± 3 289 ± 5 222 ± 6
312 ± 2 309 ± 3 341 ± 7 361 ± 3 338 ± 7 359 ± 2 355 ± 5 352 ± 3 371 ± 1 373 ± 4 343 ± 5
22.2 ± 1.8 25.8 ± 1.2 15.4 ± 0.1 7.6 ± 0.7 24.4 ± 0.8 25.0 ± 0.8 22.5 ± 1.5 19.0 ± 0.8 21.4 ± 0.8 20.2 ± 0.8 20.2 ± 1.5
NA15 d
NA15 d
J. Ren et al. / Journal of Alloys and Compounds 764 (2018) 679e683
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Fig. 1. Tensile mechanical properties for the Al-Mg-Si alloy with different SRs: (a)YS-SR; (b)Elongation-SR.
Fig. 2. The orientation distribution functions of the Al-Mg-Si alloy in different states: (a) T6; (b) T8; (c) Route 3; (d) Route 2 with SR ¼ 1.3.
Table 2 The major decomposed components of the ODFs and the Taylor factors of the Al-Mg-Si alloy. Treatments
Cube(%), M ¼ 2.45
R-Cube(%), M ¼ 2.45
Goss(%), M ¼ 2.45
R-Goss(%), M ¼ 4.9
Brass(%), M ¼ 3.17
b-Fiber(%)
M
T6 T8 Route 3 Route 2, SR ¼ 1.3
12.4 3.6 4.4 6.9
3.8 5.6 8.5 12.9
4.5 1.1 1.5 1.6
5.2 2.1 1.2 4.0
4.3 5.8 9.2 11.7
6.2 7.8 11.1 13.2
3.06 3.21 3.28 3.32
Fig. 3. Fracture morphology of the Al-Mg-Si alloy tensile test samples: (a) Route 2, 1.3 SR; (b) Route 3; (c) T6; (d) T8.
In Fig. 4-(a), without the influence of obvious dislocation contrast, a mottled contrast is apparent that may be the result of a very fine-scale dispersion of precipitates, or may potentially be a strain effect associated with the finer scale features such as clusters [24,25]. No extra diffraction spots were observed in the corresponding diffraction patterns. In Fig. 4-(b), both needle-like and lath-like precipitates were observed. Inspection of the corresponding diffraction patterns suggests that the main precipitate
phase is the needle-shaped b'', which is a critical strengthening phase in many Al-Mg-Si alloys [4]. The atomic Mg/Si ratio of the AlMg-Si alloy is 1.0, lower than the atomic ratio in Mg2Si of 2. Thus, there is a surplus of Si atoms. Accordingly, there is significant probability for the lath-like Q0 phase to be formed, since this phase is a quaternary precursor of the Q phase (Al3Cu2Mg9Si7) and is frequently found in the alloys with similar stoichiometry [26,27]. It is shown in Fig. 4-(c) and Fig. 4-(d) that the NTMP samples contain a
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Fig. 4. Bright field TEM micrographs of the Al-Mg-Si alloy, recorded near to the 〈001〉Al zone axis: (a) T4; (b) T6 peak age; (c) Route 2 with SR ¼ 1.3; (d) Route 3.
high density of dislocations, most of which are entangled together. In this study, an extensive dispersion of fine-scale precipitate particles was not observed and only undissolved constituent Mn containing phases were reported. This paper provides an initial report of the significantly enhanced tensile mechanical properties introduced by NTMP. Further work is required to investigate the detailed mechanisms operating in this highly complex situation. The preliminary inference is as follows. It is widely known that the uniform elongation varies along different directions, which is primarily attributed to the texture [28]. For the sample with a higher R-cube texture fraction (Route 2), a higher plastic strain ratio (also known as the Lankford value or r-value) can be expected [29,30]. Usually, related to the formability of the alloy, the r-value is defined as the ratio of a true strain in the transverse direction to the true strain in the normal direction during the uniform deformation stage of uniaxial tension [29,31,32], i.e., the resistance to the thickness reduction during a tension process. For aluminum sheets, rapid thickness reduction during tension will likely cause strain localization which leads to an eventual necking failure. Thus, it can be inferred that a higher R-cube texture will be beneficial to the uniform elongation. As for the increase in tensile strength, a commonly used model for the tensile yield strength (sy) model is provided in Equation (2):
sy ¼ Dsgb þMttot
(2)
where Dsgb is the strengthening contribution due to grain size effects, M is the Taylor factor, and ttot is the total critical resolved shear stress (CRSS), which is affected by such factors as solid
solution, the potential for solute clustering, and dislocationmediated strengthening [33]. Therefore, one strategy to improve strength can be increasing the Taylor factor. The high dislocation density introduced during the NTMP would also account for some of the strengthening effect. We plan further work to explore the potential role of the contribution of solute atom clusters to strengthening. This is of particular interest when considering the lack of obvious precipitation in the NTMP samples and leaves us intrigued as to the nature of the solute atom distribution in the alloy.
4. Conclusions The mechanical properties of an Al-Mg-Si alloy were significantly improved by the NTMP. The best combination of strength and ductility was observed when the SR between the rolls was 1.3. The yield strength increased by ~80 MPa over that of the T3 temper, and the ultimate tensile strength increased by ~60 MPa, while the excellent elongation of the T3 temper (>20%) is maintained. In the NTMP sample materials, the cube texture rotated toward the {001} 〈110〉 direction. This is likely due to the influence of the higher shear stress, which is more prominent in ASR. Second, the fraction of the Brass texture in these samples was higher, implying a larger equivalent strain even with the same rolling reduction. The excellent improvements of the tensile properties are thought to be due to the synergistic effect of solute clusters, dislocation substructures, and textural influences. Further work is in progress to investigate the mechanistic aspects.
J. Ren et al. / Journal of Alloys and Compounds 764 (2018) 679e683
Funding This work was financially supported by the National Natural Science Foundation of China (Grant No. 51011120052 and 50871123). SPR is grateful for the financial support from the Australian Research Council, the University of Sydney (including the Faculty of Engineering & IT 0 Materials And Structures Research Cluster' program), and technical as well as scientific support from the AMMRF (ammrf.org.au). References [1] W.S. Miller, L. Zhuang, J. Bottema, et al., Recent development in aluminium alloys for the automotive industry, Mater. Sci. Eng. A 280.1 (2000) 37e49. [2] J. Hirsch, Aluminium in innovative light-weight car design, Mater. Trans. 52.5 (2011) 818e824. [3] C. Cayron, P.A. Buffat, Transmission electron microscopy study of the b0 phase (Al-Mg-Si alloys) and QC phase (Al-Cu-Mg-Si alloys): ordering mechanism and crystallographic structure, Acta Mater. 48.10 (2000) 2639e2653. [4] H.S. Hasting, A.G. Frøseth, S.J. Andersen, et al., Composition of b' precipitates in Al-Mg-Si alloys by atom probe tomography and first principles calculations, J. Appl. Phys. 106 (12) (2009) 123527. [5] C.D. Marioara, H. Nordmark, S.J. Andersen, et al., Post-b' phases and their influence on microstructure and hardness in 6xxx Al-Mg-Si alloys, J. Mater. Sci. 41 (2) (2006) 471e478. [6] R. Vissers, M.A. van Huis, J. Jansen, et al., The crystal structure of the b' phase in Al-Mg-Si alloys, Acta Mater. 55.11 (2007) 3815e3823. [7] G. Sha, R.K.W. Marceau, X. Gao, B.C. Muddle, S.P. Ringer, Nanostructure of aluminium alloy 2024: segregation, clustering and precipitation processes, Acta Mater. 59.4 (2011) 1659e1670. [8] P.V. Liddicoat, X.Z. Liao, Y. Zhao, et al., Nanostructural hierarchy increases the strength of aluminium alloys, Nat. Commun. 1.6 (2010) 63. [9] R.K.W. Marceau, A. De Vaucorbeil, G. Sha, S.P. Ringer, W.J. Poole, Analysis of strengthening in AA6111 during the early stages of aging: atom probe tomography and yield stress modeling, Acta Mater. 61.19 (2013) 7285e7303. [10] E. Nembach, Particle Strengthening of Metals and Alloys, John Wiley & Sons, New York, 1997. [11] S. Esmaeili, X. Wang, D.J. Lloyd, et al., On the precipitation-hardening behavior of the Al-Mg-Si-Cu alloy AA6111, Metall. Mater. Trans. A 34 (13) (2003) 751e763. [12] G. Sha, H. Moller, W.E. Stumpf, J.H. Xia, G. Govender, S.P. Ringer, Solute nanostructures and their strengthening effects in Al-7Si-0.6 Mg alloy F357, Acta Mater. 60 (2) (2012) 692e701. [13] Y.J. Huang, Z.G. Chen, Z.Q. Zheng, A conventional thermo-mechanical process of Al-Cu-Mg alloy for increasing ductility while maintaining high strength, Scr. Mater. 64 (5) (2011) 382e385. [14] J.K. Kim, H.K. Kim, J.W. Park, W.J. Kim, Large enhancement in mechanical properties of the 6061 Al alloys after a single pressing by ECAP, Scr. Mater. 53
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(10) (2005) 1207e1211. [15] H.K. Kim, H.W. Kim, J.H. Cho, J.C. Lee, High-formability Al alloy sheet produced by asymmetric rolling of strip-cast sheet, Mater. Sci. Eng. A 574 (2013) 31e36. [16] S. Wronski, B. Ghilianu, T. Chauveau, B. Bacroix, Analysis of textures heterogeneity in cold and warm asymmetrically rolled aluminium, Mater. Char. 62 (1) (2011) 22e34. [17] M. Kadkhodaei, M. Salimi, M. Poursina, Analysis of asymmetrical sheet rolling by a genetic algorithm, Int. J. Mech. Sci. 49 (5) (2007) 622e634. [18] H. Gao, S.C. Ramalingam, G.C. Barber, et al., Analysis of asymmetrical cold rolling with varying coefficients of friction, J. Mater. Process. Technol. 124 (1e2) (2002) 178e182. [19] K.H. Kim, D.N. Lee, Analysis of deformation textures of asymmetrically rolled aluminum sheets, Acta Mater. 49 (13) (2001) 2583e2595. [20] F.Q. Zuo, J.H. Jiang, A.D. Shan, et al., Shear deformation and grain refinement in pure Al by asymmetric rolling, Trans. Nonferrous Met. Soc. China 18 (4) (2008) 774e777. [21] J.G. Tang, et al., Texture decomposition with particle swarm optimization method, Comput. Mater. Sci. 38 (2) (2006) 395e399. [22] W.F. Hosford, The Mechanics of Crystal and Textured Polycrystals, Oxford University Press, Oxford, 1993. [23] Y. Lü, Q. Wang, X. Zeng, et al., Effects of rare earths on the microstructure, properties and fracture behavior of MgeAl alloys, Mater. Sci. Eng. A 278 (1) (2000) 66e76. [24] M. Murayama, K. Hono, M. Saga, et al., Atom probe studies on the early stages of precipitation in AleMgeSi alloys, Mater. Sci. Eng. A 250 (1) (1998) 127e132. [25] K. Matsuda, S. Ikeno, H. Gamada, et al., High-resolution electron microscopy on the structure of Guinier-Preston zones in an Al-1.6 mass Pct Mg2Si alloy, Metall. Mater. Trans. A 29 (4) (1998) 1161e1167. [26] R.S. Yassar, D.P. Field, H. Weiland, The effect of predeformation on the b' and b' precipitates and the role of Q' phase in an Al-Mg-Si alloy; AA6022, Scr. Mater. 53.3 (2005) 299e303. [27] C. Ravi, C. Wolverton, First-principles study of crystal structure and stability of Al-Mg-Si-(Cu) precipitates, Acta Mater. 52.14 (2004) 4213e4227. [28] A.B. Lopes, F. Barlat, J.J. Gracio, et al., Effect of texture and microstructure on strain hardening anisotropy for aluminum deformed in uniaxial tension and simple shear, Int. J. Plast. 19 (1) (2003) 1e22. [29] S.B. Kang, B.K. Min, H.W. Kim, et al., Effect of asymmetric rolling on the texture and mechanical properties of AA6111-aluminum sheet, Metall. Mater. Trans. A 36 (11) (2005) 3141e3149. [30] Y. Suzuki, T. Omura, S. Hirosawa, et al., Formability and bake-hardenability of Al-Mg-Cu alloy sheets by asymmetric warm rolling, Mater. Sci. Forum 519 (2006) 1505e1510. [31] R. Narayanasamy, R. Ravindran, K. Manonmani, et al., A crystallographic texture perspective formability investigation of aluminium 5052 alloy sheets at various annealing temperatures, Mater. Des. 30 (5) (2009) 1804e1817. [32] M. Tajally, E. Emadoddin, Mechanical and anisotropic behaviors of 7075 aluminum alloy sheets, Mater. Des. 32 (3) (2011) 1594e1599. [33] T.A. Parthasarathy, S.I. Rao, D.M. Dimiduk, et al., Contribution to size effect of yield strength from the stochastics of dislocation source lengths in finite samples, Scr. Mater. 56.4 (2007) 313e316.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
An initial report on achieving high comprehensive performance in an Al-Mg-Si alloy via novel thermomechanical processing Jieke Ren a, Zhiguo Chen a, b, *, Jing Peng a, Wenjing Ma a, Simon P. Ringer c, ** a
School of Materials Science and Engineering, Central South University, Changsha 410083, PR China Department of Materials Engineering, Hunan University of Humanities, Science and Technology, Loudi 417000, PR China c Australian Centre for Microscopy & Microanalysis, School of Aerospace, Mechanical & Mechatronic Engineering, The University of Sydney, NSW 2006, Australia b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 January 2018 Received in revised form 2 May 2018 Accepted 6 June 2018 Available online 18 June 2018
A novel thermomechanical processing that significantly augments the tensile mechanical strength of an Al-Mg-Si alloy while maintaining excellent tensile ductility is reported. Our thermomechanical processing involved solution treatment followed by low-temperature hold and asymmetric rolling with thickness reduction of 5e10%, and speed ratios between 1.1 and 1.5. The resultant strength and ductility is striking: the tensile strength levels were similar to or higher than the T6 and T8 tempers, while the tensile ductility was similar to that of the T3 temper. The mechanism of the novel thermomechanical processing is thought to be the result of solute clusters, dislocation substructures, and textural microstructural influences. © 2018 Elsevier B.V. All rights reserved.
Keywords: Thermomechanical processing Asymmetrical rolling Mechanical properties Aluminum alloys
1. Introduction Improvements in the characteristics of structural alloys provide new degrees of freedom in mechanical design and enable thinner structural sections that can result in significant weight reduction [1]. The 6xxx system based on Al-Mg-Si alloys contains the most widely used aluminum alloys for automotive body applications, as they possess good corrosion resistance, welding performance, and formability [2]. However, there is still a large gap in overall performance, with the combination of strength and plasticity less than satisfactory in Al-Mg-Si alloys. The main strengthening precipitate phases in Al-Mg-Si alloys are the b'' phase and b0 phase, which are intermediate metastable precursors of the equilibrium b phase (Mg2Si) [3e6]. Like other heat-treatable aluminum alloys [7,8], the role of pre-precipitate solute clusters rich in Mg and Si are thought to play a significant role in both the hardening and evolution of the
* Corresponding author. Department of Materials Engineering, Hunan University of Humanities, Science and Technology, Loudi 417000, PR China. ** Corresponding author. Australian Centre for Microscopy & Microanalysis, School of Aerospace, Mechanical & Mechatronic Engineering, The University of Sydney, NSW 2006, Australia. E-mail addresses: [email protected] (Z. Chen), [email protected]. au (S.P. Ringer). https://doi.org/10.1016/j.jallcom.2018.06.070 0925-8388/© 2018 Elsevier B.V. All rights reserved.
microstructure in 6xxx alloys [9]. While nonequilibrium secondphase precipitate particles can increase the strength of the alloy by hindering the movement of dislocations, stress concentrations can be introduced in various ways: at grain boundaries in the case of shearable precipitates and at the precipitate/matrix interface in the case of nonshearable precipitates [10]. As a result, conventional heat treatments (such as T6 and T8 heat treatment) for Al-Mg-Si alloys usually increase strength at the cost of ductility [11]. Stress concentration appears to be less of an issue when solute atom clusters are engineered into the microstructure [7e9,12]. Thermomechanical treatments can open up new opportunities to improve the properties of aluminum alloys. However, excellent combinations of strength and ductility are difficult to achieve [13,14]. Here, we provide an initial report on the effect of a novel thermomechanical processing (NTMP) approach on the microstructure and properties of an Al-1.0 Mg-1.0 Si alloy, the content of which falls into the category of many typical aluminum alloys for automobiles.
2. Materials and methods A cold-rolled (CR) Al-1.0 Mg-1.0 Si-0.9 Cu-0.2 Zn (wt. %) alloy 2 mm sheet was used in this research. All samples were solution treated (ST) at 530 C for 1 h and then quenched in water to room
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temperature. The NTMP samples were subsequently aged at 100 C for different periods of time before being subjected to the asymmetrical rolling (ASR) process at room temperature for a total thickness reduction of 5% or 10% with one rolling pass. All NTMP samples were subjected to a final natural aging (NA) period of 15 days (d) prior to mechanical testing and characterization. In this study, the speed ratios (SR) between the rolls during the ASR were 1.1, 1.3 and 1.5. The T3 (ST, 10% strain, NA for 15 d), T4 (ST, NA for 15 d), T6 peak age (aged at 175 C for 8 h), T8 peak age (10% strain, aged at 175 C for 6 h) states are also measured for comparison. Our tensile tests were performed on an MTS 858 instrument. A constant strain rate of 8 104 s1 was used. Three parallel samples were tested for each data point with the magnitudes of the error bars being the standard deviations. The orientation distribution functions (ODF) were measured on a D/Max 2500 X-ray diffractometer. The SEM observations were performed on a Quanta 200 field emission scanning electron microscope with an operating voltage of 20 kV. The specimens for TEM observation were prepared by the standard twin-jet electropolishing method using a solution of methanol and nitric acid (3:1 by volume). The TEM observations were performed on a TECNAI G220 transmission electron microscopy, with an operating voltage of 200 kV.
observed in Fig. 2-(a) that the dominant orientations in the conventional T6 temper are the Goss texture ({011}〈100〉), R-Goss texture ({011}〈110〉), and {001}〈610〉, which deviates slightly from the cube texture ({001}〈100〉). For samples in the conventional T8 temper (Fig. 2-(b)), there is more of a Brass texture ({011}〈211〉) component, due to the rolling deformation applied following ST. The former {001}〈610〉 orientation also continues to rotate towards an R-cube texture ({001}〈110〉), forming {001}〈310〉 as a result of the minor shear stress in conventional rolling. For the sample from Route 2 with SR ¼ 1.3 (Fig. 2-(d)), the volume of the R-cube texture is significantly higher. We attribute this to the fact that the shear stress is greatly increased during the ASR process [17e19]. The intensity of the rolling textures (mainly the Brass texture) is also higher, indicating a larger equivalent strain despite the identical thickness reduction (10%) [20]. The texture decomposition was performed by means of the particle swarm optimization method [21], and the Taylor theory of crystal plasticity was used to calculate the Taylor factor (M), in which the average Taylor factor is the superposition of each textural component (Mi) weighted by its volume fraction (fi) [22], as in Equation (1):
M ¼ Sf i Mi ; i2ð1; nÞ
3. Results and discussion The tensile properties of the Al-Mg-Si alloy after NTMP using different routes are provided in Table 1. It is clear that the strength of the Al-Mg-Si alloy significantly increased by using the NTMP featuring the initial low-temperature hold. In the case of NTMP (Route 2 with SR ¼ 1.5), the ultimate tensile strength increased by over 60 MPa compared with that of the T3 alloy. Generally, the strength of aluminum alloys increases at the expense of their ductility. The T4 alloy exhibits excellent ductility but also records the lowest strength, while the T8 alloy possesses the highest strength, but records the lowest ductility. On the other hand, the NTMP techniques increase the strength of the Al-Mg-Si alloy studied here while maintaining excellent ductility. Specifically, the NTMP samples (Route 2 with SR ¼ 1.3) demonstrate an increase in tensile yield strength (YS) to 289 MPa, ultimate tensile strength (UTS) to 371 MPa, and tensile elongation to 21.4%, which is similar to that of the T3 alloy (22.2%). The relationship between the SR and tensile properties is provided in Fig. 1. Higher strength levels were achieved with higher SR values of 1.3 or 1.5. However, for both 5% and 10% ASR reduction, the elongation was the highest when the SR was 1.3. Following previous research that correlated variations in the tensile properties in materials processed with different SRs with changes in texture [15,16], we have also aimed to investigate changes in the microstructural texture. The ODFs of the Al-Mg-Si alloy are shown in Fig. 2. It can be
(1)
The major decomposed components of the ODFs as well as the Taylor factor (M) are listed in Table 2. The b-fiber components, which are textures with high M values (the Brass, S ({123}〈634〉) and Copper ({112}〈111〉) textures), are greater in Route 2 with SR ¼ 1.3, however, the cube and R-cube components are also high; thus the overall M is only slightly increased. The change in texture components is partly in agreement with a previous work, in which the texture components are significantly affected by SR, and the b-fiber textural components reach a maximum with an SR approximately 1.3 [16]. Fracture surface morphologies of the Al-Mg-Si alloy sample post-tensile testing were investigated using SEM. The results comparing NTMP (both Route 2 with SR ¼ 1.3 and Route 3), T6 and T8 tempers are provided in Fig. 3. As shown in Fig. 3-(a, b), the NTMP sample fractures are mainly composed of tough dimples, which is typical for ductile fracture. In Fig. 3-(c, d), both shear surfaces and tearing ridges were observed, indicating that the fracture mechanism is likely the combination of quasi-cleavage fracture and ductile fracture; in Fig. 3-(d), the features of quasi-cleavage fracture are more pronounced. The ductility of the conventionally treated samples is lower because the quasicleavage fracture is usually considered as a form of brittle fracture [23]. TEM images of the Al-Mg-Si alloy under NTMP (Route 2 with SR ¼ 1.3 and Route 3), are compared with overview microstructures of the T4 and T6 at conventional conditions in Fig. 4.
Table 1 Tensile mechanical properties of the Al-Mg-Si alloy after different treatments. Process
Preaging
Deformation
Final aging
YS [MPa]
UTS [MPa]
Elongation [%]
T3 T4 T6 T8 Route 1
e
100 C 16 h
NA15 d NA15 d 175 C 8 h 175 C 6 h NA15 d
Route 2
100 C 16 h
Route 3
100 C 16 h
CR 10% e e CR 10% ASR5%,1.1SR ASR5%,1.3SR ASR5%,1.5SR ASR10%,1.1SR ASR10%,1.3SR ASR10%,1.5SR CR 10%
206 ± 3 194 ± 3 212 ± 1 313 ± 3 241 ± 9 259 ± 3 249 ± 6 273 ± 5 289 ± 3 289 ± 5 222 ± 6
312 ± 2 309 ± 3 341 ± 7 361 ± 3 338 ± 7 359 ± 2 355 ± 5 352 ± 3 371 ± 1 373 ± 4 343 ± 5
22.2 ± 1.8 25.8 ± 1.2 15.4 ± 0.1 7.6 ± 0.7 24.4 ± 0.8 25.0 ± 0.8 22.5 ± 1.5 19.0 ± 0.8 21.4 ± 0.8 20.2 ± 0.8 20.2 ± 1.5
NA15 d
NA15 d
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681
Fig. 1. Tensile mechanical properties for the Al-Mg-Si alloy with different SRs: (a)YS-SR; (b)Elongation-SR.
Fig. 2. The orientation distribution functions of the Al-Mg-Si alloy in different states: (a) T6; (b) T8; (c) Route 3; (d) Route 2 with SR ¼ 1.3.
Table 2 The major decomposed components of the ODFs and the Taylor factors of the Al-Mg-Si alloy. Treatments
Cube(%), M ¼ 2.45
R-Cube(%), M ¼ 2.45
Goss(%), M ¼ 2.45
R-Goss(%), M ¼ 4.9
Brass(%), M ¼ 3.17
b-Fiber(%)
M
T6 T8 Route 3 Route 2, SR ¼ 1.3
12.4 3.6 4.4 6.9
3.8 5.6 8.5 12.9
4.5 1.1 1.5 1.6
5.2 2.1 1.2 4.0
4.3 5.8 9.2 11.7
6.2 7.8 11.1 13.2
3.06 3.21 3.28 3.32
Fig. 3. Fracture morphology of the Al-Mg-Si alloy tensile test samples: (a) Route 2, 1.3 SR; (b) Route 3; (c) T6; (d) T8.
In Fig. 4-(a), without the influence of obvious dislocation contrast, a mottled contrast is apparent that may be the result of a very fine-scale dispersion of precipitates, or may potentially be a strain effect associated with the finer scale features such as clusters [24,25]. No extra diffraction spots were observed in the corresponding diffraction patterns. In Fig. 4-(b), both needle-like and lath-like precipitates were observed. Inspection of the corresponding diffraction patterns suggests that the main precipitate
phase is the needle-shaped b'', which is a critical strengthening phase in many Al-Mg-Si alloys [4]. The atomic Mg/Si ratio of the AlMg-Si alloy is 1.0, lower than the atomic ratio in Mg2Si of 2. Thus, there is a surplus of Si atoms. Accordingly, there is significant probability for the lath-like Q0 phase to be formed, since this phase is a quaternary precursor of the Q phase (Al3Cu2Mg9Si7) and is frequently found in the alloys with similar stoichiometry [26,27]. It is shown in Fig. 4-(c) and Fig. 4-(d) that the NTMP samples contain a
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Fig. 4. Bright field TEM micrographs of the Al-Mg-Si alloy, recorded near to the 〈001〉Al zone axis: (a) T4; (b) T6 peak age; (c) Route 2 with SR ¼ 1.3; (d) Route 3.
high density of dislocations, most of which are entangled together. In this study, an extensive dispersion of fine-scale precipitate particles was not observed and only undissolved constituent Mn containing phases were reported. This paper provides an initial report of the significantly enhanced tensile mechanical properties introduced by NTMP. Further work is required to investigate the detailed mechanisms operating in this highly complex situation. The preliminary inference is as follows. It is widely known that the uniform elongation varies along different directions, which is primarily attributed to the texture [28]. For the sample with a higher R-cube texture fraction (Route 2), a higher plastic strain ratio (also known as the Lankford value or r-value) can be expected [29,30]. Usually, related to the formability of the alloy, the r-value is defined as the ratio of a true strain in the transverse direction to the true strain in the normal direction during the uniform deformation stage of uniaxial tension [29,31,32], i.e., the resistance to the thickness reduction during a tension process. For aluminum sheets, rapid thickness reduction during tension will likely cause strain localization which leads to an eventual necking failure. Thus, it can be inferred that a higher R-cube texture will be beneficial to the uniform elongation. As for the increase in tensile strength, a commonly used model for the tensile yield strength (sy) model is provided in Equation (2):
sy ¼ Dsgb þMttot
(2)
where Dsgb is the strengthening contribution due to grain size effects, M is the Taylor factor, and ttot is the total critical resolved shear stress (CRSS), which is affected by such factors as solid
solution, the potential for solute clustering, and dislocationmediated strengthening [33]. Therefore, one strategy to improve strength can be increasing the Taylor factor. The high dislocation density introduced during the NTMP would also account for some of the strengthening effect. We plan further work to explore the potential role of the contribution of solute atom clusters to strengthening. This is of particular interest when considering the lack of obvious precipitation in the NTMP samples and leaves us intrigued as to the nature of the solute atom distribution in the alloy.
4. Conclusions The mechanical properties of an Al-Mg-Si alloy were significantly improved by the NTMP. The best combination of strength and ductility was observed when the SR between the rolls was 1.3. The yield strength increased by ~80 MPa over that of the T3 temper, and the ultimate tensile strength increased by ~60 MPa, while the excellent elongation of the T3 temper (>20%) is maintained. In the NTMP sample materials, the cube texture rotated toward the {001} 〈110〉 direction. This is likely due to the influence of the higher shear stress, which is more prominent in ASR. Second, the fraction of the Brass texture in these samples was higher, implying a larger equivalent strain even with the same rolling reduction. The excellent improvements of the tensile properties are thought to be due to the synergistic effect of solute clusters, dislocation substructures, and textural influences. Further work is in progress to investigate the mechanistic aspects.
J. Ren et al. / Journal of Alloys and Compounds 764 (2018) 679e683
Funding This work was financially supported by the National Natural Science Foundation of China (Grant No. 51011120052 and 50871123). SPR is grateful for the financial support from the Australian Research Council, the University of Sydney (including the Faculty of Engineering & IT 0 Materials And Structures Research Cluster' program), and technical as well as scientific support from the AMMRF (ammrf.org.au). References [1] W.S. Miller, L. Zhuang, J. Bottema, et al., Recent development in aluminium alloys for the automotive industry, Mater. Sci. Eng. A 280.1 (2000) 37e49. [2] J. Hirsch, Aluminium in innovative light-weight car design, Mater. Trans. 52.5 (2011) 818e824. [3] C. Cayron, P.A. Buffat, Transmission electron microscopy study of the b0 phase (Al-Mg-Si alloys) and QC phase (Al-Cu-Mg-Si alloys): ordering mechanism and crystallographic structure, Acta Mater. 48.10 (2000) 2639e2653. [4] H.S. Hasting, A.G. Frøseth, S.J. Andersen, et al., Composition of b' precipitates in Al-Mg-Si alloys by atom probe tomography and first principles calculations, J. Appl. Phys. 106 (12) (2009) 123527. [5] C.D. Marioara, H. Nordmark, S.J. Andersen, et al., Post-b' phases and their influence on microstructure and hardness in 6xxx Al-Mg-Si alloys, J. Mater. Sci. 41 (2) (2006) 471e478. [6] R. Vissers, M.A. van Huis, J. Jansen, et al., The crystal structure of the b' phase in Al-Mg-Si alloys, Acta Mater. 55.11 (2007) 3815e3823. [7] G. Sha, R.K.W. Marceau, X. Gao, B.C. Muddle, S.P. Ringer, Nanostructure of aluminium alloy 2024: segregation, clustering and precipitation processes, Acta Mater. 59.4 (2011) 1659e1670. [8] P.V. Liddicoat, X.Z. Liao, Y. Zhao, et al., Nanostructural hierarchy increases the strength of aluminium alloys, Nat. Commun. 1.6 (2010) 63. [9] R.K.W. Marceau, A. De Vaucorbeil, G. Sha, S.P. Ringer, W.J. Poole, Analysis of strengthening in AA6111 during the early stages of aging: atom probe tomography and yield stress modeling, Acta Mater. 61.19 (2013) 7285e7303. [10] E. Nembach, Particle Strengthening of Metals and Alloys, John Wiley & Sons, New York, 1997. [11] S. Esmaeili, X. Wang, D.J. Lloyd, et al., On the precipitation-hardening behavior of the Al-Mg-Si-Cu alloy AA6111, Metall. Mater. Trans. A 34 (13) (2003) 751e763. [12] G. Sha, H. Moller, W.E. Stumpf, J.H. Xia, G. Govender, S.P. Ringer, Solute nanostructures and their strengthening effects in Al-7Si-0.6 Mg alloy F357, Acta Mater. 60 (2) (2012) 692e701. [13] Y.J. Huang, Z.G. Chen, Z.Q. Zheng, A conventional thermo-mechanical process of Al-Cu-Mg alloy for increasing ductility while maintaining high strength, Scr. Mater. 64 (5) (2011) 382e385. [14] J.K. Kim, H.K. Kim, J.W. Park, W.J. Kim, Large enhancement in mechanical properties of the 6061 Al alloys after a single pressing by ECAP, Scr. Mater. 53
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