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Achieve high ductility and strength in an Al–Mg alloy by severe plastic deformation combined with inter-pass annealing
Materials Science & Engineering A 598 (2014) 141–146
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Achieve high ductility and strength in an Al–Mg alloy by severe plastic deformation combined with inter-pass annealing$ Min Zha a,n, Yanjun Li a, Ragnvald H. Mathiesen b, Ruben Bjørge b, Hans J. Roven a,c a
Department of Materials Science and Engineering, Norwegian University of Science and Technology, 7491 Trondheim, Norway Department of Physics, Norwegian University of Science and Technology, 7491 Trondheim, Norway c Center for Advanced Materials, Qatar University, POB 2713 Doha, Qatar b
art ic l e i nf o
a b s t r a c t
Article history: Received 8 November 2013 Received in revised form 30 December 2013 Accepted 31 December 2013 Available online 8 January 2014
Metallic materials processed by severe plastic deformation generally possess very high strength but a rather low uniform elongation, usually o 5%. In the present Al–7Mg alloy processed by room temperature equal channel angular pressing combined with inter-pass annealing, impressive combination of high ductility (elongation 14.5%) and high strength (UTS 600 MPa) was achieved. The high ductility was rationalized by an enhanced work hardening originated from the pronounced dynamic strain aging effect and a bimodal grain structure. The high strength was mainly due to a prominent grain refinement and high dislocation density. & 2014 Elsevier B.V. All rights reserved.
Keywords: ECAP Al–Mg alloys Microstructure Ductility Work hardening Dynamic strain aging
1. Introduction It has been a long-standing goal for materials scientists to synthesize structural materials with balanced combinations of high strength and ductility [1]. However, these two properties are often mutually competitive, i.e., a material may be strong or ductile, but rarely has high strength and ductility at the same time [2]. For SPD-processed ultrafine grained (UFG) materials, plastic deformation usually localizes at the early stage of deformation, resulting in necking followed by specimen failure and hence a rather low uniform elongation [3]. The relative low ductility has severely hampered their commercial applications at ambient temperature. In order to improve the ductility of UFG alloys, socalled “mechanical” strategies via changing the testing parameters, e.g., temperature and/or strain rate, or several “microstructural” strategies via manipulation of the microstructure, have recently
☆ Prime novelty statement: This work might be the first one to achieve a high ductility in a severely plastic deformed Al–Mg alloy with a high Mg concentration by utilizing a combination of room temperature equal channel angular pressing (ECAP) and inter-pass annealing. By such an approach, high ductility (elongation 14.5%) and strength (UTS 600 MPa) were achieved simultaneously in a binary Al–7Mg alloy. The high ductility was rationalized by an enhanced work hardening originating from a high Mg content, pronounced dynamic strain aging effect and a bimodal grain structure. It is believed that the simple but efficient approach can be applied to hard-to-deform alloys, for example, Al alloys containing high content of solid solutions. n Corresponding author. Tel.: þ 47 73 59 70 92; fax: þ47 73 55 02 03. E-mail address: [email protected] (M. Zha).
0921-5093/$ - see front matter & 2014 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.12.103
been developed [4–6]. The “mechanical” strategies are somewhat restricted to low temperatures and high strain rates in order to meet the desired combination of properties. However, the “microstructural” strategies might be the most promising approach to enhance ductility of UFG Al alloys. The latter approach is based on an idea to increase the uniform deformation strain through enhancing the work hardening ability. Increased work hardening can be achieved by generating bi-modal grain structures and controlling deformation-induced nanotwins, second-phase nanoparticles and precipitates [1,3,4,7,8]. Nevertheless, most of the successful attempts so far were based on specific complex multistep processing schedules [3]. It is well known that increasing Mg contents improve strength of Al–Mg alloys by promoting grain refinement, e.g., stable grain sizes established during ECAP. For example, increasing the Mg content from 0 to 3 wt% decreases the grain size from 4 1 μm to 0.25 μm [9]. However, with the Mg content increasing up to 4 wt%, cracking and subsequent failure are likely to occur in initial ECAP passes at ambient temperature, thus limiting the ability of the material to undergo high deformation strains [10]. Promisingly, by utilizing inter-pass annealing (10 min at 200 1C) after every two passes, Furukawa et al. [11] successfully pressed an Al–5Mg–Sc alloy to 8 passes at room temperature. However, no results on mechanical properties of this interpass annealed Al alloy have been reported so far. In the present work, an inter-pass annealing treatment (120 s at 300 1C, determined according to previous work [12]) was introduced in between selected passes during room temperature ECAP of an Al–7 wt%Mg alloy, by which the alloy was processed to an equivalent
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strain of 6 without any severe cracking. Interestingly, the asdeformed alloy showed high strength and also a good ductility. Therefore, one is here attempting to understand the underlying mechanism of the simultaneous increase in strength and ductility. It is proposed that there is a potential of applying the present approach to gain optimal mechanical properties in hard-to-deform alloys, e.g., Al alloys containing high Mg solid solution contents.
2. Experimental procedure The investigated material was taken from Al–7Mg DC-cast ingots supplied by Hydro Aluminum, having the chemical composition (in wt%): Mg 7.0, Fe 0.05, Si 0.06, with Al in balance. ECAP was performed at room temperature with route Bc, using 100 19.5 19.5 mm3 bars in a 901 die, which leads to an imposed strain of about 1.0 per pass. In order to reduce the residual stresses in samples being processed, inter-pass annealing at 300 1C for 120 s was conducted between selected passes. The annealing treatments were performed using molten salt baths, with the temperature controlled within 73 1C. Immediately after annealing, the ECAP bars were quenched in water to room temperature and processed further. After processing, samples from the uniformly deformed central region of the ECAP bars were chosen for microstructure characterization. TEM foils were prepared by twinjet electro polishing in a solution of 33% nitric acid and methanol at 30 1C. The TEM observations were performed in a Philips CM30 operating at 150 KV and a JEOL 2010 at 200 kV. By measuring the length and width of about 100 individual grains from six
TEM micrographs [13], the average size d of the ultrafine grains was determined. Furthermore, EBSD analysis was carried out in a Zeiss 55VP FEG-SEM equipped with a Nordif EBSD detector and TSL OIM software, and performed at 20 kV, 20 mm working distance, 701tilt, and 0.1 μm scan steps. More details about the grain construction procedure during EBSD analysis are available in Ref. [12]. However, for the as-ECAPed samples, the quality of the Kikuchi patterns obtained from EBSD was too low for proper indexing by the OIM software. As a consequence, EBSD analysis was carried out after subsequent low-temperature annealing ( 250 1C for 30 min). Careful analysis revealed that during the low-temperature anneal some recovery occurred. It is believed, however, that the obtained EBSD information roughly represents the as-deformed structure when it comes to grain size and morphology, and that this method offers considerably more information over a larger area compared to traditional TEM analysis. Tensile tests of the processed material were performed at room temperature under a strain rate of 5 10 4 s 1. The specimens were machined from the bars in the longitudinal direction, having a gauge length of 24 mm, a cross section with a diameter of 4 mm and a total length of 40 mm (see the insert in Fig. 1(a)). In order to make the results more reliable, two specimens were tested for the finally processed condition.
3. Results and discussion As reflected by the stress–strain curves in Fig. 1, after 6 passes of ECAP, the Al–7Mg alloy has a good combination of ductility and
Fig. 1. Typical (a) engineering stress–strain curve, inserted is the specimen dimensions in [mm], (b) true stress–strain curve, inserted is the detailed view of the true stress– strain curve at large strains, and work-hardening rate curve as a function of (c) stress and (d) strain of the present Al–7Mg alloy.
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Table 1 Tensile properties of Al–Mg alloys with varying Mg concentrations processed by room temperature ECAP, i.e., comparing data from the present work (Al–7Mg, two samples, here showing the average values) and literature data. Material
Pressing passes (route)
Al–7Mg Al–3Mg [9] Al-3Mg [9] Al–2.77Mg [14] Al–2.5Mg [15] Al–1Mg [9]
6 8 6 4 4 6
(Bc) (intermediately annealed) (Bc) (Bc) (Bc) (A) (Bc)
Yield strength(0.2% offset) (sys) [MPa]
Ultimate tensile stress (sts) [MPa]
Peak true stress [MPa]
Elongation to failure [%]
True fracture strain [%]
538 400 360 – 226 240
600 – – 380 267 –
683 – – 10 – –
14.5 – – 10 – –
12.5 8 11 – 11 13
strength, i.e., 14.5% elongation, 12.5% fracture strain, UTS 600 MPa, and a peak true stress 683 MPa. Although strengthening of metallic materials is the expected outcome of SPD processing, the strength of the present processed Al–7Mg alloy is almost 2–3 times that of more dilute Al–Mg alloys (1–3 wt% Mg) processed by room temperature-ECAP [9,14,15], see Table 1. The noticeable increase in strength reflects that a high level of Mg facilitates significant strengthening in SPD Al alloys. Nevertheless, what makes the present material more interesting is the large elongation to failure (εf) and true fracture strain. Based on the definition of uniform elongation (εu), i.e., the maximum strain where homogeneous plastic flow prevails and the cross-sectional area reduces uniformly, it is observed that deformation of the present Al–7Mg continues in a uniform manner almost to the point of fracture. In other words, the present material has a uniform elongation almost equal to the elongation to failure. Here, please recall that although elongation to failures are around 10% for more dilute Al–Mg alloys not undergoing inter-pass annealing, their uniform elongations are normally o 2% [9,14,15]. Furthermore, the present material exhibits a prominent serrated-flow behavior in the stress–strain curve (Fig. 1(a) and (b)). This is attributed to the dynamic strain aging (DSA) effect, inferring to interactions between mobile Mg solute atoms and dislocations [16]. As can be seen from the magnified segment of the true stress–strain curve inserted in Fig. 1(b), the magnitude of DSA effect, in terms of the difference between the maximum and minimum stress in each serration, Δs, increases gradually with strain, and a maximum Δs value has been measured to be 25 MPa. The increasing Δs could be attributed to increased obstacle effects for dislocation movement, i.e., increased dislocation density and enhanced solute diffusivity. The latter is probably a result of increased deformation induced vacancies with increasing strain [17]. A major influence of the DSA phenomenon is to produce a higher flow stress and, importantly, a greater strain hardening at low- to intermediate strain rates than for higher rates, at which serrations do not appear [6]. Fig. 1(b) also reveals the true stress–strain curve exhibits a considerable work hardening during tensile deformation, which is a major reason for the large ductility [1,8,18]. To further characterize the work hardening ability and thus its relation with the DSA and the improved ductility, the workhardening rate, θ ¼ ∂s=∂ε, where s is the true stress and ε is the true strain, was calculated. The corresponding θ s and θ ε plots are shown in Fig. 1(c) and (d), respectively. Note that smoothing and fitting of the original tensile data were done before calculating the work hardening rate. Fig. 1(c) and (d) reveals that although the initially high work-hardening rate drops rapidly, the declining rate starts slowing down with increasing strain. Thus, the present Al–7Mg alloy has a relatively strong work-hardening ability over the whole strain range and the declined rate of θ is obviously slower than nanostructured Al alloys reported in Ref. [2]. XRD analysis has been conducted and a least squares refinement shows that present as-deformed Al–7Mg has approximately 7 wt% Mg in
solid solution. This is further confirmed by the fact that Mgcontaining precipitates have neither been observed by TEM (Fig. 2) nor been detected by XRD. Hence, the higher work hardening ability of the present severely deformed Al–7Mg is due to the high level of Mg in solid solution that promotes work hardening, as interactions between dislocations and diffused Mg solutes increase the multiplication rate of forest dislocations [19]. Meanwhile, dynamic recovery can be efficiently depressed due to a lower SFE resulting from a high level of Mg in solid solution. More importantly, it is interesting to note that after the initial decrease of θ, an apparent hump began to appear in the θ-curve at larger strains in Fig. 1(c) and (d), which means an increase in work hardening ability of the material. Please note that the onset of the small hump in Fig. 1(c) and (d) corresponds well to the onset of obvious DSA serrations in Fig. 1(b). It is therefore believed that the occurrence of DSA promoted by Mg in solid solution is in favor of the high work hardening ability and hence a high ductility in the present Al–7Mg alloy. A similar result has been reported in a nanostructured 5083 alloy, which exhibits both improved ductility and DSA at lower strain rates [5]. It should be noted, however, that the DSA phenomenon and improved ductility are not always observed simultaneously [5]. Furthermore, conventionally processed Al alloys containing Mg in solid solution normally exhibit a pronounced DSA effect combined with lower ductility at low- to intermediate strain rates ( 10 4–10 2 s 1), e.g. [20]. It should be mentioned that the slight increase in θ just before the strain at fracture (Fig. 1(d)) is just due to the stress increase caused by the final serrated yield jump. The prompt stop at top of the latter is actually the onset of final shear fracture. More investigations need to be carried out for clarification on how the DSA improves the ductility of the present Al–Mg alloy. Furthermore, the work hardening ability of a particular material is closely related to its microstructure. The representative micrographs obtained from TEM in the present case are presented in Fig. 2. These micrographs together with the corresponding selected area diffraction (SAD) patterns in Fig. 2(a) indicate that a submicron sized and near-randomly oriented polycrystalline structure has been developed in this particular area. Note that no Mg-containing precipitates, e.g., Al3Mg2 phase, were observed during the TEM investigations, indicating that most of the Mg prevails in solid solution. A majority of the ultrafine grains observed by TEM appear as elongated, having 50–100 nm width and 100–400 nm length; in addition, a small fraction of fine equiaxed grains o100 nm in size was present. The corresponding grain size distribution chart inserted in Fig. 2(a) provides an estimate of the mean grain size (equivalent diameter d) of 150 nm, i.e. even finer than reported in dilute Al–Mg alloys processed by room temperature ECAP till 8 passes. The latter usually have a mean grain size of 4200 nm [21]. Fig. 2(b) reveals that grain boundaries (GBs) of larger elongated grains are frequently curved and poorly defined, indicating a high level of internal stresses and elastic distortions of the crystal lattice due to the presence of a high density of dislocations [13]. Fig. 2
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Fig. 2. (a) Typical TEM micrograph in the ND–ED plane of the present Al–7Mg alloy, inserted are a corresponding SAED and a grain size distribution chart, (b) a high magnification of grains and (c) details of dislocation tangles inside the grain.
Fig. 3. (a) Typical FESEM-EBSD orientation map (the unit triangle denotes the crystallographic orientations) in the ND–ED plane of the present Al–7Mg alloy (upon low temperature annealing at 250 1C for 30 min), where narrow gray and coarse black lines depict low (21 oΘ o151) and high (151o Θo 1801) angle grain boundaries, respectively; (b) grain size distribution and (c) misorientation angle distribution chart derived from the orientation map (a); inserted in (b) is grain size distribution in number fraction.
M. Zha et al. / Materials Science & Engineering A 598 (2014) 141–146
(c) shows dislocation tangles inside a grain and also the fuzzy GBs, further confirming the high dislocation density and the presence of non-equilibrium GBs. Since only a limited area can be observed by TEM, FESEM-EBSD has been performed on the severely deformed Al–7Mg alloy. A typical orientation image map obtained in the ND–ED plane of the present ECAP material after low temperature annealing (see above) is presented in Fig. 3. A large number of necklace-like ultrafine- and fine grains accompanied by a considerable fraction of micrometer-sized grains containing sub-boundaries can be observed. This indicates that a bimodal grain structure was developed, as further evidenced by the (sub)grain size and misorientation distribution shown in Fig. 3(b) and (c), respectively. It should be noted that clear EBSD orientation maps could only be obtained in regions containing considerable amounts of coarse grains, as regions consisting predominantly of ultrafine grains resulted in poor index patterns. Therefore, the real mean grain size of the material is much smaller than that represented by Fig. 3. The bimodal microstructure developed in the as-deformed material can be attributed to the high level of Mg in solid solution, which favors a lowering of the stacking fault energy (SFE). This greatly depresses dynamic recovery, including rearrangement of diffuse dislocation boundaries into regular sub-grain boundaries, during deformation [12]. Therefore, a large fraction of micron sized grains could in fact survive. Generally, the SPD-induced UFG and nano-sized grains are associated with low work-hardening capability during plastic
145
deformation due to their low dislocation storage ability and efficient dynamic recovery of their non-equilibrium GBs [4]. In contrast, coarse grains having a high ability to accommodate newly formed dislocations can offer the opportunity to expeditiously reap the benefits of the strain hardening capability when ultrafine grains are saturated with dislocations. Therefore, in addition to the occurrence of DSA, the micrometer-sized coarse grains are supposed to be another primary reason for the high work hardening rate and hence good ductility in the present alloy. The relative (point to point) and accumulated (point to origin) misorientation profiles were measured along lines perpendicular and parallel to the elongated direction in two coarse grains in Fig. 3(a), and the profiles are shown in Fig. 4. Along the elongated direction, i.e., along 2 and 4 in Fig. 3(a), large misorientation gradients were observed from one side of the grain to the opposite grain boundary, although the point-to-point misorientation generally remains below 21 (Fig. 4(b) and (d)). The high accumulated misorientation gradient is generally an indication of a high dislocation density, and the density of geometrical necessary dislocations can be estimated from the accumulated misorientation gradient measured from accumulative misorientation profiles [22]. With values for θ and δ deriving from misorientation profiles in Fig. 4, based on ρ E θ/(bδ), where b¼ 2.86 10 10 m is the Burgers vector for pure Al, θ is the accumulated grain misorientation angle (in radians), and δ is the accumulation distance [22], the geometrical necessary dislocation density in these micro-meter sized grains is roughly estimated to be 1.4–15 ( 1014 m 2).
Fig. 4. Misorientation profiles measured along (a) line 1, (b) line 2, (c) line 3 and (d) line 4 indicated in Fig. 3(a), respectively.
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The lower limit might correspond to the dislocation density in grain centers and the higher limit to the dislocation density in areas along the HAGBs. Although this dislocation density might be somewhat underestimated, as the EBSD sample was subjected to a low temperature anneal, the present Al–7Mg alloy has a comparable dislocation density to an HPT-processed Al–4.1Mg alloy, where the reported values are within the range 7.6–12.8 ( 1014 m 2) [23]. Therefore, in addition to the significant grain refinement and a strong solution hardening caused by a high level of Mg in solid solution, the high strength of the present material could also be due to the high dislocation density in both the ultrafine grains and the coarse elongated micrometer sized grains. 4. Conclusions High ductility and strength were achieved simultaneously in the present Al–7Mg alloy processed by room temperature ECAP and inter-pass annealing. The superior combination of mechanical properties can mainly be attributed to the dynamic strain aging effect and a bimodal grain structure promoted by the high level of Mg solid solution in the present material. Dynamic strain aging is observed to be associated with enhanced work hardening ability, leading to an increase in both strength and ductility. In the bimodal structure, the larger grains accommodate relatively high numbers of dislocations, hence promoting higher ductility; moreover, the ultrafine grains need higher local flow stress to continue deformation, thereby contributing to higher strength. Also, the high dislocation density and Mg solid solution strengthening contribute significantly to the strength. It is believed that this simple but efficient ECAP combined with the inter-pass annealed approach can be used in the hard-to-deform alloys, for example, Al alloys containing high content of solid solutions. By manipulating the microstructure through various combinations of ECAP processing and inter-pass annealing, one has ample room and opportunities to tailor the desired properties, in favor of either strength or ductility, or optimize for a combination of both.
Acknowledgments Financial support from the SUP Project ‘Improvement’ (Pnr. 192450) financed by the Research Council of Norway is gratefully acknowledged. The authors also thank Mr. Pål C. Skaret for assisting during ECAP experiments and Hydro Aluminum for providing the experimental material. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]
I. Sabirov, M.Y. Murashkin, R.Z. Valiev, Mater. Sci. Eng. A 560 (2013) 1. Y.-H. Zhao, X.-Z. Liao, S. Cheng, E. Ma, Y.T. Zhu, Adv. Mater. 18 (2006) 2280. Y. Estrin, A. Vinogradov, Acta Mater. 61 (2013) 782. Y.M. Wang, E. Ma, Acta Mater. 52 (2004) 1699. B.Q. Han, J.Y. Huang, Y.T. Zhu, E.J. Lavernia, Scr. Mater. 54 (2006) 1175. M. Wagenhofer, M. Erickson-Natishan, R.W. Armstrong, F.J. Zerilli, Scr. Mater. 41 (1999) 1177. R.Z. Valiev, T.G. Langdon, Prog. Mater. Sci. 51 (2006) 881. H.-Y. Wang, X.-L. Nan, N. Zhang, C. Wang, J.-G. Wang, Q.-C. Jiang, Mater. Chem. Phys. 132 (2012) 248. Y. Iwahashi, Z. Horita, M. Nemoto, T.G. Langdon, Metall. Mater. Trans. A 29 (1998) 2503. V.V. Stolyarov, R. Lapovok, J. Alloys Compd. 378 (2004) 233. M. Furukawa, A. Utsunomiya, K. Matsubara, Z. Horita, T.G. Langdon, Acta Mater. 49 (2001) 3829. M. Zha, Y. Li, R.H. Mathiesen, R. Bjørge, H.J. Roven, Mater. Sci. Eng. A 586 (2013) 374. Zha, Li, Mathiesen, Bjørge and Roven, Trans. Nonferr. Met. Soc. China 2014, (in press, the 6 issue in 2014). D.R. Fang, Q.Q. Duan, N.Q. Zhao, J.J. LI, S.D. Wu, Z.F. Zhang, Mater. Sci. Eng. A 459 (2007) 137. R. Kapoor, N. Kumar, R.S. Mishra, C.S. Huskamp, K.K. Sankaran, Mater. Sci. Eng. A 527 (2010) 5246. E.L. Huskins, B. Cao, K.T. Ramesh, Mater. Sci. Eng. A 527 (2010) 1292. J.M. Robinson, M.P. Shaw, Int. Mater. Rev. 39 (1994) 113. H.-Y. Wang, E.-S. Xue, X.-L. Nan, T. Yue, Y.-P. Wang, Q.-C. Jiang, Scr. Mater. 68 (2013) 229. N.Y. Zolotorevsky, A.N. Solonin, A.Y. Churyumov, V.S. Zolotorevsky, Mater. Sci. Eng. A 502 (2009) 111. A. Fjeldly, H.J. Roven, Acta Mater. 44 (1996) 3497. M. Furukawa, Z. Horita, M. Nemoto, R.Z. Valiev, T.G. Langdon, Philos. Mag. a – Phys. Condens. Matter Struct. Defects Mech. Prop. 78 (1998) 203. D. Jorge-Badiola, A. Iza-Mendia, I. Gutiérrez, Mater. Sci. Eng. A 394 (2005) 445. M. Liu, H.J. Roven, X. Liu, M. Murashkin, R.Z. Valiev, T. Ungar, L. Balogh, J. Mater. Sci. 45 (2010) 4659.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Rapid communication
Achieve high ductility and strength in an Al–Mg alloy by severe plastic deformation combined with inter-pass annealing$ Min Zha a,n, Yanjun Li a, Ragnvald H. Mathiesen b, Ruben Bjørge b, Hans J. Roven a,c a
Department of Materials Science and Engineering, Norwegian University of Science and Technology, 7491 Trondheim, Norway Department of Physics, Norwegian University of Science and Technology, 7491 Trondheim, Norway c Center for Advanced Materials, Qatar University, POB 2713 Doha, Qatar b
art ic l e i nf o
a b s t r a c t
Article history: Received 8 November 2013 Received in revised form 30 December 2013 Accepted 31 December 2013 Available online 8 January 2014
Metallic materials processed by severe plastic deformation generally possess very high strength but a rather low uniform elongation, usually o 5%. In the present Al–7Mg alloy processed by room temperature equal channel angular pressing combined with inter-pass annealing, impressive combination of high ductility (elongation 14.5%) and high strength (UTS 600 MPa) was achieved. The high ductility was rationalized by an enhanced work hardening originated from the pronounced dynamic strain aging effect and a bimodal grain structure. The high strength was mainly due to a prominent grain refinement and high dislocation density. & 2014 Elsevier B.V. All rights reserved.
Keywords: ECAP Al–Mg alloys Microstructure Ductility Work hardening Dynamic strain aging
1. Introduction It has been a long-standing goal for materials scientists to synthesize structural materials with balanced combinations of high strength and ductility [1]. However, these two properties are often mutually competitive, i.e., a material may be strong or ductile, but rarely has high strength and ductility at the same time [2]. For SPD-processed ultrafine grained (UFG) materials, plastic deformation usually localizes at the early stage of deformation, resulting in necking followed by specimen failure and hence a rather low uniform elongation [3]. The relative low ductility has severely hampered their commercial applications at ambient temperature. In order to improve the ductility of UFG alloys, socalled “mechanical” strategies via changing the testing parameters, e.g., temperature and/or strain rate, or several “microstructural” strategies via manipulation of the microstructure, have recently
☆ Prime novelty statement: This work might be the first one to achieve a high ductility in a severely plastic deformed Al–Mg alloy with a high Mg concentration by utilizing a combination of room temperature equal channel angular pressing (ECAP) and inter-pass annealing. By such an approach, high ductility (elongation 14.5%) and strength (UTS 600 MPa) were achieved simultaneously in a binary Al–7Mg alloy. The high ductility was rationalized by an enhanced work hardening originating from a high Mg content, pronounced dynamic strain aging effect and a bimodal grain structure. It is believed that the simple but efficient approach can be applied to hard-to-deform alloys, for example, Al alloys containing high content of solid solutions. n Corresponding author. Tel.: þ 47 73 59 70 92; fax: þ47 73 55 02 03. E-mail address: [email protected] (M. Zha).
0921-5093/$ - see front matter & 2014 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.12.103
been developed [4–6]. The “mechanical” strategies are somewhat restricted to low temperatures and high strain rates in order to meet the desired combination of properties. However, the “microstructural” strategies might be the most promising approach to enhance ductility of UFG Al alloys. The latter approach is based on an idea to increase the uniform deformation strain through enhancing the work hardening ability. Increased work hardening can be achieved by generating bi-modal grain structures and controlling deformation-induced nanotwins, second-phase nanoparticles and precipitates [1,3,4,7,8]. Nevertheless, most of the successful attempts so far were based on specific complex multistep processing schedules [3]. It is well known that increasing Mg contents improve strength of Al–Mg alloys by promoting grain refinement, e.g., stable grain sizes established during ECAP. For example, increasing the Mg content from 0 to 3 wt% decreases the grain size from 4 1 μm to 0.25 μm [9]. However, with the Mg content increasing up to 4 wt%, cracking and subsequent failure are likely to occur in initial ECAP passes at ambient temperature, thus limiting the ability of the material to undergo high deformation strains [10]. Promisingly, by utilizing inter-pass annealing (10 min at 200 1C) after every two passes, Furukawa et al. [11] successfully pressed an Al–5Mg–Sc alloy to 8 passes at room temperature. However, no results on mechanical properties of this interpass annealed Al alloy have been reported so far. In the present work, an inter-pass annealing treatment (120 s at 300 1C, determined according to previous work [12]) was introduced in between selected passes during room temperature ECAP of an Al–7 wt%Mg alloy, by which the alloy was processed to an equivalent
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strain of 6 without any severe cracking. Interestingly, the asdeformed alloy showed high strength and also a good ductility. Therefore, one is here attempting to understand the underlying mechanism of the simultaneous increase in strength and ductility. It is proposed that there is a potential of applying the present approach to gain optimal mechanical properties in hard-to-deform alloys, e.g., Al alloys containing high Mg solid solution contents.
2. Experimental procedure The investigated material was taken from Al–7Mg DC-cast ingots supplied by Hydro Aluminum, having the chemical composition (in wt%): Mg 7.0, Fe 0.05, Si 0.06, with Al in balance. ECAP was performed at room temperature with route Bc, using 100 19.5 19.5 mm3 bars in a 901 die, which leads to an imposed strain of about 1.0 per pass. In order to reduce the residual stresses in samples being processed, inter-pass annealing at 300 1C for 120 s was conducted between selected passes. The annealing treatments were performed using molten salt baths, with the temperature controlled within 73 1C. Immediately after annealing, the ECAP bars were quenched in water to room temperature and processed further. After processing, samples from the uniformly deformed central region of the ECAP bars were chosen for microstructure characterization. TEM foils were prepared by twinjet electro polishing in a solution of 33% nitric acid and methanol at 30 1C. The TEM observations were performed in a Philips CM30 operating at 150 KV and a JEOL 2010 at 200 kV. By measuring the length and width of about 100 individual grains from six
TEM micrographs [13], the average size d of the ultrafine grains was determined. Furthermore, EBSD analysis was carried out in a Zeiss 55VP FEG-SEM equipped with a Nordif EBSD detector and TSL OIM software, and performed at 20 kV, 20 mm working distance, 701tilt, and 0.1 μm scan steps. More details about the grain construction procedure during EBSD analysis are available in Ref. [12]. However, for the as-ECAPed samples, the quality of the Kikuchi patterns obtained from EBSD was too low for proper indexing by the OIM software. As a consequence, EBSD analysis was carried out after subsequent low-temperature annealing ( 250 1C for 30 min). Careful analysis revealed that during the low-temperature anneal some recovery occurred. It is believed, however, that the obtained EBSD information roughly represents the as-deformed structure when it comes to grain size and morphology, and that this method offers considerably more information over a larger area compared to traditional TEM analysis. Tensile tests of the processed material were performed at room temperature under a strain rate of 5 10 4 s 1. The specimens were machined from the bars in the longitudinal direction, having a gauge length of 24 mm, a cross section with a diameter of 4 mm and a total length of 40 mm (see the insert in Fig. 1(a)). In order to make the results more reliable, two specimens were tested for the finally processed condition.
3. Results and discussion As reflected by the stress–strain curves in Fig. 1, after 6 passes of ECAP, the Al–7Mg alloy has a good combination of ductility and
Fig. 1. Typical (a) engineering stress–strain curve, inserted is the specimen dimensions in [mm], (b) true stress–strain curve, inserted is the detailed view of the true stress– strain curve at large strains, and work-hardening rate curve as a function of (c) stress and (d) strain of the present Al–7Mg alloy.
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Table 1 Tensile properties of Al–Mg alloys with varying Mg concentrations processed by room temperature ECAP, i.e., comparing data from the present work (Al–7Mg, two samples, here showing the average values) and literature data. Material
Pressing passes (route)
Al–7Mg Al–3Mg [9] Al-3Mg [9] Al–2.77Mg [14] Al–2.5Mg [15] Al–1Mg [9]
6 8 6 4 4 6
(Bc) (intermediately annealed) (Bc) (Bc) (Bc) (A) (Bc)
Yield strength(0.2% offset) (sys) [MPa]
Ultimate tensile stress (sts) [MPa]
Peak true stress [MPa]
Elongation to failure [%]
True fracture strain [%]
538 400 360 – 226 240
600 – – 380 267 –
683 – – 10 – –
14.5 – – 10 – –
12.5 8 11 – 11 13
strength, i.e., 14.5% elongation, 12.5% fracture strain, UTS 600 MPa, and a peak true stress 683 MPa. Although strengthening of metallic materials is the expected outcome of SPD processing, the strength of the present processed Al–7Mg alloy is almost 2–3 times that of more dilute Al–Mg alloys (1–3 wt% Mg) processed by room temperature-ECAP [9,14,15], see Table 1. The noticeable increase in strength reflects that a high level of Mg facilitates significant strengthening in SPD Al alloys. Nevertheless, what makes the present material more interesting is the large elongation to failure (εf) and true fracture strain. Based on the definition of uniform elongation (εu), i.e., the maximum strain where homogeneous plastic flow prevails and the cross-sectional area reduces uniformly, it is observed that deformation of the present Al–7Mg continues in a uniform manner almost to the point of fracture. In other words, the present material has a uniform elongation almost equal to the elongation to failure. Here, please recall that although elongation to failures are around 10% for more dilute Al–Mg alloys not undergoing inter-pass annealing, their uniform elongations are normally o 2% [9,14,15]. Furthermore, the present material exhibits a prominent serrated-flow behavior in the stress–strain curve (Fig. 1(a) and (b)). This is attributed to the dynamic strain aging (DSA) effect, inferring to interactions between mobile Mg solute atoms and dislocations [16]. As can be seen from the magnified segment of the true stress–strain curve inserted in Fig. 1(b), the magnitude of DSA effect, in terms of the difference between the maximum and minimum stress in each serration, Δs, increases gradually with strain, and a maximum Δs value has been measured to be 25 MPa. The increasing Δs could be attributed to increased obstacle effects for dislocation movement, i.e., increased dislocation density and enhanced solute diffusivity. The latter is probably a result of increased deformation induced vacancies with increasing strain [17]. A major influence of the DSA phenomenon is to produce a higher flow stress and, importantly, a greater strain hardening at low- to intermediate strain rates than for higher rates, at which serrations do not appear [6]. Fig. 1(b) also reveals the true stress–strain curve exhibits a considerable work hardening during tensile deformation, which is a major reason for the large ductility [1,8,18]. To further characterize the work hardening ability and thus its relation with the DSA and the improved ductility, the workhardening rate, θ ¼ ∂s=∂ε, where s is the true stress and ε is the true strain, was calculated. The corresponding θ s and θ ε plots are shown in Fig. 1(c) and (d), respectively. Note that smoothing and fitting of the original tensile data were done before calculating the work hardening rate. Fig. 1(c) and (d) reveals that although the initially high work-hardening rate drops rapidly, the declining rate starts slowing down with increasing strain. Thus, the present Al–7Mg alloy has a relatively strong work-hardening ability over the whole strain range and the declined rate of θ is obviously slower than nanostructured Al alloys reported in Ref. [2]. XRD analysis has been conducted and a least squares refinement shows that present as-deformed Al–7Mg has approximately 7 wt% Mg in
solid solution. This is further confirmed by the fact that Mgcontaining precipitates have neither been observed by TEM (Fig. 2) nor been detected by XRD. Hence, the higher work hardening ability of the present severely deformed Al–7Mg is due to the high level of Mg in solid solution that promotes work hardening, as interactions between dislocations and diffused Mg solutes increase the multiplication rate of forest dislocations [19]. Meanwhile, dynamic recovery can be efficiently depressed due to a lower SFE resulting from a high level of Mg in solid solution. More importantly, it is interesting to note that after the initial decrease of θ, an apparent hump began to appear in the θ-curve at larger strains in Fig. 1(c) and (d), which means an increase in work hardening ability of the material. Please note that the onset of the small hump in Fig. 1(c) and (d) corresponds well to the onset of obvious DSA serrations in Fig. 1(b). It is therefore believed that the occurrence of DSA promoted by Mg in solid solution is in favor of the high work hardening ability and hence a high ductility in the present Al–7Mg alloy. A similar result has been reported in a nanostructured 5083 alloy, which exhibits both improved ductility and DSA at lower strain rates [5]. It should be noted, however, that the DSA phenomenon and improved ductility are not always observed simultaneously [5]. Furthermore, conventionally processed Al alloys containing Mg in solid solution normally exhibit a pronounced DSA effect combined with lower ductility at low- to intermediate strain rates ( 10 4–10 2 s 1), e.g. [20]. It should be mentioned that the slight increase in θ just before the strain at fracture (Fig. 1(d)) is just due to the stress increase caused by the final serrated yield jump. The prompt stop at top of the latter is actually the onset of final shear fracture. More investigations need to be carried out for clarification on how the DSA improves the ductility of the present Al–Mg alloy. Furthermore, the work hardening ability of a particular material is closely related to its microstructure. The representative micrographs obtained from TEM in the present case are presented in Fig. 2. These micrographs together with the corresponding selected area diffraction (SAD) patterns in Fig. 2(a) indicate that a submicron sized and near-randomly oriented polycrystalline structure has been developed in this particular area. Note that no Mg-containing precipitates, e.g., Al3Mg2 phase, were observed during the TEM investigations, indicating that most of the Mg prevails in solid solution. A majority of the ultrafine grains observed by TEM appear as elongated, having 50–100 nm width and 100–400 nm length; in addition, a small fraction of fine equiaxed grains o100 nm in size was present. The corresponding grain size distribution chart inserted in Fig. 2(a) provides an estimate of the mean grain size (equivalent diameter d) of 150 nm, i.e. even finer than reported in dilute Al–Mg alloys processed by room temperature ECAP till 8 passes. The latter usually have a mean grain size of 4200 nm [21]. Fig. 2(b) reveals that grain boundaries (GBs) of larger elongated grains are frequently curved and poorly defined, indicating a high level of internal stresses and elastic distortions of the crystal lattice due to the presence of a high density of dislocations [13]. Fig. 2
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Fig. 2. (a) Typical TEM micrograph in the ND–ED plane of the present Al–7Mg alloy, inserted are a corresponding SAED and a grain size distribution chart, (b) a high magnification of grains and (c) details of dislocation tangles inside the grain.
Fig. 3. (a) Typical FESEM-EBSD orientation map (the unit triangle denotes the crystallographic orientations) in the ND–ED plane of the present Al–7Mg alloy (upon low temperature annealing at 250 1C for 30 min), where narrow gray and coarse black lines depict low (21 oΘ o151) and high (151o Θo 1801) angle grain boundaries, respectively; (b) grain size distribution and (c) misorientation angle distribution chart derived from the orientation map (a); inserted in (b) is grain size distribution in number fraction.
M. Zha et al. / Materials Science & Engineering A 598 (2014) 141–146
(c) shows dislocation tangles inside a grain and also the fuzzy GBs, further confirming the high dislocation density and the presence of non-equilibrium GBs. Since only a limited area can be observed by TEM, FESEM-EBSD has been performed on the severely deformed Al–7Mg alloy. A typical orientation image map obtained in the ND–ED plane of the present ECAP material after low temperature annealing (see above) is presented in Fig. 3. A large number of necklace-like ultrafine- and fine grains accompanied by a considerable fraction of micrometer-sized grains containing sub-boundaries can be observed. This indicates that a bimodal grain structure was developed, as further evidenced by the (sub)grain size and misorientation distribution shown in Fig. 3(b) and (c), respectively. It should be noted that clear EBSD orientation maps could only be obtained in regions containing considerable amounts of coarse grains, as regions consisting predominantly of ultrafine grains resulted in poor index patterns. Therefore, the real mean grain size of the material is much smaller than that represented by Fig. 3. The bimodal microstructure developed in the as-deformed material can be attributed to the high level of Mg in solid solution, which favors a lowering of the stacking fault energy (SFE). This greatly depresses dynamic recovery, including rearrangement of diffuse dislocation boundaries into regular sub-grain boundaries, during deformation [12]. Therefore, a large fraction of micron sized grains could in fact survive. Generally, the SPD-induced UFG and nano-sized grains are associated with low work-hardening capability during plastic
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deformation due to their low dislocation storage ability and efficient dynamic recovery of their non-equilibrium GBs [4]. In contrast, coarse grains having a high ability to accommodate newly formed dislocations can offer the opportunity to expeditiously reap the benefits of the strain hardening capability when ultrafine grains are saturated with dislocations. Therefore, in addition to the occurrence of DSA, the micrometer-sized coarse grains are supposed to be another primary reason for the high work hardening rate and hence good ductility in the present alloy. The relative (point to point) and accumulated (point to origin) misorientation profiles were measured along lines perpendicular and parallel to the elongated direction in two coarse grains in Fig. 3(a), and the profiles are shown in Fig. 4. Along the elongated direction, i.e., along 2 and 4 in Fig. 3(a), large misorientation gradients were observed from one side of the grain to the opposite grain boundary, although the point-to-point misorientation generally remains below 21 (Fig. 4(b) and (d)). The high accumulated misorientation gradient is generally an indication of a high dislocation density, and the density of geometrical necessary dislocations can be estimated from the accumulated misorientation gradient measured from accumulative misorientation profiles [22]. With values for θ and δ deriving from misorientation profiles in Fig. 4, based on ρ E θ/(bδ), where b¼ 2.86 10 10 m is the Burgers vector for pure Al, θ is the accumulated grain misorientation angle (in radians), and δ is the accumulation distance [22], the geometrical necessary dislocation density in these micro-meter sized grains is roughly estimated to be 1.4–15 ( 1014 m 2).
Fig. 4. Misorientation profiles measured along (a) line 1, (b) line 2, (c) line 3 and (d) line 4 indicated in Fig. 3(a), respectively.
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The lower limit might correspond to the dislocation density in grain centers and the higher limit to the dislocation density in areas along the HAGBs. Although this dislocation density might be somewhat underestimated, as the EBSD sample was subjected to a low temperature anneal, the present Al–7Mg alloy has a comparable dislocation density to an HPT-processed Al–4.1Mg alloy, where the reported values are within the range 7.6–12.8 ( 1014 m 2) [23]. Therefore, in addition to the significant grain refinement and a strong solution hardening caused by a high level of Mg in solid solution, the high strength of the present material could also be due to the high dislocation density in both the ultrafine grains and the coarse elongated micrometer sized grains. 4. Conclusions High ductility and strength were achieved simultaneously in the present Al–7Mg alloy processed by room temperature ECAP and inter-pass annealing. The superior combination of mechanical properties can mainly be attributed to the dynamic strain aging effect and a bimodal grain structure promoted by the high level of Mg solid solution in the present material. Dynamic strain aging is observed to be associated with enhanced work hardening ability, leading to an increase in both strength and ductility. In the bimodal structure, the larger grains accommodate relatively high numbers of dislocations, hence promoting higher ductility; moreover, the ultrafine grains need higher local flow stress to continue deformation, thereby contributing to higher strength. Also, the high dislocation density and Mg solid solution strengthening contribute significantly to the strength. It is believed that this simple but efficient ECAP combined with the inter-pass annealed approach can be used in the hard-to-deform alloys, for example, Al alloys containing high content of solid solutions. By manipulating the microstructure through various combinations of ECAP processing and inter-pass annealing, one has ample room and opportunities to tailor the desired properties, in favor of either strength or ductility, or optimize for a combination of both.
Acknowledgments Financial support from the SUP Project ‘Improvement’ (Pnr. 192450) financed by the Research Council of Norway is gratefully acknowledged. The authors also thank Mr. Pål C. Skaret for assisting during ECAP experiments and Hydro Aluminum for providing the experimental material. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]
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