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1 Materials Science Forum Online: ISSN: , Vols , pp doi: / Trans Tech Publications, Switzerland Influence of Scandium Addition on the High Temperature Compressive Strength of Aluminium Alloy 7010 A.K.Mukhopadhyay *, K.S.Prasad and A.Dutta Defence Metallurgical Research Laboratory, Kanchanbagh, Hyderabad , India * Corresponding author s address: ashim_mukhopadhyay@yahoo.com Keywords: Aluminium alloy 7010; Sc addition; homogenization; compression test; transmission electron microscopy Abstract: The influence of Sc addition on the high temperature compressive strength of a commercial alloy 7010 (hereafter termed base alloy) has been examined. The base alloy, and the base alloy with 0.23 wt% Sc were cast, homogenized and subjected to compression tests at temperatures ranging from 300 to 450 o C and strain rates of 10-3, 10-2, 10-1 and 1 sec -1. It is shown that Sc addition to the base alloy increases the compressive flow stress under these deformation conditions. The increase in peak flow stress is nearly 3-6 times the peak flow stress of the base alloy at temperatures o C over the strain rate range investigated. Whilst, at temperatures 400 o C, the flow stresses decrease significantly irrespective of the strain rate used. Transmission electron microscopy (TEM) revealed that a combination of (1) increased nucleation frequency of dispersoids, (2) evolution of smaller subgrain size, and (3) refinement of alloy phases in the Al-Zn-Mg-Cu system contribute to superior strengthening in the alloy containing Sc. Whilst, it is primarily a combination of coarsening and instability of the alloy phases in the Al-Zn-Mg-Cu system that dramatically reduces the flow stresses in both the alloys at temperatures 400 o C. Introduction Compression test data are routinely and successfully utilized to develop hot workability maps for high strength Al alloys. It is the purpose of this study to highlight that a comparison of compressive flow stress data can further provide a basis for realizing the variations in relative hot strengths of the alloys of interest. In the context of the present study, the influence of Sc on various properties of wrought Al alloys is first reviewed as follows in order to establish the scope of the work and to develop a basis for the analysis of the results obtained in this work. The beneficial effects of scandium on the grain refinement of Al and its alloys were established in mid 1960 s in the erstwhile Soviet Union [1]. The studies showed that an addition of about 0.2wt% Sc plays an effective role in refining the as-cast grain structure of various Al alloys [1]. It was established that it is the combined action of Sc and Zr that brings about the considerable grain refinement [1,2]. Combined alloying with Sc and Zr refines the grain structure of casting and weldment, leading to remarkable improvements in the weldability of non-weldable, high strength 7xxx series Al alloys [3-5]. Combined additions of Sc and Zr improve workability and deformability of various wrought Al alloys [6,7]. An important microstructural feature of Al alloys containing both Sc and Zr is the formation of Al 3 Sc x Zr 1-x dispersoids that are fine and more uniformly distributed compared to either Al 3 Sc or Al 3 Zr i.e. when Sc or Zr is added singly [1,3,6,7]. The Al 3 Sc x Zr 1-x dispersoids have proven most effective in pinning grain and subgrain boundaries, thus providing stability to the refined cast grain structure during homogenization [8] and inhibiting recrystallization during mechanical and thermal processing of the alloys post homogenization [3,6,9]. Studies have shown that the recrystallization temperature in these alloys is so high that these alloys remain unrecrystallized up to the temperature of 600 o C [6,9]. Using binary Al-Sc alloys, several investigators have pointed out that Sc is the most potent strengthening element on a per atom basis available for use in Al base alloys. The increase in the 0.2% proof stress per 1 at% Sc addition may reach 1000 MPa [6]. The source of the strengthening is coherent Al 3 Sc dispersoids having L1 2 crystal structure with a lattice parameter that is 1.3% higher than that of the matrix Al [6]. A similar assessment of strengthening due to Al 3 Sc x Zr 1-x dispersoids is yet to be performed. Several recrystallization studies on commercial Al alloys containing both Sc and Zr have shown that the composite Al 3 Sc x Zr 1-x dispersoids are more stable against coarsening as compared to the Al 3 Sc dispersoids in the high temperature range of commercial interest [1,3,6,9]. However, it is not known to All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of Trans Tech Publications, (ID: , Pennsylvania State University, University Park, USA-11/05/16,02:34:37)

2 872 Aluminium Alloys ICAA10 what extent or whether the presence of Al 3 Sc x Zr 1-x dispersoids would impart strength and stability to the microstructure of commercial Al alloys at temperatures 300 o C. It is also not known in what way or whether the morphology of the alloy phases in the constituent Al-Zn-Mg-Cu system is modified by the presence of Sc and contributes to the process of hardening at these high temperatures. It is the purpose of this study to evaluate the influence of scandium addition on the high temperature strength of Al-Zn-Mg- Cu-Zr base alloy 7010 by compression tests under varying deformation conditions. Experimental Procedure The 7010 Al alloy used in this work (i.e. the base alloy) had the composition (wt%) of Al-6.3 Zn- 2.3 Mg-1.5 Cu-0.14 Zr-0.04 Fe-0.02 Si (Fe and Si, associated with the primary Al, are present as impurities) wt% Sc was added to the base alloy in order to examine the influence of Sc addition on the high temperature strength of the alloy. The reason for the selection of the particular alloy composition for the base alloy is that the alloy when suitably processed gives rise to improved combination of mechanical properties and stress corrosion cracking (SCC) resistance [3]. Previous studies further showed that about 0.23wt% Sc addition to the base alloy caused desired refinement of the as-cast grain structure, and produces a uniform and fine distribution of Al 3 Sc x Zr 1-x dispersoids upon homogenization [3,8]. The alloys were produced by conventional ingot metallurgical route. The alloys were prepared from high purity Al, Zn and suitable Al-Cu, Al-Mg, Mg-Zr and Al-Sc master alloys in an induction furnace under argon atmosphere. The as-cast alloys were homogenized in an air-circulating furnace. The homogenized ingots were then scalped in order to remove the oxides formed on the surfaces. The resultant materials were then utilized for machining of the compression test samples having the dimensions of 10 mm dia 15 mm height i.e. having a L/D ratio of 1.5 [10]. For the present work, the reason for the selection of the homogenized samples was to contain only Al 3 Zr and Al 3 Sc x Zr 1-x dispersoids in the starting materials (i.e. base alloy and Sc-bearing alloy, respectively) and to evaluate the strength of such alloys under various deformation conditions. The alloy samples were compressed to 50% of their heights at temperatures of 300, 350, 400 and 450 o C and at strain rates of 10-3, 10-2, 10-1 and 1sec -1. The compression tests were carried out using a 100 kn DARTEC servo hydraulic testing machine. The machine generated load-stroke diagrams were converted to true stress true strain plots. The grain structures of as-cast, homogenized and compression-tested samples were examined using optical microscopy. Transmission electron microscopy (TEM) was employed to examine a combination of the distribution of dislocations, dispersoids and the alloy phases, together with the development of subgrain structure in the alloy samples subjected to the compression tests. Electron diffraction was further employed to evaluate the lattice parameter of the composite Al 3 Sc x Zr 1-x dispersoid and to compare this lattice parameter with the known lattice parameters of Al 3 Zr and Al 3 Sc dispersoids. Specimens for TEM were prepared by electrolytic polishing using 30% nitric acid and 70% methanol at 35 o C. Thin foils were examined on a Philips EM 430T electron microscope operating at 300 kv. In order to facilitate comparison between the transmission electron micropgraphs examined, the micrographs showing the distribution of dispersoids were imaged using two-beam orientations obtained near the <110> Al zone, and all other transmission electron micrographs were imaged in <110> Al orientation. Also, unless otherwise stated, transmission electron micrographs obtained from the samples deformed at different temperatures but at the strain rate of 1sec -1 and strain value of ε =0.6 are presented in this article. This was to examine, compare and discuss the microstructures responsible for attaining the highest values of compressive flow stresses at the deformation temperatures investigated. Results and Discussion Figures 1 and represent optical micrographs showing grain structure in the homogenized base alloy and Sc-bearing alloy, respectively. Figure 1 reveals that the refined as-cast grain structure of the Scbearing alloy is retained after homogenization [8], and it is due to the pinning of grain boundaries by numerous, fine Al 3 Sc x Zr 1-x dispersoids that form in the alloy during homogenization. Figure 1 represents a transmission electron micrograph obtained from the homogenized Sc-bearing alloy. The micrograph was imaged near <110> Al zone using a two-beam orientation. The micrograph does show a uniform and fine distribution of the coherent Al 3 Sc x Zr 1-x dispersoids in contrast to the known,

3 Materials Science Forum Vols inhomogeneous distribution of Al3Zr dispersoids in the base alloy. The coherent precipitates in Figure 1 exhibit Ashby Brown strain contrast [11]. Figure 1(d) represents a selected area electron diffraction pattern (SAEDP) with B = [011] obtained from the fine Al3ScxZr1-x dispersoids present in a similar morphology in a different region of the same thin foil. The superlattice reflections observed in the SAEDP are due to the L12 structure of the fine Al3ScxZr1-x dispersoids with a = nm. The lattice parameter of the composite Al3ScxZr1-x dispersoid is, therefore, smaller compared to that of Al3Sc i.e. a = nm [12], and this is consistent with the atomic diameter of Zr being smaller than that of Sc [13]. It follows that the misfit strain (%) with matrix Al for the composite Al3ScxZr1-x dispersoids (i.e. about 0.5%) is intermediate between those of Al3Zr and Al3Sc dispersoids. While this explains the increased nucleation frequency of the composite Al3ScxZr1-x dispersoids compared to Al3Sc, the explanation for the sluggish nucleation frequency of Al3Zr dispersoids in Al alloys, therefore, appears to be of diffusivity origin [14]. (d) Figure 1. Optical micrographs showing grain structure in homogenized base alloy and Sc-bearing alloy. Transmission electron micrograph showing the distribution of the Al3ScxZr1-x dispersoids in the homogenized Sc-bearing alloy, (d) SAEDP with B = [011] obtained from Al3ScxZr1-x dispersoids present in the homogenized Sc-bearing alloy. a b c d Figure 2. True stress vs. true strain curves for the base alloy at 300oC, 350oC, 400oC and (d) 450oC for strain rates varying from 10-3 to 1 sec-1. a b c d Figure 3. True stress vs. true strain curves for the Sc-bearing alloy at 300oC, 350oC, 400oC and (d) 450oC for strain rates varying from 10-3 to 1 sec-1. Figures 2 through (d) represent true stress vs. true strain curves for different strain rates (ranging from 10-3 to 1sec-1) and temperatures (ranging from 300 to 450oC) for the base alloy. The flow stress increases with increasing strain rates. Also, at all temperatures, the work softening is pronounced at higher strain rates i.e and 1sec-1. For the lower strain rates i.e.10-3 and 10-2 sec-1, the rate of work softening decreased at 300 and 350oC, and it is minimal at 400 and 450oC at the lower strain rates. Figures 3 through (d) show true stress vs. true strain curves for different strain rates & temperatures for the Sc-bearing alloy. The flow stresses due to two successive strain rates often show values close to each other, although in all cases, the peak and steady state flow stresses are registered at values increasing with increasing strain rates. Further, in contrast to the base alloy, the work softening is pronounced at all strain rates up to the temperature of 400oC. At 450oC, the rate of flow softening decreased at least at the lower strain rates.

4 874 Aluminium Alloys ICAA10 The most remarkable feature of the flow stress curves as shown in Figures 2 and 3 is the dramatic high temperature strengthening of the Sc containing alloy as compared to the base alloy. A comparison of flow stress values at ε = 0.5 between the base alloy and the Sc-bearing alloy at various temperature-strain rate combinations pointed out that at all temperatures and strain rates, the flow stress values of Sc containing alloys are 2-3 times greater than those of the base alloy. Further, the peak flow stress values of the Sc-containing alloy at lower temperatures i.e. 300 and 350oC and the strain rate of 10-3 sec-1 reached 46 times those of the base alloy. On the other hand, the significant reductions in the flow stresses for both the alloys at or above 400oC, irrespective of strain rate, are evident in the Figures. Figures 4 and represent optical micrographs showing the grain structure developed in the base alloy and the Sc-bearing alloy, respectively, when deformed at temperature, strain rate combinations of 350oC, 10-3 sec-1. Pancaking of the grains is a noticeable feature, but no sign of any recrystallization could be noticed in the micrographs. The use of increased deformation temperature of 450oC also did not induce recrystallization in the base alloy, and of course, not in the Sc-bearing alloy [Figure 4]. Figure 4. Optical micrographs showing development of grain structures when deformed at temperature-strain rate combinations of 350oC, 10-3 sec-1, base alloy, 350oC, 10-3 sec-1, Sc-bearing alloy, 450oC, 10-3 sec-1, Sc-bearing alloy. Figures 5 and represent transmission electron micrographs showing development of subgrains in the base alloy and Sc-bearing alloy, respectively; both the alloys were deformed at 300oC and at the strain rate of 1sec-1. It was revealed that the process of subgrain formation in the Sc-bearing alloy is relatively sluggish because of the presence of uniformly distributed, fine Al3ScxZr1-x dispersoids. Also, the average size of the subgrains present in the Sc-bearing alloy (a couple marked by arrows) is remarkably small i.e. about 0.5 µm compared to 3.5 µm for the subgrains present in the base alloy. This causes the flow stress to rise in the case of Sc-bearing alloy. Figures 5 and (d) show the development of subgrains in the base alloy, and the Sc-bearing alloy (a couple marked by arrows), respectively when deformed at 400oC and at the strain rate of 1sec-1. With increasing temperature, the subgrain boundaries became clearly defined and the subgrain size increased. In the case of the Sc-bearing alloy, the average size of the subgrains increased to about 1 µm compared to about 4 µm for the base alloy. This increase in subgrain size would contribute toward the reductions in the flow stresses at temperatures 400oC. (d) Figure 5. Transmission electron micrographs showing the development of subgrains in base alloy and Sc-bearing alloy, both deformed at 300oC, 1sec-1, base alloy and (d) Sc-bearing alloy, both deformed at 400oC, 1sec-1. The distribution of dispersoids in the base alloy and the Sc-bearing alloy, both deformed at 450oC and at the strain rate of 1sec-1, is shown in Figures 6 and, respectively. The transmission electron micrographs reveal an inhomogeneous distribution of fewer Al3Zr dispersoids in the base alloy compared

5 Materials Science Forum Vols to the uniform distribution of fine Al3ScxZr1-x dispersoids in the Sc-bearing alloy. Figure 6 reconfirms the greater nucleation frequency of Al3ScxZr1-x compared to Al3Zr in Al alloys, and the stable nature of the dispersoids in the Scbearing alloy at high temperatures. This would, therefore, cause greater strengthening in the Scbearing alloy at 450oC. Yet another noticeable feature of Figure 6 is the apparent absence of coarse alloy phase particles in Figure 6, when compared with Figure 6. This particular aspect is discussed in more details as follows. Figure 6. Transmission electron micrographs showing the distribution of dispersoids in base alloy, and Sc-bearing alloy, when deformed at 450oC, 1sec-1. Figures 7 through (d) present transmission electron micrographs showing the distribution of alloy phase particles, mainly η[mg(zncual)2], formed under different deformation conditions in the alloy samples. The morphology of phase particles formed in the base alloy deformed at the strain rate of 1sec-1 and at temperatures 300 and 450oC is shown in Figures 7 and, respectively. The very considerable reduction in the number density of the phase particles with temperature may be noted. The morphology of the phase particles present in the similarly deformed Sc-bearing alloy is shown in Figures 7 and (d), respectively. Comparison of Figure 7 with Figure 7 shows that the size of the phase particles formed in the Scbearing alloy is an order of magnitude smaller as compared to the phase particles formed in the base alloy. Further, compared to the base alloy, the number density of the phase particles formed in the Sc-bearing alloy at 300oC is considerably higher. Consideration of Figures 7 and (d) shows that the influence of Sc on the morphology of η plates persists in the alloy even when deformed at 450oC. These results demonstrate that Sc addition has the remarkable influence of refining the morphology of the alloy phases (d) Figure 7. Transmission electron micrographs showing the morphology of the alloy phase precipitates formed in base alloy at 300oC, 1 sec-1, base alloy at 450oC,1 sec-1, Sc-bearing alloy at 300oC,1 sec-1, and (d) Sc-bearing alloy at 450oC,1 sec-1. in the constituent Al-Zn-Mg-Cu system that contributes to the superior strengthening of the Sc-bearing alloy observed at temperatures up to 350oC. Further, it is a combination of higher rates of coarsening, and instability of the alloy phases at temperatures 400oC that reduces the flow stresses in both the alloys. The present results may further be discussed in light of the known hot working behavior of Al alloys [15] as follows. The hot working behavior of the present alloys will have influenced by a combination of (1) morphology and the stable nature of the dispersoids, (2) changing type, morphology and stability of the alloy phases due to exposure to different temperature, strain rate and strain combinations, and (3) dynamic recovery through subgrain formation. The hot working begins with a quickly developed microstructure where, the precipitates (alloy phases) are fine and numerous and the solute atmospheres are highly concentrated. Such microstructures are highly sensitive to both temperature and strain changes. With increasing strain, the phase particles coarsen and partly dissolve, giving rise to a flow curve with a characteristic peak and work softening toward a steady state. The more pronounced peaks (and superior hardening) observed in the case of Sc-bearing alloy are now understood to be due to a combination of the formation of remarkably finer morphology of the alloy phase particles, the presence of a greater

6 876 Aluminium Alloys ICAA10 number density of uniformly distributed, fine dispersoids, and the evolution of smaller subgrains. With increasing deformation temperature, the level of hardening and the intensity of the peaks reduce and this is primarily associated with the decreasing number density of the alloy phase particles in the alloys. Further, it should be noted that the formation of subgrains through dynamic recovery contributes to work softening, whilst work softening due to dynamic recrystallization (DRX) is ruled out, because recrystallization was detected neither in the base alloy nor in the Sc-bearing alloy. Summary and Conclusions 1.The influence of 0.23 wt% Sc on the high temperature strength of homogenized Al alloy 7010 was investigated by compression tests. The results demonstrate that Sc addition increases the compressive flow stress. The increase in peak flow stress is nearly 3-6 times the peak flow stress of the base alloy over the temperature range of o C and the strain rate range of 10-3 to 1 sec -1. At temperatures 400 o C, the flow stress decreases considerably irrespective of the strain rate utilized. 2.Compared to the base alloy, the significant increase in flow stress in the Sc-bearing alloy at temperatures o C is attributed to a combination of the formation of remarkably refined alloy phases in the constituent Al-Zn-Mg-Cu system, the presence of a greater number density of uniformly distributed, fine dispersoids, and the evolution of smaller subgrains. 3.At temperatures 400 o C, the morphology of the dispersoids remains unaltered. Whilst, the recovery process progresses with increasing subgrain size, and the number density of the alloy phase particles decreases considerably leading to significant reductions in the flow stresses of Sc-bearing alloy as well as base alloy. Acknowledgements The authors wish to acknowledge the financial support of Defence Research and Development Organisation (DRDO), Government of India. References [1] L.S.Toropova, D.G.Eskin, M.L.Kharakterova, T.V.Dobatkina: Advanced Aluminium Alloys Containing Scandium: Structure and Properties, Gordon and Breach Science Publishers, The Netherlands (1998). [2] V.G.Davydov, V.I.Yelagin, V.V.Zakharov, T.V.Rostova: Metallovedenie I Termicheskaya Obrabotka Metallov, Vol.8 (1996), p. 25. [3] A.K.Mukhopadhyay, G.M.Reddy, K.S.Prasad, S.V.Kamat, A.Dutta, C.Mondol: J.T.Staley Honorary Symposium on Al Alloys (ed. M.Tiryakioglu), Indianapolis, USA, 5-8 November 2001, ASM International (2001), p. 63. [4] A.K.Mukhopadhyay and G.M.Reddy: Mater. Sci. Forum, Vol (2002), p [5] G.M.Reddy, A.K.Mukhopadhyay and A.S.Rao: Sci. and Technol. of Welding and Joining, Vol.10 (2005), p.432. [6] Y.V.Milman, D.V.Lotsko, O.I.Sirko: Mater. Sci. Forum, Vol (2000), p [7] G.I.Eskin: Technologia Legkikh Splavov, Vol. 2 (2000), p.17. [8] K.S.Prasad, A.K.Mukhopadhyay and Vydehi Joshi: Z.Metallkunde, Vol.11 (2004), p [9] B.Forbord, H.Hallem and K.Marthinsen: Proc. ICAA 9 (eds. J.F.Nie, A.J.Morton and B.C.Muddle), Inst. of Mater. Engg. Australasia Ltd. (2004), p [10] ASTM E9-89a, ASTM Annual Book of Standard, (2004), p.110. [11]P.B.Hirsch, A.Howie, R.Nicholson, D.W.Pashley and M.J.Whelan: Electron Microscopy of Thin Crystals, R.E.Krieger Publishing Company, Florida (1977), p [12] R.W.Hyland, Jr., M.Asta, S.M.Foiles and C.L.Rohrer: Acta Mater., Vol. 46 (1998), p [13] W.Hume-Rothery and G.V.Raynor: The structure and Metals of Alloys, The Institute of Metals, Monograph and Report Series No. 1 (1962), p. 92. [14] L.Lae, P.Guyot and C.Sigli: Proc. ICAA 9 (eds. J.F.Nie, A.J.Morton and B.C.Muddle), Inst. of Mater. Engg. Australasia Ltd.(2004), p.281. [15] H.J.McQueen: Hot Deformation of Al Alloys (eds. T.G.Langdon, H.D.Merchant, J.G.Morris and M.A.Zaidi), The minerals, Metals and Materials Society (1991), p.31 and p.105.

7 Aluminium Alloys ICAA / Influence of Scandium Addition on the High Temperature Compressive Strength of Aluminium Alloy /