Improvement of Bake-Hardening Response of Al-Mg-Cu Alloys by means of Nanocluster Assist Processing (NCAP) Technique

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1 Materials Science Forum Online: ISSN: , Vols , pp doi: / Trans Tech Publications, Switzerland Improvement of Bake-Hardening Response of Al-Mg-Cu Alloys by means of Nanocluster Assist Processing (NCAP) Technique Shoichi Hirosawa 1,a, Tomoya Omura 2,b, Yoshikazu Suzuki 3,c and Tatsuo Sato 1,d 1 Dept. of Metallurgy and Ceramics Science, Tokyo Institute of Technology, O-okayama, Meguro-ku, Tokyo , Japan. 2 Graduate student, Tokyo Institute of Technology, Japan (Present: Ricoh Co., Ltd., Japan). 3 Furukawa-SKY Aluminum Corp., Technical Research Division, Uwanodai, Fukaya-shi, Saitama , Japan. a shoichi.hirosawa@mtl.titech.ac.jp, b omura@mtl.titech.ac.jp, c suzuki.yoshikazu@furukawa -sky.co.jp, d sato@mtl.titech.ac.jp Keywords: Al-Mg-Cu alloy, nanoclusters, 3DAP, HRTEM, microalloying, bake-hardening Abstract. In this work, the bake-hardening (BH) response of an Al-3.0Mg-1.0Cu (in mass%) alloy has been improved by the small addition of Ag as a good example of our proposed Nanocluster Assist Processing (NCAP) technique. From the detailed observation through high resolution transmission electron microscopy (HRTEM), it is found that the origin of the increased hardness in the Ag-added alloy is attributed to the densely and uniformly formed Z phase at the expense of Guinier-Preston- Bagaryatsky (GPB) zones and the S phase. It is new findings that the Z phase is formed even in the ternary alloy although the chemical composition lies in the (α+s+t) phase field. Based on the threedimensional atom probe (3DAP) results, furthermore, it is suggested that nanoclusters of Mg, Ag and/or Cu provide effective nucleation sites for the Z phase, whereas nanoclusters of Mg and Cu do less. Such unique characteristics of Ag are clearly seen in the newly constructed interaction energy map (IE map). Introduction Non-heat treatable Al-Mg alloys have been utilized as a light-weight body sheet material for automotives due to their well-balanced properties of mechanical strength, ductility and pressformability [1]. Their inferior strengths to those of heat treatable Al-Mg-Si alloys can be increased by the addition of copper because Al-Mg-Cu alloys give rise to the precipitation strengthening phenomena [2]. This improvement partially meets the requirement of manufacturing processes in which good age-hardenability during a paint-bake treatment at ~443K, termed as a bake-hardening (BH) response, is favorable to prevent bake-softening occurring in Al-Mg alloys. However, a more increased BH response without diminishing good formability is fascinating, and therefore the continuous development in chemical compositions for new alloys has been undertaken. The small addition (e.g. 0.1mol%) of some alloying elements has been recognized to be effective to control precipitate microstructures and the resultant alloy properties. Ringer et al. [3, 4] reviewed the effects of microalloying additions in Al-Cu(-Mg) and Al-Zn-Mg alloys and summarized their wide range of atom probe work, showing the experimentally clarified distribution of solute and microalloying elements. Furthermore, the present authors [5, 6] have theoretically predicted the atomic behaviors of microalloying elements in Al-Zn, Al-Cu(-Mg) and Al-Mg-Si alloys in the form of a scatter plot, in which interatomic interaction energies among solutes and/or vacancies are mapped out for various microalloying elements. The proposed process termed as the Nanocluster Assist Processing (NCAP) [6] is based on the control of nano-scale clusters (nanoclusters) by means of the addition of preferably predicted microalloying elements. In this work, the BH response of an Al-3.0Mg-1.0Cu (in mass%) alloy has been improved by the small addition of Ag as a good example of our proposed NCAP technique. The origin of the hardness increase during the BH treatment is verified from the experimentally obtained nano-scale 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-06/03/16,10:33:25)

2 216 Aluminium Alloys ICAA10 microstructures through high resolution transmission electron microscopy (HRTEM) and a threedimensional atom probe (3DAP) analysis. Employing a first-principles calculation based on a fullpotential Korringa-Kohn-Rostoker (KKR) Green s function method [7, 8], more accurately estimated interaction energy map is constructed to represent the unique characteristics of Ag with Mg and Cu. Experimental Procedure The chemical compositions of the alloys used in this work are listed in Table 1. The cold-rolled sheets were solution-treated in a salt bath at 793K for 25s, followed by iced-water quenching. An aging treatment was carried out in an oil bath at 443K for various times. Micro-Vickers hardness was measured with a load of 500g. The foils for TEM observation were prepared by the electrolytic twin-jet polishing technique in 20vol% nitric acid with 80vol% methanole at ~253K. TEM microstructures were observed using a JEM3010 microscope at an accelerating voltage of 300kV and HRTEM images along the [001] α direction were taken by the many-beam tecnique under the axial illumination. The chemical analysis of precipitated particles was performed with a probe size of ~15nm by an energy-dispersive X-ray (EDX) spectrometer attached to the microscope. Needleshaped samples for the 3DAP analysis were prepared by electropolishing in 25vol% perchloric acid with 75vol% acetic acid and in a solution of 2vol% perchloric acid in 2-butoxyethanol. The 3DAP analysis was carried out at a specimen temperature of 20K with a pulse fraction of 20% under the UHV condition (<10-10 mbar) by an energy-compensated atom probe with a delay-line detector. Table 1 Chemical compositions of the alloys used in this work [mass%]. Alloy Mg Cu Ag Si Fe Ti Al Ternary bal. Ag-added bal. Experimental Results BH Response The isothermal aging curves of hardness at 443K are shown in Fig.1 for the Al-Mg-Cu ternary and Ag-added alloys. The two-stage hardening behavior, similar to those in Al-Cu-Mg(-Ag) alloys with high Cu-to-Mg ratios [3, 9-11], is observed in the two alloys; i.e. the first rapid hardening within 0.1ks and the subsequent hardening after a plateau. The BH response of the alloys can be estimated as the difference in hardness between as-quenched (A.Q.) and aged for (e.g.) 1.2ks, and therefore the addition of Ag is found to improve the BH response significantly; e.g. from HV16 to HV29. Fig.1 Isothermal aging curves of hardness at 443K for the Al-Mg-Cu ternary and Ag-added alloys.

3 Materials Science Forum Vols Precipitate Microstructures The corresponding TEM images along the [001]α direction are shown in Fig.2. The microstructures at the plateau stage of hardness (not shown here) exhibit no clear evidence of any precipitates contributing to the first rapid hardening. On the other hand, the microstructures at the subsequent hardening stage consist of rod-shaped Guinier-PrestonBagaryatsky (GPB) zones, the lath-shaped S (or S) phase and equiaxed precipitates (Fig.2). In the ternary alloy, a large amount of the S phase are heterogeneously formed on dislocations and a quite limited number of equiaxed precipitates (shown by arrowheads) coexist in the matrix, suggesting that the primary strengthening phase at the second hardening stage is GPB zones. In contrast, the small addition of Ag drastically changes the formation rate, spatial distribution and volume fraction of these precipitates; i.e. a substantial amount of smaller equiaxed precipitates at the expense of GPB zones and the S phase. The EDX analysis of the peak-aged Ag-added alloy (Fig.2(d)) confirms that Ag is incorporated in the equiaxed precipitates with a Mg:Cu:Ag ratio of approximately 6:3:1. Therefore, a remarkable microstructural change by microalloying Ag is the origin of more rapid and increased hardening in the Ag-added alloy (Fig.1). Note that the terminology of microalloying originates not only in the trace amount but also in no or slight effect on the phase diagram regions [5, 6]. In this work, more detailed observation has been performed for the equiaxed precipitates. Fig.3(a) and (d) shows typical HRTEM images of equiaxed precipitates and the surrounding matrix in the Al-Mg-Cu ternary and Ag-added alloys aged at 443K for 345.6ks. The precipitates are distinguished from the matrix by the different arrangement of dots, and therefore assigned ~20nm in size for the ternary alloy and ~10nm for the Ag-added alloy, respectively. It is also seen that the interface between the precipitate and the matrix is coherent along the [100]α and [010]α directions, independent of the existence of Ag. Furthermore, although systematic changes in contrast are not necessarily observed, the arrangement of dots within the precipitates appears to exhibit some periodicity, as shown by the squares in the magnified figures (Fig.3(b) and (e)). This periodicity is confirmed by the corresponding FFT spectra (Fig.3(c) and (f)) where intensified reflections from individual precipitates locate at almost every tenth between fundamental 220 spots. This strongly suggests that the precipitates possess five times the length of the lattice constant of the Al matrix. S S S GPB Fig.2 TEM micrographs with diffraction patterns for the Al-Mg-Cu ternary ((a), (c)) and Ag-added ((b), (d)) alloys aged at 443K for ((a), (b)) and 1209ks ((c), (d)).

4 218 Aluminium Alloys ICAA10 Fig.3 HRTEM micrographs and FFT spectra obtained from individual equiaxed precipitates in the Al-Mg-Cu ternary ((a)-(c)) and Ag-added ((d)-(f)) alloys aged at 443K for 345.6ks. Discussion Identification of Precipitated Phases According to the phase diagram of the Al-Cu-Mg system, the ternary alloy used in this work lies in the (α+s+t) phase field at 463K [9]. On the other hand, the Ag-added alloy is expected to produce the recently designated Z phase, which had been considered as the T phase, based on a result of Chopra et al. for a similar Al-Mg-Cu-Ag alloy with a high Mg-to-Cu ratio [12]. The structural comparison between the T and Z phases is listed in Table 2 [12, 13]. The HRTEM microstructures observed in this work suggest that the equiaxed precipitates have a cubic structure with a lattice parameter of ~2nm, five times of the lattice constant of the Al matrix, independent of the existence of Ag (Fig.3). The periodicity is satisfied if the Z phase is formed with a cube-on-cube orientation relationship with the matrix; i.e. (100) z //(100) α and [010] z //[010] α. This assumption would be justified from the fact that the positions and intensities of reflected spots from the individual precipitates (Fig.3(c) and (f)) are almost identical to those of a convergent-beam electron diffraction pattern from the coarsened Z phase with the same orientation relationship (Fig.2(c) in [12]). In contrast, the obtained FFT spectra (Fig.3(c) and (f)) do not rationally match the diffraction pattern from the T phase, which gives rise to faint spots near 1/3 220 with superimposed double diffractions at 1/3 and 2/3 220 if viewing along the [001] α direction. The coherent interface between the precipitate and the matrix is also explained by the small misfit arising from the difference in lattice parameters between the Z phase and the Al matrix (i.e. ~1.0%). Therefore, it can be concluded that the Z phase is formed even in the ternary alloy and the Ag addition stimulates the formation of the Z phase at the expense of competitively formed GPB zones and the S phase. Phase Table 2 Structural comparison between the T and Z phases. Crystal structure Point group Lattice constant [nm] Number of atoms per unit-cell T - Mg 32 (Al,Cu) 49 bcc Im3 _ [13] Z - unknown fcc Fm3m unknown [12] Ref.

5 Materials Science Forum Vols (a) (b) Fig.4 3DAP maps of (a) Al-Mg-Cu ternary and (b) Ag-added alloys aged at 443K for 3.6ks. Only Mg, Cu and Ag atoms are depicted for better visibility; i.e. (a)mg+cu, (b)mg+cu+ag and (c)ag. Origin of Hardening during BH Treatment The age-hardening behavior of the investigated alloys exhibits two-stage increase in hardness; i.e. the first rapid hardening within 0.1ks and the subsequent hardening after a plateau (Fig.1). The remarkable change in TEM microstructures (Fig.2) well explains the second hardening stage, as firstly suggested by Vietz et al. [9], but the origin of the first rapid hardening is still controversial [3, 10, 11, 14-17]. In this work, the 3DAP analysis has been applied to the alloys at the plateau stage, where no clear evidence of any precipitates contributing to the first rapid hardening had been observed by TEM. Fig.4 shows 3DAP maps of the Al-Mg-Cu ternary and Ag-added alloys aged at 433K for 3.6ks. Only Mg, Cu and Ag atoms are depicted in the 3DAP maps for better visibility. As opposed to the nearly homogeneous distribution of Mg and Cu atoms in the ternary alloy (Fig.4(a)), small clusters with a diameter of 3-4nm (nanoclusters) are clearly observed in the Ag-added alloy (Fig.4(b)). The quantitative estimation of chemical compositions reveals that Ag atoms are incorporated into the nanoclusters with Mg:Cu:Ag ratios of 6.7:1.1:2.2 for 1.8ks aging (not shown here) and 6.4:1.8:1.8 for 3.6ks aging (Fig.4(b)). Therefore, by comparing with the chemical compositions of the Z phase; e.g. Mg:Cu:Ag = 6:3:1 (Fig.2(d)) and Al-(20-25)at.%Mg-20at.%Cu -(~2-5)at.%Ag [18], it is found that nanoclusters of Mg, Ag and/or Cu atoms form first, and then develop into the Z phase at the second hardening stage. Although the similar compositional analysis has not been performed for the ternary alloy, it is plausible to consider that nanoclusters of Mg, Ag and/or Cu provide effective nucleation sites for the Z phase, whereas nanoclusters of Mg and Cu do less. Interaction Energy Map for NCAP The effectiveness of microalloying Ag is well interpreted by the interatomic interaction energies with both Mg and Cu. Fig.5 shows our newly constructed interaction energy map (IE map; N.B. previously named as OP map [5, 6]) for various microalloying elements X in Al-Mg-Cu alloys. This revised IE map graphically compares the two-body interaction energies in Al between Mg-X (E II Mg-X) and between Cu-X (E II Cu-X), which are all estimated not by a previously utilized regular solution approximation [5, 6], but by a first-principles calculation based on a full-potential KKR Green s function method [7, 8]. Although no effect of misfit strain arising from the difference in the atomic sizes is taken into account, it is predicted from the systematic changes in E II Mg-X and E II Cu-X that only seven elements including Ag possess attractive interactions with Mg and Cu simultaneously (i.e. in the case of E II Mg-X <0 and E II Cu-X <0). Among them, however, only Ag has a large solubility limit in Al enough to obtain the super-saturated solid solution, whereas the others are practically insoluble in Al, restricting the potential atomic behaviors as microalloying elements. Therefore, it is supposed that the two unique characteristics of Ag; i.e. simultaneous interactions with Mg and Cu together with the large solubility limit in Al, are the origins of the densely and uniformly formed Z phase through a heterogeneous nucleation effect by the nanoclusters of Mg, Ag and/or Cu. Note that such a NCAP technique can be applied to other alloy systems; e.g. Al-Cu(-Li) + Mg, Al-Cu-Mg + Ag, Al-Zn +Ag, Al-Zr + Si and so on [6].

6 220 Aluminium Alloys ICAA10 Fig.5 Interaction energy map (IE map) for various microalloying elements X in Al-Mg-Cu alloys, showing the two-body interaction energies in Al between Mg-X atoms (E II Mg-X) and between Cu-X atoms (E II Cu-X). Summary (1) The BH response of an Al-3.0Mg-1.0Cu (in mass%) alloy can be improved by the small addition of Ag due to the remarkable change in precipitate microstructures from GPB zones and the S phase to a substantial amount of the equiaxed Z phase. (2) The formation of nanoclusters of Mg, Ag and/or Cu is responsible for the densely and uniformly formed Z phase through the heterogeneous nucleation effect (i.e. the NCAP technique). (3) The IE map proposed in this work is quite useful to show and identify which are likely to behave favorably as microalloying elements in a given alloy system. Acknowledgements The authors gratefully thank Prof. T. Hoshino, Shizuoka University, and Dr. F. Nakamura, Tokyo Institute of Technology, for the computational assistance. This study has been conducted as Nanotechnology Metal Project supported by New Energy and Industrial Technology Development Organization (NEDO) and The Japan Research and Development Center for Metals (JRCM). References [1] For example, J.Hirsch: Proc. 9th Int. Conf. on Aluminum Alloys (2004), p.15. [2] Y.Suzuki, M.Matsuo, M.Saga and M.Kikuchi: Materials Science Forum Vol (1996), p [3] S.P.Ringer and K.Raviprasad: Materials Forum Vol.24 (2000), p.59. [4] I.J.Polmear and S.P.Ringer: J. Japan Inst. Light Metals Vol.50 (2000), p.633. [5] S.Hirosawa, T.Sato, A.Kamio and H.M.Flower: Acta mater. Vol.48 (2000), p [6] T.Sato, S.Hirosawa, K.Hirose and T.Maeguchi: Metall. Mater. Trans. Vol.34A (2003), p [7] T. Hoshino and F. Nakamura: J. Metastable and Nanocrystalline Materials Vols (2005), p.237. [8] F.Nakamura, T.Hoshino, S.Tanaka, K.Hirose, S.Hirosawa and T.Sato: Trans. Materials Research Society of Japan (2006), in print. [9] J.T.Vietz and I.J.Polmear: J. Inst. Metals Vol.94 (1966), p.410. [10] S.P.Ringer, T.Sakurai and I.J.Polmear: Acta mater. Vol.45 (1997), p [11] L.Reich, S.P.Ringer and K.Hono: Phil. Mag. Lett. Vol.79 (1999), p.639. [12] H.D.Chopra, B.C.Muddle and I.J.Polmear: Phil. Mag. Lett. Vol.73 (1996), p.351. [13] J.H.Auld and B.E.Williams: Acta Cryst. Vol.21 (1966), p.830. [14] A.M.Zahra, C.Y.Zahra, C.Alfonso and A.Charai: Scripta Materiala Vol.39 (1998), p [15] S.P.Ringer, S.K.Caraher and I.J.Polmear: Scripta Materiala Vol.39 (1998), p [16] P.Ratchev, B.Verlinden, P.De Smet, P. Van Houtte, Acta mater. Vol.46 (1998), p [17] L.Kovarik, P.I.Gouma, C.Kisielowski, S.A.Court and M.J.Mills: Mater. Sci. Forum Vols (2002), p [18] S.P.Ringer, G.C.Quan and T.Sakurai: Mater. Sci. Eng. A Vol.250 (1998), p.120.

7 Aluminium Alloys ICAA / Improvement of Bake-Hardening Response of Al-Mg-Cu Alloys by Means of Nanocluster Assist Processing (NCAP) Technique / DOI References [7] T. Hoshino and F. Nakamura: J. Metastable and Nanocrystalline Materials Vols (2005), p /