The effects of trace Sc and Zr on microstructure and internal friction of Zn Al eutectoid alloy

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1 Materials Science and Engineering A 370 (2004) The effects of trace Sc and Zr on microstructure and internal friction of Zn Al eutectoid alloy B.H. Luo, Z.H. Bai, Y.Q. Xie Department of Materials Science and Engineering, Central South University, Changsha , PR China Received 12 July 2002 Abstract The damping properties of Zn 22 wt.% Al alloys without and with Sc (0.55 wt.%) and Zr (0.26 wt.%) were investigated. The internal friction of the determined by the microstructure has been measured in terms of logarithmic decrement (δ) using a low frequency inverted torsion pendulum over the temperature region of C. An internal friction peak was separately observed at about 218 C in the Zn Al alloy and at about 195 C in Zn Al Sc Zr alloy. The shift of the δ peak was found to be directly attributed to the precipitation of Al 3 (Sc, Zr) phases from the alloy matrix. We consider that the both internal friction peak in the alloy originates from grain boundary (GB) relaxation, but the grain boundary relaxation can also be affected by Al Sc Zr intermetallics at the grain boundaries, which will impede grain boundary sliding. In addition, Al Sc Zr intermetallics at the grain boundaries can pin grain boundaries, and inhibit the growth of grains in aging, which increases the damping stability of Zn 22 wt.% Al alloy Published by Elsevier B.V. Keywords: Internal friction; Zn Al eutectoid alloys; Sc; Zr 1. Introduction Zn 22 Al eutectoid alloy has much theoretical and practical value. Experiments [1 3] have confirmed that the damping in this alloy originates from its phase-interface motion, i.e. the high damping of the alloy is produced by the grain boundary (GB) sliding which can be explained by the diffusive flux on a boundary between phases like /, / and /. But there is a problem, even at room temperature phase (Al-rich phase) and phase (Zn-rich phase) in the alloy coarsen with time, which can cause a change of the alloy s damping capacity, or internal friction of the alloy decreases with time [4]. The elements B, Ti, Zr, La, Ce can improve the alloy s mechanical properties and corrosion-resistance at a large scale [5], but they can not increase the alloy s damping stability. Researches [6 8] show that Sc in aluminum alloys is an element that can disperse strengthen, refine grains and prohibit recrystalization. Aluminum alloys with trace Sc have high strength, plasticity, thermal-resistance, corrosion-resistance Corresponding author. Tel.: ; fax: address: lbh@mail.csu.edu.cn (B.H. Luo). and good weldability. Simultaneous adding trace Sc and Zr in aluminum alloys can improved its mechanical properties greatly. In this paper, the damping capacity and damping stability with time are studied and has the Zn 22 wt.% Al 0.55 wt.% Sc 0.26 wt.% Zr alloy, and compared with the Zn 22 wt.% Al eutectoid alloy. 2. Experimental procedure Aluminium and zinc of wt.% purity and Al Sc, Al Zr master alloys were used to produced Zn Al alloy ingots by molten and water-cooled casting in air. The sample s analyzed composition is as follows: 1 #, Zn 22 wt.% Al; 2 #, Zn 22 wt.% Al 0.55 wt.% Sc 0.26 wt.% Zr. After homogenization at 350 C for 8 h, the ingots were rolled to 4 mm thick at 270 C. After that the samples were soluted treat in the solid state at 370 C for 1 h, and then quenched into water at the temperature of 10 C. Finally, the samples were rolled to 1 mm thick at 80 C, with a total reduction of 75%. The rolled pieces of 1 # and 2 # alloys were then cut into 60 mm 2.5mm 1 mm specimens used in the experiment. The damping value are measured at the elevated temperature on a low frequency inverted torsion /$ see front matter 2003 Published by Elsevier B.V. doi: /j.msea

2 pendulum after each quenched and completely decomposed specimens is preannealed at 240 C for 1 h. This is done in order to eliminate the nonequilibrium quenching interstress and the probable structure changes due to the nonequilibrium interphase solute atom content accommodation. The internal friction curves were obtained at a heating rate of 5 K/min, a strain amplitude of and a frequency of 1 Hz. Microstructure characteristic of the alloy was investigated with optical microscope (OM), D500 X-ray diffractometer and H-800 transmission electron microscope (TEM). B.H. Luo et al. / Materials Science and Engineering A 370 (2004) Experimental results 3.1. Microstructure Fig. 1 shows the microstructure of differently treated alloys. Fine equiaxed grains of and phases are observed. The light and dark phases are and phases, respectively. Fig. 1(c) indicates the microstructure of the as-cast alloy with Sc and Zr addition after etching by mixed acid, from which small square or trigonal particles of the second phases can be found at grain boundaries and within some grains, identified as Al 3 (Sc, Zr) composite particles from X-ray diffraction results (Fig. 2), which agree with that of reference [6]. It is shown that simultaneous adding of trace Sc and Zr can obviously refine the grain size. In Fig. 2. XRD patterns of Zn 22 wt.% Al alloy (1 # alloy) and Zn 22 wt.% Al 0.55 wt.% Sc 0.26 wt.% Zr alloy (2 # alloy) (a) 2 # alloy, 50 h aging at room temperature; (b) 1 # alloy, 50 h aging at room temperature. addition, after 370 C/h solution treatment and then 80 C rolling treatment, both alloys have smaller equiaxed grains with superplastic characteristics. Fig. 3 presents the TEM microstructures of the two studied alloys at final treated conditions. In these two studied alloys, dislocations are rarely discovered. Comparing Fig. 3(a) with Fig. 3(b) and Fig. 1. Optical microstructure of differently treated alloys (a) 1 # alloy, as-cast; (b) 1 # alloy, as-rolled, after solution treatment; (c) 2 # alloy, as-cast; (d) 2 # alloy, as-rolled, after solution treatment.

3 174 B.H. Luo et al. / Materials Science and Engineering A 370 (2004) Fig. 3. TEM microstructures of two alloys (a) 1 # alloy, 50 h aging at room temperature; (b) 1 # alloy, h aging at room temperature; (c) 2 # alloy, 50 h aging at room temperature; (d) 2 # alloy, h aging at room temperature. Fig. 3(c) with Fig. 3(d), the equiaxed grains of and phases in Zn Al alloy coarsen with the time at room temperature. But in Zn Al Sc Zr alloy, the grain coarsening does not occur, and Al 3 (Sc, Zr) particles mainly distribute at grain boundaries and partially spread dispersively within the grains Internal friction Internal friction of Zn Al eutectoid alloy and Zn Al Sc Zr alloy The two alloys relation of internal friction and temperature in the region of C for rolled alloys after Fig. 4. Internal friction vs. temperature of both alloys (f = 1 Hz).

4 B.H. Luo et al. / Materials Science and Engineering A 370 (2004) Fig. 5. Internal friction vs. time (h) of both alloy (f = 1 Hz). solution treatment are shown in Fig. 4. Simultaneous adding of trace Sc and Zr to Zn Al eutectoid alloy can obviously increase internal friction of alloys at the temperature of C. A broad internal friction peak superimposed on a large high-temperature background is observed at temperatures about 218 C and about 195 C, respectively, in the two alloys. On the other hand, the internal friction peak height of Zn Al Sc Zr alloy is lower than that of Zn Al eutectoid alloy The effect of aging time at room temperature on internal friction stability Fig. 5 shows internal friction v s aging time at room temperature of both alloys. At the beginning, the two alloys have similar change tendency, and the internal friction shows little change. And then alloys internal friction decreases with the elapsed-time. But the internal friction of sample 1 # decreases more than that of sample 2 #, and the internal friction of sample 2 # is almost unchanged when aging time exceeds h at room temperature, which shows that adding trace Sc and Zr can improve the stability of Zn Al eutectoid alloy s internal friction. 4. Discussion Adding trace Sc and Zr can strongly refine Zn Al alloy s as-cast microstructure (as shown in Fig. 1). The researches [8] suggested minor Sc and Zr mainly exist as two kind of Al 3 (Sc, Zr) intermetallic compound particles. One is primary Al 3 (Sc, Zr) precipitated forms from the melt during solidification. Primary Al 3 (Sc, Zr) is an ideal crystal nucleus and can greatly decrease the grain size of as-cast alloys. The other is secondary Al 3 (Sc, Zr) precipitated during homogenization, which is small, dispersive, and coherent to the matrix. In this study, Al 3 (Sc, Zr) particles can be seen in Fig. 3(c). The Al 3 (Sc, Zr) particles strongly pin grain boundaries, which effectively restrain migration of boundaries (as shown in Fig. 3(c) and (d)). However, the alloy without the adding of Sc and Zr gradually recrystallized after solution treatment (as shown in Fig. 3(a) and (b)), and phases are coarsened with time elapsed at room temperature. The studies [1 3] show that superplastic and high damping behaviour of Zn Al eutectoid alloy are attributed to the structure of the alloy, which is composed of two phases, i.e. Al-rich phase (soft phase) and Zn-rich phase (hard phase). In addition, the damping of the multiphase Zn Al alloy originates from the phase-interface linear viscous motion. The finer the grain of Zn Al eutectoid alloy is, the higher the alloy s internal friction. Therefore, trace Sc and Zr can increase the alloy s damping property. The Zn Al eutectoid alloy has a high internal friction peak (δ max is about 1.2) at 218 C( 0.6T m ) as shown in Fig. 4. There are numerous examples of strong thermally activated relaxation peaks at around 0.5T m in many polycrystalline metals and alloys [9,10]. These have usually been interpreted in terms of stress relaxation at the grain boundaries, either due to boundary sliding [11] or boundary-migration [12] mechanism. It is well known that an internal friction peak is observed at 290 C in polycrystalline aluminum, and this peak results from grain boundary relaxation [13]. The structure of the studied alloys remains in equiaxed grains and has low dislocation density after superplastic deformation (as shown in Fig. 3). Therefore, a GB sliding process is dominant during superplastic deformation [14,15]. It is suggested that the damping peak in the present alloys also originates from a reversible grain boundary sliding mechanism. Simultaneously we consider that internal friction peak of the Zn Al Sc Zr alloy with a height lower than that of the Zn Al eutectoid alloy and occur at a lower temperature is also related to grain boundary relaxation. This peak is affected by Al 3 (Sc, Zr) particles at the grain boundary. Al 3 (Sc, Zr) particles have high thermal stability [16], and these particles do not coarsen and can strongly pin grain boundaries. They block grain boundary sliding and, as a result, the internal friction peak of the Zn Al Sc Zr alloy is reduced and shifted to lower temperature. The improved internal friction stability of Zn Al eutectoid alloy is also attributed to the addition of trace Sc and Zr, which makes the alloy s microstructure more stable. Zn Al eutectoid alloy s damping value decrease with elapsing time (as shown in Fig. 5). This is caused by the growth of subgrains and coarsing of grains of phase and phase [4]. Al 3 (Sc, Zr) particles have surprising high thermal stability, they can strongly pin grain boundary, and prohibit the formation, growth and merging of subgrains, which stabilize alloy s damping property. Trace Sc and Zr added in Zn Al can increase internal friction of Zn Al alloy at room temperature, and can stabilize the alloy s damping property at a high level is rarely discovered in the history of Zn Al alloys, this opens up a new way of making alloys of high and stable damping value. 5. Conclusions (i) Trace Sc and Zr can obviously refine the microstructure of Zn Al eutectoid alloys. Thermally stable Al 3 (Sc,

5 176 B.H. Luo et al. / Materials Science and Engineering A 370 (2004) Zr) intermetallic phases at the grain boundaries can effectively inhibit the coarsen of the grains. (ii) Zn Al eutectoid alloys has high damping value. Simultaneous adding of trace Sc and Zr to Zn Al eutectoid alloy can increase the alloys damping value in the temperature region of C, and makes the peak temperature of internal friction shift from 218 to 195 C. It is considered that the origin of the internal friction peaks observed in the alloys can be explained by grain boundary relaxation, which is affected by Al 3 (Sc, Zr) intermetallic compound particles at the grain boundaries. The change of peak position and peak height is also attributed to the effect of Al 3 (Sc, Zr) intermetallic compound particles, which block grain boundary sliding. (iii) Zn 22 wt.% Al 0.55 wt.% Sc 0.26 wt.% Zr alloy has higher internal friction stability than that of Zn 22 wt.% Al alloy at room temperature, which is attribute to Al 3 (Sc, Zr) particles in the alloys pining grain boundaries and blocking the coarsen of the grains. Acknowledgements This work was supported by Laboratory of Internal Friction and Defects in Solids at Chinese Academy of Sciences in Hefei, as well as by professor Zhu Jingsong from Nanjing university. References [1] K. Nuttall, J. Inst. Metals 99 (1971) 266. [2] X. Zhu, J. Appl. Phys. 67 (12) (1990) [3] Y. Torisaka, S. Kojima, Acta Metal. Mater. 39 (5) (1991) 947. [4] M. Hinai, S. Sawaya, H. Masumoto, J. Jpn. Inst. Metals 55 (6) (1991) 715. [5] S. Pang, Changsha, Masters dissertation, Central South University, 1994, p. 10. [6] I.I. Velichko, G.V. Dodin, Int Cont: Scandium and Prospects of its Use, Moscow, 1994, p. 14. [7] B.A. Parker, Alloys, J. Mater. Sci. 30 (1995) 452. [8] V.G. Davydov, Mater. Sci. Eng. A 280 (2000) 30. [9] K.M. Entwistle, in: B. Chalmers, A.G. Quarrell (Eds.), The Physical Examination of Metals, Edward Arnold, London, 1960 (Chapter X). [10] A.S. Nowick, Prog. Metal Phys. 4 (1953) 1. [11] T.S. Ke, Phys. Rev. 71 (1947) 533. [12] G.M. Leak, Proc. Phys. Soc. 78 (1961) [13] T.S. Ke, P. Cui, C.M. Su, Phys. Status Solidi A 84 (1984) 157. [14] R.C. Gifkins, J. Inst. Metals 95 (1967) 373. [15] D. Lee, Acta Metals 17 (1969) [16] K. Yu, W.X. Li, J. Nonferrous Metals 9 (1999) 709 (in Chinese).