MANUFACTURING AND EVALUATING CU-BASED SHAPE MEMORY ALLOY BY HOT EXTRUSION OF PM SAMPLES MADE BY MECHANICAL ALLOYING

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1 MANUFACTURING AND EVALUATING CU-BASED SHAPE MEMORY ALLOY BY HOT EXTRUSION OF PM SAMPLES MADE BY MECHANICAL ALLOYING Sajjad Pourkhorshidi, Mohammad Naeimi, Nader Parvin, Seyed Mohammad Mahdi Zamani, Hamid Ebrahimnia Department of Mining and Metallurgical Engineering, Amirkabir University of Technology, Hafez Ave, Tehran, Iran Abstract: Powder Metallurgy method is used to produce Cu-Based Shape memory alloys with maximum desired properties. It is simpler to use P/M to produce near-net shape alloy products and give better controllability of the composition and grain sizes. In this study, ball milling process is applied to convert the elemental powder mixtures of Cu, Al, and Ni into pre-alloyed powders. The effect of milling time on composition and particle size distribution is studied, and the pre-alloyed powders are compacted by pressing and then hot extrusion to obtain the final alloy sample. Strain recovery of bended samples is studied to evacuate the effect of milling on shape memory properties. It has shown that milling time has a variable influence on density of samples, which can affect the shape memory recovery due to amount of porosity. Keywords: shape memory alloy; powder metallurgy; mechanical alloying; strain recovery; particle size distribution 1. INTRODUCTION: Because of low price and high recovery force (only secondary to Ni Ti alloy), Cu-based shape memory alloys are the most promising in practical use. Among the Cu-based shape memory alloys, Cu Al Ni alloy has higher thermal stability than that of Cu Zn Al alloy[1].this system therefore being selected in the present study. For the preparation of SMAs, the conventional casting method has difficulties in controlling the grain size. Coarse grains will weaken mechanical properties of alloys, which is related to their large elastic anisotropy and large grain size [2]. Moreover, the composition change during casting can shift the transformation temperature. It has been reported that mechanical alloying(ma) and powder metallurgy (P/M) with hot isostatic press (HIP) can be used to fabricate Cu-based SMAs[3-6]. P/M can reduce the hot working processes in fabricating the near-net shape products and usually give better control of grain size[7,8]. MA produces pre-alloyed powders which can shorten the sintering time. However, no published work has been reported on the preparation of CAN SMA by MA and the conventional P/M with cold compaction and hot extrusion. In this study, high energy planetary ball milling was applied to convert the Cu-Al-Ni elemental powder mixture to pre-alloyed powders. Conventional P/M with cold compaction was used to produce bulk Cu- Al-Ni based SMA from the pre-alloyed powders by MA. The purpose of this study is to reveal the effect of MA on the microstructure and properties of Cu-AI-Ni SMAs produced by general milling. pressing and furnace equipment in a materials laboratory. 2. EXPERIMENTAL PROCEDURE: A Cu-Al-Ni Alloy powder was prepared by mechanical alloying of elemental powders in a planetary mill for several durations. In MA process, a Fritsch-puluarisate planetary ball mill with stainless steel vial was used. The detail of milling process is shown in Table 1.

2 Table 1. Detailed milling processing of the mixture powder The powder mixtures were mechanically alloyed for 5, 10, 15 and 20 hours. The specification of elemental powders and initial powder mixture is shown in Table 2. The compaction process consisted of a 2 ton handoperated hydraulic press and a single-acting piston die of 13mm bore diameter. The pre-alloyed powders were compacted at a pressure of 900 MPa in the die to form cylindrical-shaped green compacts. The sintering procedure of green compacts was performed in a tube furnace at 900 C for 3 hrs under a protective Hydrogen gas atmosphere. followed by furnace cooling. Sintered compacts were hot-extruded at 700 C with an extrusion ratio of 10:1. For some samples, an annealing has done in 950 C in 1, 5 and 10 hours following quenching in 25 C water. For characterization of particle size distribution, sieving method according to ASTM E11 was used, and density measurement of extruded compacts was according to ASTM designation X-ray diffraction study was made using Philips-pw 1140/90 with Cu K α radiation. The morphology of the hot extruded samples was observed using Philips-XL30 scanning electron microscope (SEM) and OLYMPUS BH-2 optical microscope. Differential Scanning Calorimetric (DSC) test has done by TA-Instrument DSC2010 with 20 C/min heating and cooling rate. Shape memory effect testing was performed by bending rods with diameter of 1mm made by machining extruded samples around a 2mm in diameter holder. The maximum deformation strain ε, was determined as 33% for D=2mm. Table 2. Specification of elemental powders 3 RESULTS AND DISCUSSIONS 3.1 The structural evolution during mechanical alloying: In Figure 2 the structural evolution during mechanical alloying is represented. The pattern of the 1hr MA curve was regarded as a reference condition where the peaks of starting components Cu, Al, Ni appeared. The peaks of Al and Ni were lower than that of Cu because of their small amount in the overall composition. As the MA time increased, the intensity of all the diffraction peaks was decreased and became broader. After MA for 15hrs, the intensities of the Ni peak and most of the Al peaks became almost indistinguishable. A slight shift in the position of the Cu peak to smaller angles, indicated that there is a diffusive process of the Al and Ni in the Cu matrix. After MA for 20hrs, the shift of the Cu peak was more significant and only the peaks of a single phase with FCC structure appeared. The lattice parameters of these phases were found to be close to that of the Cu element. The X-ray patterns, which was in a good agreement with the observation by

3 Kancyoshi et. Al [9], lead to this conclusion that a FCC single phase solid solution was formed after 20hr MA of the elemental powder mixtures.. Fig 2. XRD Patterns of samples mechanically alloyed for different times Figure 3 illustrates the changes in the X-ray pattern of Cu-Al-Ni quenched compact fabricated from powder mixture milled for 15hr, extruded and quenched after solution thermal treatment in 950 C for 1hr. Fig 3. XRD Pattern of 15hr alloyed sample after solution treatment. Points indicated with 1, relate to β 1 (18R) phase and points indicated with 2, relate to α solid solution phase. β 1 is the desired phase in the structure for shape memory effect. Absence of γ2 (Cu 9 Al 4 ) peaks indicates the low amount of this intermetallic phase which is harmful for the mechanical properties of the alloy. Evolution of particle size distribution is shown in Figure 4 By increasing milling time to 15hr, deformation of uniaxial particles to layer-shaped particles, causes an increase in particle size. After milling for 20hr, fracture of particles caused by work-hardening, increased the amount of fine particles.

4 Fig 4. Particle size distribution of MAed powders in different milling time. It can be concluded that before 15hr, the dominating mechanism is in fracture and cold welding, and after milling for 20hr, fracture of work-hardened particles happens which is a sign of extensive milling. SEM micrographs which are shown in Figure 5 confirm the mentioned interpretations. On the basis of XRD and SEM results, several steps in mechanical alloying process can be considered: In the beginning of process, joining of particles by cold-welding, occurred simultaneous by with fracture of particles caused by strain hardening; The amount of fine and coarse particles increased at the same time. Meanwhile, particle joining takes the lead, and layer composites appear in the form of coarse particles on surface of the balls. By increasing milling time, fracture caused an increase in fine semi-axial composite particles, and diffusion of elements in the matrix increases by localized elevation in temperature. Also, low driving force channels for diffusion of atoms will be provided by strain hardening. Appearance of clear active surface by fracture makes it more easy to welding between particles. This process continues until the powder becomes homogenous. Figure 5. SEM results for samples MA-ed for different times. a. 0 hr, b. 5hr, c. 10 hr, d. 15 hr, e. 20 hr.

5 3.2 Effect of milling and solution treatment on shape memory effect: Figure 6 shows changes in strain recovery by milling time, and Figure 7 shows changes in strain recovery by solution treatment. More β 1 martensite could be produced by increasing milling time and solution treatment because of more Al and Ni diffusion in Cu matrix, leading to have more desired composition, increasing efficiency of milling process and desired shape memory effect. Decrease in 20hr milled sample could be because of more strain hardening and porosity. The diagram of solution time-strain recovery shows that efficiency of solution treatment in first stages is more than efficiency in high solution times, maybe because of appearing martensites other than β 1, like γ 1 and γ 2 phases. Figure 6. Shape memory properties of final samples made from powders milled in different times. Figure 7. Shape memory properties of sample made from 15hr milled mixture and heat treated for different times. 4. CONCLUSIONS: MA can be successfully applied to prepare Cu, Al, Ni and pre-alloyed powders. A single phase of FCC structure with lattice parameter close to that of Cu is produced after high energy MA for 15h. By increasing milling time, the positions of Cu diffraction peaks moved towards low-angles and broadened the width of Cu diffraction peaks.

6 MA caused a notable evolution in particle size and shape. With increasing milling time, joining and fracture of particles happen simultaneously, and made a wide distribution of particle size. REFERENCES: [1]. V. Recarte, J.I. Perez-Landazabal, A. Ibarra, M.L. N o, J.S. Juan, Mater. Sci. Eng. A 378 (2004) 238. [2]. S. Miyazaki, K. Otsuka, ISIJ Inter. 29 (1989) 353. [3]. M. Igharo, and J.V. Wood: Powder Met.. 28(1985) 13 I-139. [4]. Y.D. Kim and CM. Wayman; Scripta Metall., 24(1990) [5]. R.B. P erez-s aez, V. Recarte, O.A. Ruano, M.L. N o, J.S. Juan, Adv. Eng. Mater. 2 (2000) 49. [6]. A. Ibarra, P.P. Rodriguez, V. Recarte, Mater. Sci. Eng. A 370 (2004) 492. [7]. A. Ibarra, P.P. Rodriguez, V. Recarte, J.I. Perez-Landazabal, M.L. N o, J.S. Juan, Mater. Sci. Eng. A 370 (2004) 492. [8] J. San Juan, R.B. P erez-s aez,v. Recarte, M.L.N o, G. Caruana, M. Lieblich, O. Ruano, J. Phys. IV JP 5 (1995) 919. [9] T. Kaneyoshi, T. Takahashi, Y. Hayashi, M. Motoyama, J. Jpn. Inst. Met. 56 (1992) 517.