DRX LMF

7/26/2019 DRX LMF

http://slidepdf.com/reader/full/drx-lmf 1/5

Materials Science and Engineering A 417 (2006) 225–229

Cu–Al–Ni–Mn shape memory alloy processed by mechanicalalloying and powder metallurgy

Z. Li ∗, Z.Y. Pan, N. Tang, Y.B. Jiang, N. Liu, M. Fang, F. Zheng

School of Materials Science and Engineering, Central South University, Room 410,

Changsha 410083, PR China

Received in revised form 22 October 2005; accepted 22 October 2005

Abstract

The Cu–Al–Ni–Mn shape memory alloy has been fabricated by mechanical alloying and vacuum hot pressing and hot extrusion. SEM and X-raydiffraction analysis have been used to characterize the pre-alloyed powders and the hot extruded sample solution treated at 850 ◦C for 10 min and

then water-quenched. The shape memory recovery of the quenched sample is measured to be 100% as it is recovered in boiling water for 40 s

after deformed to 4.0%, and the shape memory recovery of the sample remains 100% as it is subjected to deforming and recovering for 100 times

cycling.

© 2005 Elsevier B.V. All rights reserved.

Keywords: Cu–Al–Ni–Mn shape memory alloy; Powder metallurgy; Mechanical alloying

1. Introduction

There has been a major interest on Cu-based shape mem-

ory alloys (SMAs) mainly due to their low cost and rela-tive ease of processing [1–3]. Among Cu-based shape mem-

ory alloys, Cu–Al–Ni alloy has higher thermal stability than

that of Cu–Zn–Al alloy [4–7]. Moreover, addition of Mn in

Cu–Al–Ni SMA has been reported enhancing the thermoelastic

and pseudoelastic properties [8]. So, the Cu–Al–Ni–Mn system

is selected in this study.

It is difficult to maintain the desired chemical compositions

andcontrol thegrain size of Cu-based SMAs by theconventional

casting method. In general, the composition change will shift the

transformation temperature and coarse grains will weaken the

mechanical properties of alloys.

It has been reported that mechanical alloying (MA) and pow-der metallurgy (P/M) with hot isostatic press (HIP) can be

used to fabricate Cu-based SMAs [9–12]. It is simpler to use

P/M to produce near-net shape alloy products and give bet-

ter controllability of the composition and grain sizes. However,

the shape memory effect of the SMAs fabricated by the com-

bining method of MA and the conventional P/M is declining

∗ Corresponding author. Tel.: +86 731 8830264; fax: +86 731 8876692.

E-mail address: [email protected] (Z. Li).

very fast with the cycling of deformation [9]. In this study, the

high-energy planetary ball milling is applied to convert the ele-

mental powder mixtures of Cu, Al, Ni and Mn into pre-alloyed

powders. The pre-alloyed powders are compacted by vacuumhot pressing and then hot extrusion to obtain the final alloy

sample.

The purpose of this work is to study the effect of MA on

the microstructure of the pre-alloyed mixture of powders and

the properties of the Cu–Al–Ni–Mn SMA prepared through ball

milling, plus vacuum hot pressing and hot extrusion.

2. Experimental procedures

2.1. Preparation of pre-alloyed powders

In the MA process, a QM-1F high-energy planetary ball mill

with four stainless steel vials was used. Each vial contained

hardened steel balls of different sizes (6, 10 and 20 mm in diam-

eter). The ball-to-powder weight ratio (BPR) used was15:1. The

specification of elemental powders and the initial mixture were

shown in Table 1.

The sealed vials were evacuated and then filled with argon to

avoid oxidation of the mixture. The mixture was then mechan-

ical milled for 1, 5, 15, 25, 35 and 45 h at a speed of 300 rpm,

respectively.

0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved.

doi:10.1016/j.msea.2005.10.051

7/26/2019 DRX LMF


226 Z. Li et al. / Materials Science and Engineering A 417 (2006) 225–229

Table 1

Specification of elemental powders and mixture

Cu powder Al powder Ni powder Mn powder

Size (mesh) 200 200 200 200

Purity (%) 99.0 99.9 99.9 99.0

Composition of

mixture (wt.%)

81 12 5 2

2.2. Vacuum hot pressing and hot extrusion

The vacuum hot pressing consisted of a 30 t hydraulic press

and a two-action piston die of 30 mm bore diameter. The pre-

alloyed powders mechanical milled for 45 h were compacted

at 850 ◦C in vacuum of 10−1 Pa under pressure of 30 MPa for

120 min. The green compact (a relative density over 94.8%) in

the evacuation of Cu-capsule was hot extruded at 900 ◦C. The

extrusion ratio was 90:1.

2.3. Measurement of shape memory effect

A strip specimen of 20 mm in length×2 mm in width×

0.5 mm in thickness was cut off from the hot-extruded rod. The

strip was solution-treated at 850◦C for 10minand thento water-

quenched. The phase transformation temperatures M s, M f , As

and Af were measured by the electrical resistance method to be

30, 15, 46 and 65 ◦C, respectively. The quenched strip was bent

to 90◦ (see Fig. 1) at room temperature. The bent strip was put

into boiled water for 40 s and then the angle θ 1 was measured.

The maximum deformation strain, ε, was pre-determined as

4% for D = 12 mm. The recovery, η, and the deformation strain,

ε, can be calculated as the follows:

η =

90◦ − θ ◦1

90◦

× 100% (1)

and

ε =

t

D+ t

× 100% (2)

Where, t was the specimen thickness and D was the diameter

of curvature [9].

2.4. X-ray diffraction analysis

A small amount of milled powders was removed after certainmilling time (at 5, 15, 25, 35 and 45 h) from the container in an

argon glove box and investigated using X-ray diffraction (XRD)

with Cu K radiation in Dmax-2500 diffractometer.

Fig. 1. The arrangement of shape memory effect measurement.

The structure of the hot-extruded sample solution-treated at

850 ◦C for 10 min then water-quenched was investigated using

X-ray diffraction also.

2.5. SEM observation

The morphology and the elemental mapping images of the

hot-extruded sample solution-treated at 850 ◦C for 10 min then

water-quenched was observed using scanning electron micro-

scope (SEM) Sirion 200 equipped with EDAX GENESIS 60.

3. Results and discussions

3.1. The structural evolution during MA

The structural evolution during MA of the Cu, Al, Ni and

Mn powder mixture is shown in Fig. 2. The pattern of 1 h MAed

powders is taken as the reference condition where the diffraction

peaks of initial components of Cu, Al, Ni and Mn appear. The

Fig. 2. XRD patterns of powder after different milling time.

Fig. 3. Changes of crystallite size andmicro-strain of Cuas a functionof milling

time.

7/26/2019 DRX LMF


Z. Li et al. / Materials Science and Engineering A 417 (2006) 225–229 227

Fig. 4. Changes of peak positions and lattice parameters of Cu matrix as a

function of milling time.

Fig. 5. The XRD pattern of sample hot-extruded and solution-treated at 850 ◦C

for 10 min then water-quenched 1-122M, 2-202M, 3-0018M, 2201, 4-128M 5-

208M, 6-1210M (2010M) 7− 1214M, (2014M), 8-2221, 9− 2016M, 10-2020M

(1220M), 11-4001, 12-040M, 13-320M, 14 − 1226M, 15 − 3212M (4221).

Fig. 7. Change of the shape memory recovery, η, as a function of deforming

cycling.

intensities of diffraction peaks of Mn, Al and Ni are lower than

that of Cu diffraction peaks because of their small amount in

the overall composition. The height of all the diffraction peaks

decrease with balling time while the width of peaks increase. It

is caused by the decrease of crystallite size and the increase of

micro-strain due to a high stresses evolving during milling balls

impacts. Considering that the increase of peak widths mainly

consists of the Cauchy part that caused by decrease of crystal-

lite size and Gauss part caused by increase of micro-strain, the

XRD profiles can be fitted by Pearson-VII function. Then the

average crystallite size D and micro-strain ε can be calculatedby substituting the fitted integral breadth into D = Kλ/βf

C cos θ

and ε = (1/4)βf G cot θ , where βf

C is Cauchy integral breadth

and βf G is Gauss integral breadth, the constants K is 0.9 and λ is

0.15405 nm. The changes in crystallite size and micro-strain of

Cu matrix as a function of milling time are shown in Fig. 3. The

crystallite size decreases while the micro-strain increases with

balling time.

Fig. 6. TEM images of the martensite structure of quenched sample (a) BF image and (b) corresponding diffraction pattern with zone axis of [0 10].

7/26/2019 DRX LMF


228 Z. Li et al. / Materials Science and Engineering A 417 (2006) 225–229

The intensity of the Ni diffraction peaks and most of Al and

Mn diffraction peaks become indistinguishable after MA 15 h. A

slight shift of position of Cu diffraction peaks (see Fig.4) reflects

the diffusion of Al, Ni and Mn into Cu matrix. The position of

Cu diffraction peaks moved towards low-angle indicates that the

lattice parameters of Cu matrix increase with milling time.

Plastic deformation and interdiffusion of elements control

the formation of amorphous and crystal phase [13], the mixing

of elements is accelerated by diffusion of Al, Ni and Mn along

dislocations and grain boundaries of Cu solid solution. After

35 h MA, only the diffraction peaks of a single phase with fcc

structure appear (see Fig.2). The X-raydiffraction results, which

agree well with the observation of Kaneyoshi et al. [14], lead to

the conclusion that a single phase solid solution is formed after

35 h MA of the elemental power mixtures. Increasing themilling

time to 45 h, theXRD pattern where thespectrumis consisting of

Cu diffraction peaks superimposed on top of broad background

presents. This broad background is associated with amorphous

phase [15].

3.2. The structure of quenched sample

The XRD pattern of the hot-extruded sample solution-treated

at 850 ◦C for 10 min then water-quenched is given in Fig. 5.

The characteristic diffraction peaks of martensite structure, as

suggested by Saburi and Wanyman [16] and Kubo and Shimizu

Fig. 8. SEM microstructures and elemental mapping images taken from the (b) of hot-extruded sample and solution-treated at 850◦C for 10min then

water-quenched.

7/26/2019 DRX LMF


Z. Li et al. / Materials Science and Engineering A 417 (2006) 225–229 229

[17] for 18R, canbe found existing between two theta of 35–85◦.

The 1 phase can be found existing also.

The microstructure and electron diffraction pattern of the

martensite above are showed in Fig. 6. The information of

the martensite structure can be obtained from the photos. (1)

Sub-structure of the martensite is the basal-stacking fault (see

Fig. 6a). (2) Thedistance between the two strong reflection spots

(the reciprocal unit interlayer spacing (RUIS)) has been divided

into three parts by two weak reflection spots and can be charac-

terized with a structure of long period stacking order (LPSO).

This characteristic structure is similar to that of M18R marten-

site found in Cu30–Au25–Zn45 [17].

3.3. Shape memory effect (SME) of quenched sample

The shape memory effect of the strip cut off from the hot-

extruded sample solution-treated at 850 ◦C for 10min then

water-quenched is measured with the results showed in Fig. 7.

The sample is undergone 4.0% deformation and then recovered

in boiling water for 40 s. The shape memory recovery, η, ismea-sured to be 100%. The η of the sample has remained 100%

as it is subjecting to deformation and recovering for 100 times

cycling. This is much higher than that of the MAed Cu–Al–Ni

shape memory alloy [9]. The reported initial recovery of the

Cu–Al–Ni shape memory alloy was 68% [9] and dropped to

30% in the second test where the deformation strain used was

only 1%. The internal cracks and the stressconcentration around

pores causeda rapid degradationof theSME in MAed Cu–Al–Ni

shape memory alloys [9].

The scanning electron microscopy and elemental mapping

images of the hot extruded sample solution-treated at 850◦C for

10 min then water-quenched are presented in Fig. 8. It is worthto point out that the metallurgical bonding among the matrix

particles has replaced the mechanical bonding or engagement

(see Fig. 8a and b). Furthermore, the internal cracks and pores

are not found in the photos. We can draw a conclusion that the

employment of vacuum hot pressing and hot extrusion here does

help eliminating the existence of pores and promoting metallur-

gical bonding among the particles of Cu matrix. The average

grain size of solution-treated sample is about 3 m, where a

small part of grains have merged together and grown up. There

is self-accommodation configuration of martensite inside some

of those merged grains (see Fig. 8a and b).

The elemental mapping images taken from the Fig. 8b

present in Fig. 8c–e (the size of elemental mapping image isthe same as that of Fig. 8b) allow to conclude that the dis-

tribution of the chemical composition is homogeneous. The

average composition determined by energy dispersive spec-

troscopy is 81.6 wt.% Cu, 11.8 wt.% Al, and 4.9 wt.% Ni and

1.7 wt.% Mn.

4. Conclusions

(1) MA can be successfully applied to prepare Cu, Al, Ni and

Mn pre-alloyed powders. A single phase of fcc structure

with lattice parameter close to that of Cu is produced after

MA for 35 h.

(2) Metallurgical bonding among the particles in Cu matrix has

been found to replace the mechanical bonding or engage-

ment after hot extruded at 900 ◦C with the extrusion ratio of

90:1.

(3) Crack- and pore-free and homogeneous samples can be pro-

duced through vacuum hot pressing and hot-extrusion.

(4) The shape memory recovery of the hot-extruded sample

solution-treated at 850 ◦C for 10 min then water-quenched

is measured to be 100% as it is recovered in boiling water

bath for 40 s after deformed to 4.0%, and the shape memory

recovery of the sample remains 100% as it was subjected to

deforming and recovering for 100 times cycling.

Acknowledgements

This study was supported by the Postdoctoral Science Foun-

dation of China (2004036427) and Ph.D. Programs Foundation

of ministry of education of China (20040533069).

References

[1] R. Kainuma, S. Takahashi, K. Ishida, Metall. Mater. Trans. A 27A (1996)

2187.

[2] Z. Li, M.P. Wang, Trans. Nfsoc. 12 (1) (2002) 6.

[3] G. Zak, A.C. Kneissl, G. Zatulskij, Scripta Mater. 34 (3) (1996) 363.

[4] V. Recarte, I. Hurtado, J. Herreros, M.L. No, J. San Juan, Scripta Mater.

34 (1996) 255.[5] V. Recarte, R.B. Perez-Saez, M.L. No, J. San Juan, J. Phys. IV 7 (1997)

329.

[6] T. Daricek, J. Lasek, N. Zarubova, J. Phys. IV 11 (PR8) (2001) 179.

[7] V. Recarte, J.I. Perez-Landazabal, A. Ibarra, Mater. Sci. Eng. A 378

(2004) 238.

[8] M.A. Morris, T. Lipe, Acta Metall. 42 (1994) 1583.

[9] S.M. Tang, C.Y. Chhung, W.G. Liu, J. Mater. Process. Technol. 63

(1997) 307.

[10] A. Ibarra, P.P. Rodriguez, V. Recarte, Mater. Sci. Eng. A 370 (2004)

492.

[11] R.B. Perez-Saez, V. Recarte, O.A. Ruano, M.L. No, J. San Juan, Adv.

Eng. Mater. 2 (2000) 49.

[12] J. San Juan, R.B. Perez-Saez, V. Recarte, M.L. No, G. Caruana, M.

Lieblich, O. Ruano, J. Phys. IV 5 (1995) 919.

[13] R.B. Schwarz, R.R. Petrich, C.K. Saw, J. Non-Cryst. Solids 76 (1985)281.

[14] T. Kaneyoshi, T. Takahashi, Y. Hayashi, J. Jpn. Inst. Met. 56 (1992)

517.

[15] H. Huang, P.G. McCormic, J. Alloys Compd. 256 (1997) 258.

[16] T. Saburi, C.M. Wanyman, Acta Metall. 27 (1979) 979.

[17] H. Kubo, K. Shimizu, Martensite Trans. JIM 17 (1976) 330.

Documents

DRX LMF