Shape memory characteristics in powder metallurgy TiNi alloys

Scripta METALLURGICA et M A T E R I A L I A

Vol.

24, pp. 2 3 3 5 - 2 3 4 0 , 1990 P r i n t e d in the U . S . A .

Pergamon Press plc All rights reserved

SHAPE MEMORY CHARACTERISTICS IN POWDER METALLURGY TiNi ALLOYS

* ** ***

H. Kato*, T. Koyari*, S. Miura*, K. Isonishi** and M. Tokizane*** Department of Engineering Science, Faculty of Engineering, Kyoto University, Kyoto b06, Japan. Faculty of Education, Ibaragi University, Mite, Ibaragi 310, Japan. Department of Mechanical Engineering, Faculty of Science and Engineering, Ritsumeikan University, Kyoto 603, Japan. (Received

September

24,

1990)

Introduction In recent years considerable attention has been focused upon "near net fabrication" of TiNi alloys which possess excellent properties with regard to both mechanical strength and shape memory characteristics. However, until now, 'TiNi shape memory alloys have. not been very successfully produced by powder metallurgical methods and the stress-strain characteristics have not been investigated in detail. The present authors will report that by using prealloyed powder made by the Plasma Rotating Electrode Process (P-REP) method [1], the TiNi shape memory alloy can indeed be produced by the hot pressing method. P-REP is a recently developed technology in the production of powders, and has been recognized as an effective process in the preparation of powders of titanium and titanium alloys, which are active to impurities, oxygen and nitrogen, etc, since this method effectively suppresses impurity contamination. In the case of production of TiNi shape memory alloy, severe restrictions are imposed on impurity contents because of large sensitivity of transformation temperatures. For example, the presense of l at.Z oxygen lowers Ms by 92.6K [2]. Another merit of this method is that the obtained powders have uniform shape and size. The applications of a standard powder metallurgy method in the production of shape memory alloys has been reported by several authors [3-6]. These studies, however, laid emphasis on the improvement of fatigue properties by the refinement of grain size by powder metallurgy in CuZnAl alloys[3], and on the method to obtain bulk alloys of a desired composition through the mixture of powders with different compositions in TiNi alloys [4][~][6]. While mechanical properties in P/M alloys have been studied to some extent, a TiNi P/M alloy having shape memory properties comparable to those of typical alloys grown from a melt has not yet been reported. The purpose of the present study is to evaluate the shape memory properties of a P/M TiNi alloy made of solidified P-REP powders by comparing its stress-strain behavior with that of a TiNi shape memory alloy grown from a melt. Experimental From a prealloyed ingot prepared in a high frequency induction furnace, TiNi powders were produced by P-REP. The obtained powders had a uniform size of ]50ttm in diameter on average under the appropriate condition in P-REP; the anode diameter was 50mm and the rotating velocity

** Formerly with Nippon Steel Welding Products & Engineering Co. Ltd., Tsukigi, Chuo-ku, Tokyo 104, Japan.

2335 0 0 3 6 - 9 7 4 8 / 9 0 $3.00 + .00 C o p y r i g h t ( c ) 1990 Pergamon P r e s s p l c

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was 9Xl0Srpm in a He atmosphere. The consolidation of these powders by hot isostatic pressing (HIP), at 1073K and 180HPa for 2h, achieved a compact, having a relative density of 99.6Z. Table 1 shows the results of the chemical analysis of the prealloyed ingot electrode in P-REP, powders and an as-HIP'ed alloy. It was found that the weight Ti, Ni and impurities were not altered during the processes, P-REP and HIP.

TABLE

I. Composition

of the P/M TiNi Alloy

Ti Prealloyed

Ingot

Ni

O

used as the fractions of

(at.Z) C

50.0

49.7

0.17

0.13

P-REP Powder

50.1

49.6

0.17

0.13

As-HIP'ed Alloy

50.1

49.7

0.14

0.14

Standard TiNi alloy, used for a reference, was prepared as follows; a prealloyed ingot was hot-rolled and cold-rolled into a plate, and then annealed at 1023K. Hereafter, we will call this alloy the 'hot-roiled alloy'. These alloys were cut into specimens of rectangular shape with a spark cutting machine and a slitting wheel. Surfaces of the specimens were mechanically and electrolytically polished. For the surface observation, boundaries between adjacent powder particles and grain boundaries inside powder particles were revealed by etching using a solution of HF: HNOs: glycerine (]:1:8). Electrical resistivity measurements were made by the conventional four point method. Tensile tests were performed with an Instron type testing machine at a strain rate of 5.6X10-4/sec.

Results and Discussion The etched surfaces of the P/M alloy are shown in figure ](a) and (b). Figure 1(a) shows that powder particles became polygon-shaped after the HIP treatment and several grains existed in a particle. The mean grain size was 63~m, which was determined by the cross-cut method. This magnitude is almost identical with the mean grain size of a TiNi alloy grown from a melt. Figure 1(b) shows a scanning electron micrograph of the triple point of the particle boundaries. As to be expected in a high density P/H alloy, pores at this triple point were too small to be observed.

FIG. I. (a) Optical showing boundaries

and (b) SEM micrographs of the etched surfaces of an as-HIP'ed P/M alloy between powder particles and grain boundaries within each particle.

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POWDER TiNi ALLOYS

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Figures 2 and 3 show the variation of electrical resistance with temperature in the hot-rolled alloy and the P/H alloy, respectively. The Hs and Hf temperatures, at which martensitic transformation started and finished during decreasing temperature, were indicated by the arrows in each figure. It is known that precipitations of the Ti or Ni rich phases from the ~-TiNi phases occur by annealing at the two phase region, TizNi and TiNi3 [TJ. However, since the composition of the TiNi alloys used in the present study was close to the stoichiometric composition for a B2 type intermetallic compound, precipitation reactions were slow and the compositional change in the ~ matrix was small even after annealing for 24h at 673K. As a results, the Ms and Hf temperatures in the P/N alloy and in the hot-rolled alloy were not significantly altered by annealing. The well defined maximum at the Ms temperature in the electrical resistance curves and the small hysteresis width, Ms-Mr, comparable to that of the hot-rolled alloy, indicate that the compositional homogeneity in the prealloyed ingot was preserved in the P-REP powders.

ol

P/M

Ti-50.Oot.%Ni

~

~ x

o.6~ (a)

hot-rolled

0°%0.0.. . . .

Mf,

o.t ;oO',~'

%~oo°O°O o °"

.9

)

200

300

400

0.8 x

g °° -_~%og ,:o,=,~o, =°"

1.0

m w

o= .

1o2...-'"

0 ooo ° •

..o,

'

~°'°'"'"

. . . .

i

,

1.4

(b)

1.3 1.2 I.I 1.0

o. " o" i

673K24h ,

,

,

i

. . . .

I

I .7 1.6 1.5

Oo

f `c,

1.4

•

,

,

i

250

.....

2° " "

• ~'°" . . . .

i

,

i

I

Temperature

L

I

,

,

350

. . . .

300

i . . . . . 400

",2

!

i.z

:o

~'~ I

~:"%.

24 h

o~ °° ~',/ o.e %o 1.2 t . , . . . .°~"i . . . . "° ° , . j i , i . , ' . , , 200 3O0 400 TemDeroture / K

1.4 1 1.3~

873K 24h

300

24h

~°

200

o

: 'I(°) bJ

~s

~"

~

~

I

400

/ K

FIG.2. Temperature dependence o f t h e electrical resistance in Ti-50.0at.ZNi hot-rolled alloy, (a) no heat treatment, (b) annealed at 673K for 24h and quenched into ice brine, and (c) annealed at 873K for 24h and quenched into ice brine. Open and closed circles indicate the electrical resistance measured in the cooling and heating runs, respectively.

FIG.3. Temperature dependence of the electrical resistance in the P/H alloy, (a) as-HIP'ed, (b) annealed at 673K for 24h and quenched into ice brine, and (c) annealed at 873K for 24h and quenched into ice brine. Open and closed circles indicate the electrical resistance measured in the cooling and heating runs, respectively.

Figures 4-7 show stress-strain curves of the hot-rolled alloy and the P/H alloy after three different heat treatments. Before deformation at each temperature, a specimen was heated to 373K and cooled to the testing temperature. Therefore the initial state of a specimen before deformation corresponds to the electrical resistance curve of the cooling run in figure 2 or 3. In these curves, the residual strain after deformations was completely recovered by heating the specimen to 373K, namely both of these alloys exhibited a shape memory effect (S.H.E.). However superelasticity or an R-phase transition was not observed; this result is in accordance with the study on cold worked TiNi alloys by Hiyazaki,et.al.[7], since the TiNi alloys used in the present study had nearly equiatomic compositions and were heat treated above I023K.

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P/M As HIP'ed alloy T i - 5 0 . O a t . % Ni hot-rolled

253 K

212K

O

277 K

2°° f / 7

~ lOOI//

~ I00 0

0 .... 0 1.0 . . 1.0

0

. . . . 0 1.0 0 Strain (%)

1.0

0 1.0 0 Strain (%)

1.0

0 1.0 0 Strain (%)

1.0

0 1.0 0 Strain .1%1

1.0

2001290K

lO0~,y 552K

0''' ' 0 1.0

.,. :oO:i m

OV " /. , 0 1.0

, 0

,

1.0

0

IOOIA

1.0

0

w

Strain (%)

annealed at 67:5 K

'~ ~.200 r~ /209K ~

,

,

FIG.5. Stress-strain curves in the as-HIP'ed P/H alloy.

FIG.4. Stress-strain curves in the Ti-50.0at.ZNi hot-rolled alloy annealed at 873K for 24h followed by quenching into ~ce brine.

P/M

i

0 I°0

~gK

P/M

a n n e a l e d • at 8 7 3 K

~ 200

K

"

I00

== i o o

0 ~-200 ~r

1.0

0 1.0 Strain (%)

0

1.0 2.0

m

O 0 ' I'.0"

0 1.0' Strain (%)

0

1.0'

a

318K

"~

200 t 305K

=

'°°~7

= =oo

00

i,,,/7.

1.0 2.0 0 1.0 Strain (%)

0

1.0 ~'.0

FIG.6. Stress-strain curves in the P/H alloy annealed at 673K for 24h followed • by quenching into ice brine.

O0

1.0'

0 1.0' 0 ' I~0' Strain I%1

FIG.7. Stress-strain c u r v e s i n t h e P/N a l l o y a n n e a l e d a t 873K f o r 24h f o l l o w e d by q u e n c h i n g i n t o i c e b r i n e .

12

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The temperature dependence of yield stress determined from the stress-strain curves of the hot-rolled alloy and the P/N alloy is shown in figures 8 and 9, respectively. With decreasing temperature, the yield stress tends to decrease above the Ms temperature, and tends to increase below Ms, showing a well-defined minimum in yield stress at the Ms temperature. This indicates that lattice softening occurs near Ms, and the deformation above the Ms proceeds by the formation of stress induced martensite (S.I.H.).

Ti-5Oot%Ni

Ti-50.Oat.%Ni hot-rolled alloy 400

o as H I P ' e d

o annealed at 673K for 2411 A 873K 2 4 h ~ 6 7 3 K 2411 300

PIM alloy

500

/

• 6 7 3 K 24h

400

• 873K 24h

O

3OO

//

D.

zoo

200 IO0

0

" °e°~%.o o o

,~/

•-. ?'/

.

.

.

.

200

i

,

250

/

• 8 7 3 K 24h

,

.

,

m

e o

I00

. . . .

300

i

.

,

,

o

°oo

°°

o//

"''.. °o

350 250 300 350 Temperature I K

Temperature / K

FIG.8. Temperature dependence of yield stress in the Ti-50.0at.ZHi hot-rolle~ alloy.-

FIG.9. Temperature in the P/H alloy.

dependence

of yield stress

Therefore, the yield stress at deformation above the Hs temperature is the critical stress to induce martensite, a M , from the bcc matrix. It was found that the slope ' d a M / d T , is constant in the temperature range above Ms. Using this yalue, the latent heat to form S.I.H., AH, can be estimated from the Clausius-Clapeyron relationship given by, AH = T - ( d G M / d T ) - H * ~

/p .

H e r e , T i s t h e t e m p e r a t u r e , H i s t h e m o l e c u l a r w e i g h t , A~ i s t h e b r a n s f o r m a t i o n s t r a i n and p i s the density. Table 2 indicates the calculated l a t e n t heat of s t r e s s - i n d u c e d transformation at t h e Ms temperature determined from t h e electrical resistivity measurement. Corresponding to the changes in the Hs temperatures AH could also be altered by annealing at 673K and at 873K. It can be seen that the magnitudes of 8H in the P/H alloy were within the same range as those in the hot-rolled alloy.

TABLE 2.

Latent Heat of Stress Induced Transformation

Alloy

Heat Treatment

Hs (K)

dGM/dT (MPa/K)

AH at Hs (J/mol)

Hot-rolled

673K 24h 673K 24h ~ 873K 24h 873K 24h

316 313 312

7.29 6.27 5.00

II.9XI0 z I0.IX10 z 8.1X10 z

P/H A l l o y

As-HIP'ed 673K 24h 873K 24h

304 308 305

5.84 8.95 7.33

9.2XI02 14.|X10 z ll.6XlO 2

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At temperatures below Ms, the alloys had thermally transformed partially to martensite above Hf, and completely below Hf before the deformation. As shown in figures 8 and 9, the yield stress between the Hf and Ms temperatures decreased rapidly with increasing temperature in both of these alloys. This change was smooth below Hf, at which temperature range it has been known that deformation proceeds by the coalescence of variants of thermally-induced martensite. It can be seen from these figures that the stress required for this process was almost equal in the P/H and hot-rolled alloys. Thus it can be concluded that the P/M alloy and the hot-rolled alloy had the same mechanical properties associated with the martensitic transformation. However, these results are for the early stage of deformation, that is, at the initiation of the formation of S.I.H. or the variant coalescence. When the P/H alloy is deformed to a relatively larger strain, mechanical properties different from those of the hot-rolled alloy should be obtained if local plastic deformation around particle boundaries occurs, and it should degrade the recoverability of shape strain due to S.H.E. To determine the maximum recoverable strain due to S.H.E., the gauge lengths of a tensile specimen before deformation and after shape recovery by heating to 373K were compared. Stress-strain curves in the hot-rolled alloy and the P/H alloy are shown in figures ]0 and I], respectively. In the hot-rolled alloy 6.2Z strain was recovered due to S.H.E., while in the P/H alloy 5.2Z strain was completely recovered. When the latter alloy had been deformed by 8.2Z at the same temperature, the residual strain after unloading was 6.1Z. 5.8Z of this strain was recovered due to S.H.E., as indicated by the dotted line in figure II. Such a large amount of shape recovery due to S.H.E., comparable to that of the alloy grown from a melt, has never been observed in the P/M processed TiNi alloy.

As HIP'ed P/M alloy

500

40O

Ti-50.Oat.%Ni hot-rolled alloy ~I.~250[_ 200 Deformedot 316K

300

2oo

/ /

,,

,

,

I

2

I00

i

,/,/,

........................

0

J

g

Z//I

= ,,o I

vt

Deformed at 512K

3 4 5 Strain (%)

L__J

6

o 7

o

ii

12545

6789

Strain I%1

FIG.IO. S t r e s s - s t r a i n c u r v e in t h e Ti-50.Oat.gNi hot-rolled alloy annealed 873K f o r 24h and q u e n c h e d i n t o i c e b r i n e . D o t t e d l i n e s show t h e s t r a i n r e c o v e r e d by h e a t i n g t h e s p e c i m e n t o 373K.

FIG.ll. Stress-strain curve in the a s - H I P ' e d P/M a l l o y . D o t t e d l i n e s show t h e r e c o v e r e d s t r a i n by h e a t i n g t h e s p e c i m e n to 373K.

Acknowledgement We a r e g r a t e f u l t o The Furukawa E l e c t r i c and h o t - r o l l e d a l l o y .

Co. L t d . f o r t h e p r o v i s i o n

of a TiNi p r e a l l o y e d

ingot

References l. 2. 3. 4. 5. b. 7. 8.

K . I s o n i s h i , M.Kobayashi and H . T o k i z a n e , T e t s u - t o - H a g a n e . 75, 9 9 ( 1 9 8 9 ) . Y . S h u g o , S.Hanada and T.Honma, B u l l u t e i n of RITU. 4 1 , 3 5 ( 1 9 8 5 ) . J . J a n s s e n , F . W i l l e m s , B . V e r e l s t and J . M a e r t e n s and L . D e a l e y , J . d e P h y s i q u e . C4, 8 0 9 ( 1 9 8 2 ) . Y . S e k i g u c h i , K . F u n a m i , H.Hunakubo and Y . S u z u k i , i b i d , 2 7 9 ( 1 9 8 2 ) . W . A . J o h n s o n , J . A . D o m i g u e and S . R . R e i c h m a n , i b i d , 2 8 5 ( 1 9 8 2 ) . W . A . J o h n s o n , J . A . D o m i g u e , S . R . R e i c h m a n and F . E . S c z e r z e n i e , i b i d , 2 9 ] ( 1 9 8 2 ) . D.M.Poole and W.Hume-Rothery, J . I n s t . M e t a l s . 83, 4 7 3 ( 1 9 3 4 ) . S . M i y a z a k i , Y.Ohmi, K . O t s u k a and Y . S u z u k i , J . d e P h y s i q u e . C4, 2 5 5 ( 1 9 8 2 ) .