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Materials Science and Engineer

Mechanical properties of Ti–Nb biomedical shape

memory alloys containing Ge or Ga

Tomonari Inamuraa,*, Yusuke Fukuib, Hideki Hosodaa, Kenji Wakashimaa, Shuichi Miyazakic

aPrecision and Intelligence Laboratory, Tokyo Institute of Technology, 4259 Nagatsuta, Midori, Yokohama 226-8503, JapanbMatsushita Electrical Industrial Co., Ltd., 3-1-1 Yakumo, Moriguchi, Osaka, 570-8501, Japan

cInstitute of Materials Science, University of Tsukuba, Tennodai 1-1-1, Tsukuba, Ibaraki 305-8573, Japan

Received 26 November 2004; accepted 10 January 2005

Available online 2 June 2005

Abstract

Mechanical properties of Ti–24 mol% Nb-based shape memory alloys (SMA) containing 3 mol% Ga or Ge were characterized in this

paper as a part of our systematic work for the development of h-Ti based biomedical shape memory alloys. The alloys, called TiNbGa and

TiNbGe, were produced by severe cold-rolling followed by a solution treatment at 1273 K for 1.8 ks. The apparent phase was h (bcc) at RT in

both the alloys. It was revealed by X-ray diffraction pole figure analysis that a {112}hb110�h recrystallization texture was well developed in

TiNbGa. However, a {001}hb110�h deformation texture still remained in TiNbGe even after the solution treatment. Martensite transformation

temperatures were significantly lowered by the addition of Ge, compared to Ga and Al additions. TEM-EDX observation revealed that (Ti,

Nb)5Ge3 particles are formed in TiNbGe regardless of the solution treatment. The (Ti, Nb)5Ge3 particles were judged to be an ineffective

strengthener, because significant hardening was not recognized in the flow-stress of TiNbGe. TiNbGa exhibited a large shape recovery of

about 2% above RT in the strain– temperature curves during thermal cycles under external stress. The TiNbGe alloy exhibited superelasticity

of 3.5% at RT.

D 2005 Elsevier B.V. All rights reserved.

Keywords: h-Titanium alloy; Ti–Nb alloy; Superelasticity; Biomedical; Texture; Shape memory alloy

1. Introduction

Ti–Ni shape memory alloys (SMAs) are widely used in

biomedical applications such as artificial dental root,

orthodontic wires, catheters and stents due to their superior

properties in the shape memory effect and the superelasticity

[1]. However, since Ti–Ni alloys contain Ni of about 50

mol%, the possibility of Ni-hypersensitivity has been

pointed out. To the authors’ knowledge, no report exists

in the literature about the onset of Ni-hypersensitivity

attributed by Ti–Ni. However, it has yet to be determined

that Ni-hypersensitivity due to Ti–Ni does not occur.

Therefore, new practical shape memory alloys consisted of

non-toxic elements only, especially Ni-free, have been

0928-4931/$ - see front matter D 2005 Elsevier B.V. All rights reserved.

doi:10.1016/j.msec.2005.01.025

* Corresponding author.

E-mail addresses: inamura@pi.titech.ac.jp (T. Inamura),

fukui.yusuke@jp.panasonic.com (Y. Fukui).

required in the biomedical area. It is known that CP-Ti

and titanium alloys such as Ti–6Al–7Nb posses high

biocompatibility and biocorrosion resistance as well as

moderate mechanical properties [2–4]. Some h-titaniumalloys have thermoelastic martensitic transformation from h(bcc) to aW (c-centered orthorhombic [5]) and are candidate

materials to replace Ti–Ni alloys in biomedical applications.

Therefore, our group has systematically investigated the

shape memory effect and superelasticity of Ti–Nb based

and Ti–Mo based alloys containing various ternary ele-

ments belonging to 13- and 14-groups in the periodic table

[6–12]. These alloys exhibit high corrosion resistance and

biocompatibility, at least comparable to Ti–Ni, and are

applicable for biomedical use [13,14]. One of the character-

istic behavior of Ti–Nb based alloys is that severe cold-

rolling of more than 90% is possible. In our recent study,

Ti–24 mol% Nb–3 mol% Al alloy was found to exhibit

superelastic strain more than 4.5% by a formation of a well-

ing C 25 (2005) 426 – 432

T. Inamura et al. / Materials Science and Engineering C 25 (2005) 426–432 427

developed recrystallization texture [10,11]. In this study,

microstructures and shape memory properties of severely

cold-rolled TiNbX (X=Ga, Ge) alloys were investigated.

200nm200β

-121β

-1

ωω

2

121β-

(a) (b)

(c)

Fig. 1. (a) Bright field image, (b) selected area diffraction pattern and (c)

key-diagram of (b) taken in TiNbGa-ST. Electron beam direction was

b012̄�h//b11–26�N.

2. Experimental procedures

The chemical compositions of the alloys were Ti–24

mol% Nb–3 mol% Ga and Ti–24 mol% Nb–3 mol% Ge.

The alloys are termed TiNbGa and TiNbGe, hereafter.

Ingots of 30 g in weight were fabricated by an arc melting

method in Ar–1% H2 atmosphere. The weight change of

the ingots before and after the arc-melting was negligible

and no chemical analysis was conducted. The ingots were

sealed into evacuated (¨10�3Pa) quartz tubes and homo-

genized at 1273 K for 7.2 ks followed by quenching into

water. Cold-rolling of 99% reduction in thickness was

carried out for each ingot. Final thickness of the materials

after the cold-rolling was about 0.1mm. The materials after

the cold-rolling are termed as TiNbGa-CW and TiNbGe-

CW hereafter. The cold-rolled materials were sealed into

evacuated quartz tubes after making specimens for

measurements and then solution-treated at 1273 K for

1.8 ks followed by quenching into water. The solution-

treated materials are termed as TiNbGa-ST and TiNbGe-

ST, hereafter.

The specimens for the X-ray pole figure measurements

were finished by electro-polishing using a solution of 6%

perchloric acid+35% buthanol+59% methanol under 20 V

at 230 K. Since TiNbGa-ST has a relatively high trans-

formation temperatures as shown later, TiNbGa-ST was

heated up by heat-gun to eliminate aW-martensite formed

during the electro-polishing at 230 K. X-ray pole figure

measurements were conducted at RT with CuKa1 radiation

using Philips X’pert MRD system. The beam size of the X-

ray was set to be 1 mm�1 mm in the pole figure

measurements and a Ge-monochromator was used. The

range of elevation angle was from 0- to 85-. The reflectionsused were 110h, 200h and 112h.

Thin foils for transmission electron microscopy (TEM)

observations were prepared by a twin-jet polishing techni-

que at 230 K using a solution of 5% sulfuric acid+2%

hydrofluoric acid+93% methanol solution. TEM observa-

tions were conducted at 200 kV using a Philips CM200

equipped with an energy dispersive analysis of X-ray

(EDX).

Tensile properties were evaluated through loading–

unloading tensile tests with a constant strain increment

performed at ambient temperature using a Shimadzu

Autograph 500NI. Stain rate was set to be 5�10�4/s

and the constant strain increment was 1% per loading–

unloading cycle. The tensile specimens with a gauge

length of 10 mm were made by mechanical cutting and

damaged surface layer was removed by mechanical

polishing. The tensile direction of the specimens was set

to be the rolling-direction (RD).

Dynamic mechanical analysis (DMA), a kind of thermal

mechanical analysis (TMA), was also carried out to measure

the transformation behavior using a NETZSCH DMA242C.

The principle of the DMA is described elsewhere [15]. The

heating/cooling rate was set to be 5 K/min and the

temperature range of measurement was set to be 133 K to

423 K. The gage size of DMA specimens was 10 mm in

length, 1 mm in width and 0.1 mm in thickness. The loading

direction for DMA was also set to be RD. The deformation

mode was the tension under the static force (mean force) of

2, 4, 6 and 8 N. The ratio between static force and dynamic

force was kept a constant of 4:3. The frequency of dynamic

force was 2 Hz. It should be noted that it is often difficult to

measure the transformation temperature of Ti-base SMAs by

differential scanning calorimetry (DSC) due to small trans-

formation heat.

3. Results and discussion

3.1. TiNbGa

Fig. 1(a) shows TEM bright field image of TiNbGa-ST.

A corresponding selected area diffraction pattern taken from

b012̄�h//b11–26�N and its key-diagram are shown in Fig.

1(b) and (c), respectively. The apparent phase was h at RT

in TiNbGa-ST. It was confirmed that no additional phase

except for athermal N was formed in TiNbGa-ST. However,

dark-field observations using reflections of N failed to

image athermal N particles clearly. No remnants of order–

disorder transformation in h-phase were detected by TEM.

Then, the apparent phase of TiNbGa at RT is h-phase. TEMmicrostructure of TiNbGa-ST was similar to the micro-

structure observed in TiNbAl-ST in our previous study

[10,11]. However, a small amount of aW-martensite was

observed at the edge of the thin-foil and was suggested to be

stress-induced during thin-foil preparation. This residual

martensite suggests that the austenite finish temperature (Af)

is higher than RT in TiNbGa-ST.

Filled : β1Open : β2

(b) RD

TD

001011111

RD

TD

(a)

30 60 90

TiNbGa-ST011β

Fig. 2. 110h X-ray pole figure of (a) TiNbGa-ST and (b) stereographic

projection of {112}hb110�h texture. RD and TD are vertical and horizontal

to the paper, respectively. Note that the RD is b110�h, TD is b111�h and ND

is b112�h. There are two variants of h-phase (indicated as h1 and h2) in the

texture.

200 250 300 350 4000

20

40

60

80

100

Temperature, T/K

App

lied

stre

ss, σ

/MP

a MsMf As Af

dσ / dT~1.4 MPa/K

236K 262K 329K 357K

TiNbGa-ST

Fig. 4. The relationships between the applied stress and the transformation

temperatures of TiNbGa-ST obtained from Fig. 3. Linear dependence

corresponding to the Clausius–Clapeyron relationship was seen.

T. Inamura et al. / Materials Science and Engineering C 25 (2005) 426–432428

Fig. 2(a) shows 110h-pole figure obtained from

TiNbGa-ST. The pole figure was not normalized by the

diffraction intensity obtained from the same alloy com-

posed of sufficiently small grains with random crystallo-

graphic orientations. RD is vertical to the paper, and the

18MPa

35MPa

53MPa

70MPa

1%

100 150 200 250 300 350 400 450

Temperature, T /K

Str

ain,

ε

18MPa

35MPa

53MPa

70MPa

Ms

Mf

As

A fcooling

heating

Fig. 3. Strain– temperature curves of TiNbGa-ST taken by DMA. The

indicated stresses are the mean applied stresses during a cycle of dynamic

stress with the frequency of 2 Hz. The heating/cooling rate was 5 K/min.

transverse direction (TD) is horizontal to the paper as

indicated in the figures. Dark and dense parts in the pole

figures correspond to high intensity of diffracted X-ray. A

recrystallization texture was clearly observed in Fig. 2(a).

The background intensity (white region in the pole

figures) was about 10 counts/s and the maximum intensity

was more than 40000 counts/s in Fig. 2(a). Therefore, the

ratio between peak and background (P/B ratio) is about

4000 and the recrystallization texture was deduced to be

well developed. The observed recrystallization texture was

found to be {112}hb110�h type texture as schematically

shown in Fig. 2(b). As seen in the stereoprojection in Fig.

2(b), the material was mainly consisted with two h-variants (denoted by h1 as the filled markers and h2 as the

open markers) which are twin-related to each other. This

result is in good agreement with that obtained in TiNbAl-

ST in our previous study [11].

Fig. 3 shows strain– temperature relationships of

TiNbGa-ST obtained by DMA. The strain employed is the

averaged strain during dynamic loading cycles in DMA. It

should be noted that these strain–temperature curves are

similar to that obtained under constant stresses. The

indicated stresses in Fig. 3 are also the mean stress (= static

stress). It was clear in the cooling curves that a large shape

change starts at a temperature which corresponds to Ms and

that the shape change finishes at another temperature which

corresponds to Mf for each cycle. For the heating process,

the shape recovery starts at As and finishes at Af. In SMAs,

these transformation temperatures depend on applied stress

Table 1

Martensitic transformation temperatures of TiNbGa-ST

Mf Ms As Af

TiNbGa 236 K 262 K 329 K 357 K

TiNbAl 154 K 201 K 214 K 269 K

Those of TiNbAl-ST are also listed as a comparison [10].

0

200

400

600

800

1000

0 2 8 10

TiNbGa-ST: RTS

tres

s, σ

/ M

Pa

Strain, ε (%)4 6

Fig. 5. Stress– strain curve of TiNbGa-ST at RT obtained by a loading–

unloading tensile test with a constant strain increment of 1%.

T. Inamura et al. / Materials Science and Engineering C 25 (2005) 426–432 429

and usually obey the Clausius–Clapeyron relationship. The

relationships between transformation temperatures and

applied stress are plotted in Fig. 4. By extrapolating these

Fig. 6. TEM bright field image of (a) TiNbGe-CW and (b) TiNbGe-ST.

TEM-EDX revealed that the second phase was (Ti, Nb)5Ge3. Inset in (b) is

a selected area diffraction pattern taken from b012�h.

transformation temperatures into the horizontal axis corre-

sponding to the ‘‘stress-free’’ condition, the transformation

temperatures under zero stress can be evaluated and these

transformation temperatures are listed in Table 1 in addition

to those of TiNbAl as a comparison [10]. It was found that

the effect of Ga addition on the decrease of the trans-

formation temperatures was smaller than that of Al addition.

The transformation temperature hysteresis between Ms and

As is 67 K for TiNbGa and is much lager than that of 13 K

for TiNbAl. Af of this alloy is higher than RT, therefore, this

alloy with the same composition does not have a potential to

exhibit superelasticity at RT in nature. These results indicate

that modification in composition is necessary to bring out

superelasticity at RT in TiNbGa.

The stress–strain curve obtained by loading–unloading

tensile tests with the constant strain increment for TiNbGa-

ST at RT is shown in Fig. 5. At the first cycle of

deformation, no significant pseudoelasticity was observed.

This result is consistent with the transformation temper-

atures As and Af being higher than RT listed in Table 1.

However, pseudoelastic behavior was observed during

2nd–4th deformation cycles. The nature of the observed

pseudoelasticity is suggested to be not due to superelasticity

but due to reversible movement of twin-boundaries in

TD

(a)

30 60 90

TiNbGe-CW011β

RD

TD

(b)

30 60 90

TiNbGe-ST011β

RD

Fig. 7. 110h XRD pole figure of (a) TiNbGe-CW and (b) TiNbGe-ST. RD

and TD are vertical and horizontal to the paper, respectively.

0

200

400

600

800

1000

0 4 6 8 10

TiNbGe-ST: RT (298K)S

tres

s, σ

/ M

Pa

Strain, ε (%)

(a)

0

200

400

600

800

1000

Str

ess,

σ /

MP

a

TiNbGe-ST: 253K

(b)

0

200

400

600

800

1000

0 4 6 8 10

Str

ess,

σ /

MP

a

TiNbGe-ST: 193K

(c)

2

2

0 4 6 8 10Strain, ε (%)

Strain, ε (%)

2

Fig. 8. Selected stress– strain curves of TiNbGe-ST taken at (a) RT (298 K),

(b) 253 K and (c) 193 K.

0

1

2

3

4

5

6

0 6 8 10

Sup

erel

astic

str

ain,

εS

E (

%)

TiNbGe-ST

RT

193K

253K

(a)

0

1

2

3

4

5

6

Tra

nsfo

rmat

ion

stra

in, ε

T(%

)

Maximum applied strain, εMAX

(%)

Maximum applied strain, εMAX

(%)

TiNbGe-ST

RT

253K

193K

(b)

2 4

0 6 8 102 4

Fig. 9. (a) Superelastic strain (ESE) and (b) transformation strain (ET) of

TiNbGe-ST as functions of maximum applied strain (EMAX) at various

temperatures.

T. Inamura et al. / Materials Science and Engineering C 25 (2005) 426–432430

martensites which were induced by the deformation [16].

The shape recovery strain by unloading is about 2% after the

5th cycle and the maximum flow stress is about 470 MPa.

The maximum flow stress of TiNbAl with the similar

texture is about 400 MPa under similar tensile tests

described in Ref. [10].

3.2. TiNbGe

Fig. 6(a) shows TEM bright field image of TiNbGe-CW

and many particles of the second phase were confirmed.

Fig. 6(b) shows TEM bright field image of TiNbGe-ST and

the particles of the second phase were also observed. TEM-

EDX analysis showed that averaged composition of the

particles was 60 mol% Ti–5 mol% Nb–35 mol% Ge and

then the second phase was suggested to be (Ti, Nb)5Ge3.

Similar results are reported in h-Ti alloys that (Ti, V)5Si3particles are formed in the Ti–V–Si alloys [17]. The

matrix was h-phase and contains athermal N at RT as

shown in the selected area diffraction pattern taken from

b012̄�h//b11–26�N which is inserted in Fig. 6(b). However,

the reflections from athermal N seemed to be weak,

compared to those in Fig. 1(b).

Fig. 7(a) and (c) show 110h-pole figure of TiNbGe-CW

and TiNbGe-ST, respectively. The texture of TiNbGe-CW

was mainly {001}hb110�h deformation texture and is

similar to that observed in TiNbAl [11]. Textures existed

in TiNbGe-ST were seen to be mainly a mixture of

{112}hb110�h-type and {001}hb110�h-type texture, judging

T. Inamura et al. / Materials Science and Engineering C 25 (2005) 426–432 431

from Figs. 2(a) and 7(a). The existence of coarse (Ti,

Nb)5Ge3 particles should disturb the movement of grain

boundaries during recrystallization. The development of

the recrystallization texture should be affected by the

presence of these (Ti, Nb)5Ge3 particles. By considering

the results obtained in our previous study on TiNbAl, the

types of deformation texture and recrystallization texture

seemed to be essentially independent of the ternary

elements of Ga, Ge and Al.

Fig. 8(a), (b) and (c) show the selected stress–strain

curves of TiNbGe-ST at RT(298 K), 253 K and 193 K,

respectively. It was clearly confirmed that TiNbGe-ST

exhibits superelasticity at the temperatures below RT. The

maximum flow stress at RT was about 580 MPa and is

slightly higher than that of TiNbGa (470 MPa) and TiNbAl

(400 MPa) [11]. This must be due to the presence of (Ti,

Nb)5Ge3 particles. However, since significant strengthening

was not recognized, the (Ti, Nb)5Ge3 particles are supposed

to be ineffective strengthener of the alloy. As seen in the

series of the stress–strain curves in Fig. 8, the super-

elasticity becomes clearer with decreasing test temperature.

Fig. 9(a) shows the relationships between ESE and EMAX at

various temperatures. The superelastic strain (ESE) is defined

as the total recovery strain. ET is the transformation strain

which is calculated to be ESE minus elastic strain EE It was

seen that ESE increases with decreasing test temperature. ESEat 193 K was 5% and it was comparable to that of TiNbAl at

RT. Fig. 9(b) shows the relationships between ET and EMAX.

It was clearly seen that the transformation strain ETincreased with decreasing test temperature. ET at 193 K

was 2.8% and is comparable to that of TiNbAl [11]. It is

suggested that the lattice deformation strain of TiNbGe is

close to that of TiNbAl [11]. The stress to induce martensite

(0.2% flow stress) at each test temperature is plotted in Fig.

10 as a function of test temperature. According to the

Clausius–Clapeyron relationships, Ms of TiNbGe-ST at

zero stress was determined to be 55 K. It was found that the

Temperature, T/K

0 100 200 300 400

Str

ess,

σS

IMT/M

Pa

300

250

200

150

100

50

0

dσ /dT ~1MPa/K

TiNbGe-ST

Ms

55K

Fig. 10. Relationship between the stress to induce martensitic trans-

formation (rSIMT) and temperature in TiNbGe-ST. rSIMT was evaluated to

be 0.2% flow stress in the stress– strain curve obtained at each test

temperature.

addition of Ge significantly decreased these martensitic

transformation temperatures, as compared to Al and Ga

additions. These results indicate that the transformation

temperatures of TiNbGe-ST are much lower than RT. Then,

the stress to induce martensite becomes high value due to

the Clausius–Clapeyron relationship. As a result, the stress

for inducing martensite transformation exceeds to the

critical stress for slip around RT. The increases in ESE and

ET with decreasing test temperature are due to the increase in

critical stress for slip and volume fraction in stress-induced

martensite. Modification in composition, i.e., decrease in Nb

and/or Ge contents, should bring out better superelasticity at

RT in the Ti–Nb–Ge system.

4. Conclusions

Microstructures and mechanical properties of Ti–24

mol% Nb–3 mol% Ga (TiNbGa) and Ti–24 mol% Nb–3

mol% Ge (TiNbGe) shape memory alloys were investigated

and following conclusions were obtained.

TiNbGa is parent h-phase containing athermal N at RT.

Recrystallization texture of TiNbGa was {112}hb110�h after

the severe cold-rolling followed by the solution treatment

and the texture is the similar to that observed in the TiNbAl

alloy. The effect of Ga-addition on decrease in the

transformation temperatures is smaller than that of Al-

addition. TiNbGa alloy does not exhibit superelasticity at

RT because the reverse martensitic transformation temper-

atures are higher than RT. However, clear shape recovery

was confirmed above RT in the strain–temperature curves

obtained during thermal cycling under external stress in

dynamic mechanical analysis (DMA).

TiNbGe is consisted with parent h-phase containing

athermal N and coarse (Ti, Nb)5Ge3 particles. The (Ti,

Nb)5Ge3 particles are not very effective on strengthening of

the alloy. The apparent texture of TiNbGe after solution

treatment was a mixture of {112}hb110�h-type recrystalliza-

tion texture and {001}hb110�h-type deformation texture.

This should be due to the presence of coarse (Ti, Nb)5Ge3particles which interrupts the motion of grain boundaries

during recrystallization. The effect of Ge-addition on

decrease in the transformation temperatures is larger than

those of Ga- and Al-additions. The transformation temper-

atures of TiNbGe are much lower than RT. TiNbGe

exhibited clear superelasticity at temperatures lower than

RT and the largest superelastic strain in the present test

conditions was 5% at 193 K.

Acknowledgments

This work was partially supported by Furukawa Techno

Material Co. Ltd., Osawa Scientific Studies Grants Founda-

tion, Grant-in-Aid for Fundamental Scientific Research

(Wakate B: No. 16760566) and the 21st COE program

T. Inamura et al. / Materials Science and Engineering C 25 (2005) 426–432432

from the Ministry of Education, Culture, Sports, Science

and Technology, Japan.

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