Modeling hydrogen transport causing environmental embrittlement in Co3Ti alloys

InrermetaNics 3 (1995) 115-123 0 1995 Elsevier Science Limited Printed in Great Britain. All rights reserved 0966-9195/95/%09.50 ELSEVIER

Modeling hydrogen transport causing environmental embrittlement in CosTi alloys A. Kimura,* H. Izumi, H. Saitoh & T. Misawa Muroran Institute of Technology, Muroran 050, Japan

(Received 16 February

1994; revised version received 25 May 1994; accepted 27 May 1994)

Effects of deformation rate on the fracture energy of Co21.5 at.% Ti, C-23.1 at.% Ti and Co21 at.% Ti-3 at.% Fe were investigated at temperatures ranging from 77 to 573 K. Environmental embrittlement was observed at temperatures between 200 and 420 K in Co-21.5 at.% Ti and Co23.1 at.% Ti. The lower critical temperature decreases with decreasing deformation rate. The deformation rate dependence of fracture mode is accounted for on the basis of dislocation transport of hydrogen, which results in the estimation of the activation energies of hydrogen bulk diffusion to be 0.39 eV for C-21.5 at.% Ti and 0.37 eV for Co23.1 at. % Ti. Apparent activation energy of hydrogen diffusion was also estimated by measuring the depth of intergranular cracking after deformation at 77 K following hydrogen charging under various conditions of the charging temperature and period. The resulting activation energy is about 0.4 eV for Co-21.5 at.% Ti, which is similar to that obtained by modeling dislocation transport of hydrogen, Key words: CosTi, environmental tions, hydrogen diffusion.

embrittlement,

INTRODUCTION

many other metals. Unfortunately, measurements of hydrogen diffusivity in intermetallic compounds are very limited and nothing has been studied on grain boundary short circuit diffusion. Hydrogen transport by dislocations has been studied for many metals such as iron,839 steel” and nickels,97”P’4 and shown to be accompanied by serrated flow in a certain region of temperature and strain rate, depending on the material. The region where serrations appear is generally bounded by lower and upper critical temperatures at each strain rate. The activation enthalpy associated with the lower critical temperature is considered to involve that for hydrogen diffusion.9-‘4 Takasugi and Izumi reported in their study of hydrogen effects on the tensile deformation behavior of Co3Ti alloy3 that ductility was significantly influenced by the strain rate and test temperature, and the ductility loss is accompanied by the change in the fracture mode from ductile fracture to intergranular cracking. Kimura et ~1.‘~ also showed that the temperature region where the ductility loss was observed was widened with decreasing deformation rate. They found that the environmental embrittlement of CosTi alloys occurred only after the macroscopic yielding and corre-

Environmental embrittlement of intermetallic compounds has been considered to be due to the penetration of hydrogen. 1-S One of the reasons which supports this is as follows: ductility significantly depends on strain rate and test temperature, showing more embrittlement at slower strain rate. Recently, authors have clearly shown that there is a critical hydrogen partial pressure to cause brittle fracture of CosTi alloys at ambient temperature6 and that the brittle fracture in air is due to hydrogen atoms which are produced via the decomposition of water vapor.7 Although many studies have been performed on the environmental embrittlement of intermetallics, little has been made clear on the mechanism of hydrogen transport. This is partly due to the lack of a database of hydrogen behaviors, such as hydrogen solubility and diffusivity, in intermetallic compounds. Hydrogen in intermetallics can be transported by means of lattice diffusion, grain boundary diffusion and/or transportation by dislocations similar to *Present address: Institute for Materials University, Sendai, Sendai 980-77, Japan.

hydrogen transport by disloca-

Research, Tohoku 115

A. Kimura, H. Izumi, H. Saitoh, T. Misawa

116

sponding acoustic emission.‘6~17 This suggests that hydrogen transportation involves the massive dislocation motion. In this work, possible mechanisms of hydrogen transport in CosTi have been inspected and finally the modeling of hydrogen transport by dislocations are represented to interpret the strain rate dependence of the fracture behavior of the alloys.

EXPERIMENTAL The materials used were C-21.5 at.% Ti, Co-23.1 at.% Ti and Co-2 1 at. % Ti-3 at. % Fe, which were prepared by the arc-melting method in an argon atmosphere from cobalt (99.9 wt.%), sponge titanium (99.7 wt.%) and iron (99.99 wt.%). The alloy buttons were homogenized in a vacuum of about 1.3 x lop3 Pa at 1323 K for 2 days, followed by rollings into O-7 mm thick sheets at 773 K with intermediate annealings. After machining into 10 mm square sheets, the specimens were recrystallized at 1173 K for 5 h in the vacuum. The average grain size of these alloys was the same at 7 pm. Finally, specimens were sanded into 0.5 mm thickness. The hydrogen content of these alloys was less than 0.5 ppm (w/w). In order to investigate the dependence of ductility on the test temperature and deformation rate precisely, small punch (SP) tests were carried out at temperatures ranging from 77 to 573 K at cross head speeds of 0.02, 0.2 and 2.0 mm/min. The SP test, which was a kind of bulge test, was carried out with the specimen holder shown in Fig. 1. The specimen, grasped by top and bottom die, was deformed in bulge deformation mode by loading the specimen with the steel ball and puncher. The tests were performed in air which was heated by a surrounding heater or cooled by liquid nitrogen. The test temperature was controlled within f 1°C. Ductility was evaluated from the SP fracture energy, which was

obtained by calculating the area beneath the loaddisplacement curve. To estimate activation energies of hydrogen diffusion in Co3Ti alloys, the depth of intergranular cracking after hydrogen charging, which corresponds to the hydrogen penetration distance, was measured following the different charging conditions of temperature and period. ‘s,t9 After hydrogen was cathodically charged into specimens in a solution of 1N H2S04 with an addition of 10 mg/litre of NaAs02 at a current density of 500 A/m’, the specimens were deformed at 77 K and their fracture surfaces were observed by scanning electron microscopy @EM). X-ray measurements using CL& radiation were carried out at room temperature after hydrogen charging to attempt to detect hydride formation.

RESULTS Deformation rate dependence Dependences of SP fracture energy on the test temp erature and deformation rate in Co-21.5 at.% Ti,

1

6 2 is

4

f s

2

1

61

6 -

Clamping Screws

\

4 Crosshead Speed 2.4mm Steel Ball pecimen (10x10x0.5mm)

Thermocouple

0

loo

200

300

400

500

600

Temperature ( K)

Fig 1. Schematic

diagram

of the specimen testing.

holder for SP

Fig. 2. The dependence of SP-fracture energy on the test temperature and the deformation rate: (a) O-21.5 at.% Ti; (b) C-23.1 at.% Ti; (c) Cc+21 at.% Ti-3 at.% Fe. (See text for the broken line in (b) and dotted lines in (a) and (b).)

Modeling hydrogen transport causing environmental embrittlement in CoJTi alloys

at.% Ti and Co-21 at.% Ti-3 at.% Fe are shown in Fig. 2(a), (b) and (c), respectively. In Co 21.5 at.% Ti and 23.1 at.% Ti, the SP fracture energies markedly depend on the temperature and deformation rate. The SP fracture energy of Co21.5 at.% Ti gradually decreases with increasing temperature from 77 to 300 K followed by an increase in the SP fracture energy above 300 K, which is one of the characteristics of the hydrogen embrittlement observed in metals such as iron, steels and nickel. High ductility at lower temperature is considered to be due to the immobility of hydrogen atoms. The SP fracture energy reduces as deformation rate decreaseqespecially at low temperatures. Since SP fracture energy is large in a vacuum of 1.3 x 1OA Pa even at 300 K,17 the reduction of SP fracture energy in air at around 300 K can be regarded as an environmental effect. The environmental embrittlement of CosTi alloys C-23.1

117

is considered to be due to hydrogen atoms which were produced via the decomposition of water vapor in the atmosphere.6,7 In Co-23.1 at.% Ti, the feature of the temperature dependence of SP fracture energy is different from that for Co-21 *5 at.% Ti, showing that the SP energy considerably reduces below 200 K. In contrast to these alloys, Co-21 at.% Ti-3 at.% Fe does not suffer from environmental embrittlement. Fracture of the alloy occurs in a ductile manner under any conditions in this experiment. Figure 3 shows the dependence of fracture mode of Co-23.1 at.% Ti on the deformation temperature. The fracture mode above 420 K, where the SP energy is rather high is completely ductile. With decreasing temperature to around 250 K, the SP energy decreases and the fracture mode changes to intergranular cracking. Further decrease in the temperature causes the recovery of ductility because the partial transition of fracture

Fig. 3. SEM observations of fractured specimens and their fracture surfaces of Co-23.1 at.% Ti after deformation at (a) 77 K, (b) 200 K, (c) 285 K and (d) 523 K at a deformation rate of 0.2 mm/min. Upper, middle and bottom ones are of overall, specimen surface and fracture surface respectively.

118

A. Kimura, H. Izumi, H. Saitoh. T. Misawa

mode from intergranular cracking to ductile fracture occurs. Below 200 K, the SP fracture energy again decreases with decreasing temperature, and the fracture mode at 77 K becomes a mixture of intergranular and cleavage. The SP energy no longer depends on deformation rate around 77 K, suggesting that the reduction of ductility at this temperature is not due to environmental effects, but might be due to the inherent low temperature brittleness of the compound. The low temperature embrittlement of Co-23.1 at.% Ti is not discussed in this work. It is considered that the depressed area below the broken line in Fig. 2(b) represents the reduction in the ductility caused by hydrogen embrittlement. As clearly shown in Fig. 2(a) and

(b), such environmental loss in ductility increases with decreasing deformation rate. Since the ductility loss is due to hydrogen-induced intergranular cracking and the brittleness increases with increasing hydrogen, the amount of hydrogen at grain boundaries appears to increase with decreasing deformation rate.

Cathodic charging

In order to estimate the activation energies of hydrogen diffusion in CosTi alloys, the depth of the intergranular cracking, namely, apparent hydrogen penetration distance, was measured for the specimens deformed at 77 K after hydrogen charging at different conditions of the charging period and temperature. An example of a fracture surface of G-21.5 at.% Ti is shown in Fig. 4. The specimen of Fig. 4 was cathodically charged with hydrogen at 323 K for 2 h. Although the Co-21.5 at.% Ti does not suffer from hydrogen embrittlement at 77 K and fracture occurs in a ductile mode when hydrogen is not charged, brittle intergranular cracking is observed at surfaces of specimens charged with hydrogen, as shown in Fig. 4(b) and (c). The brittle fracture was limited only to the surface strained in tension mode, and the boundary where the fracture mode changed was clear enough to measure the depth of the brittle fracture region. The extent to which the brittle fracture spread into the inside of specimen depended on cathodic char-

.”

A

3

2 I

P E C z

(4 0 40

”

”

..“...’

“““’

..

1

(b) 30.

E

///

0’ lo2

Fig. 4. Fracture surfaces of Co21.5

at.% Ti deformed at 77 K after hydrogen charging at 323 K for 2 h: (a) overall, (b) fracture surface (low magnification), and (c) the boundary where fracture mode changes.

-+353K -323K +-295K

I lo3

IO

IO5

Charging Time

(s)

Fig. 5. Depth of intergranular

IO6

cracking after deformation at 77 K immediately after hydrogen charging under various conditions of the charging period and temperature: (a) Co 21.5 at.% Ti; (b) Co-21 at.% Ti-3 at.% Fe.

119

Modeling hydrogen transport causing environmental embrittlement in CojTi alloys

t N

=F _ N

X

-o-iOpm:Ed-0.42eV

.

3

-o-20pm:Ed=0.48eV

1 O-IS-

lo-‘2 :

W

I

I

I

I

I

I

40

50

60

70

60

90

26

-c-

1 Ovm

Ed - 0.24 eV

-a-

20um.

Ed - 0.41 eV

2.8

3.0

2.6

(degree)

Fig. 7. Effects of hydrogen charging on the X-ray diffraction profile of Q-21.5 at.% Ti: (a) charged with hydrogen at 323 K for 2 h; (b) before charging. 3.2

3.4

3.6

lOoO/T(K“)

Fig. 6. Log x2/t-l/Tplots of (a) Co-21.5 at.% Ti and (b) C* 21 at.% Ti-3 at.% Fe.

ging temperature and period, as shown in Fig. 5, where (a) is for Co-21.5 at.% Ti and (b) is for Co21 at.% Ti-3 at.% Fe, indicating that the depth of intergranular cracking becomes large as the charging temperature and period increase. This result suggests that the Arrhenius-type equation, involving a diffusion controlled process, is appropriate to the charging temperature and time dependence of the depth of intergranular cracking, and thus the data obtained at all charging temperatures are expected to be superimposed by normalizing the time by a factor: t x exp (-AH/kT), where t is the charging time, T is the charging temperature, and AH is the activation enthalpy for the process, namely, hydrogen diffusion. Figure 6 shows the results of least square analyses of the data, which yield AH = 0.42 and O-48 eV for the depth = 10 and 20 pm, respectively, for Co-21.5 at.% Ti, and AH = 0.34 and O-41 eV for the depth = 10 and 20 pm, respectively, for Co-21 at.% Ti-3 at.% Fe. No examination was carried out for Co-23.1 at.% Ti, since it broke in a brittle manner at 77 K even without hydrogen, which

made it difficult to measure the depth of brittle fracture produced by cathodic charging. The estimated hydrogen diffusivities, with the assumption that x = 2&tt), and the activation energies for hydrogen diffusion of the alloys, are listed in Table 1. Figure 7 shows the X-ray diffraction profiles from the specimen surface of Co-21.5 at.% Ti (a) with and (b) without cathodic charging with hydrogen at 323 K for 2 h. After hydrogen charging, the (220) and (311) reflections decreased in intensity and additional reflections were observed, which could be indexed as due to a Liz structure with a 2.6% larger lattice parameter than Co-21 -5 at.% Ti and was considered to correspond to the CosTi hydride.

DISCUSSION Mechanism of hydroger. transport The estimation of hydrogen diffusivity from the results obtained here is difficult, since the concentration profile of hydrogen in the presence of CosTi hydride formed on the specimen surfaces by the cathodic charging is quite ambiguous.

Table 1. The estimated values of hydrogen diffusivity (DH-cath.), which are evaluated from the cathodic charging experiment and the activation energies of hydrogen diffusion in Co,Ti alloys Alloy C*21.5at.%Ti C*23.1at.%Ti C*2lat.%Ti-3at.%Fe

Ed-cath. (eV)

Ed-disl. (eV)

0.42 (0.48) Not available 0.34 (0.41)

0.39 0.37 Not available

D,-cath.(295

K) (m*/s) ____

1.1 x lo-l4 Not available 1.5 x lo-l4

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A. Kimura, H. Izumi, H. Saitoh, T. Misawa

However, the apparent activation energy obtained for hydrogen diffusion could be effective in discussing the mechanism of hydrogen transport in CosTi alloys. Tien et al.*’ derived the following equation to represent the critical velocity of dislocation to move with a hydrogen cloud, Vc, in their study on hydrogen transport by dislocations in metals: Vc = DH/kT

x EB130b

(1)

where DH is the diffusivity of hydrogen, that is Doexp (-Ed/kT), T is the deformation temperature, EB is the binding energy of hydrogen to the dislocation, and b is the Burgers vector. If the dislocation velocity, which is given by the Orowan relation, exceeds the Vc, dislocation transport of hydrogen no longer occurs. The critical velocity of dislocation ( Vc) can be correlated with the critical cross head speed of bulge deformation (io) and can be given by the following equation: & = apbVc

=apb(EB/30b) x (DolkT)

exp(-EdkT)

(2)

where a is a constant, and p is the dislocation density. Finally, the following equation is available: log( T&)

=

log C - (Ed/2.3 k T)

(3)

where C is a constant, and Ed is the activation energy of hydrogen bulk diffusion. The log(T*o) -l/Tplots for Co-21.5 at.% Ti and Co-23.1 at.% Ti, whose SP energies are 3 and 2, respectively, are shown in Fig. 8. When the data were unavailable for the value of 2 J, the extrapolated values are used, as shown in Fig. 2(a) and (b), by dotted lines. The slopes of the lines corresponding to the values for the activation energies for the relevant processes and they are 0.37 eV for Co-23.5 at.% Ti and 1

10.

10.2. 3

‘E g lo.:

3.

IO'.4 _ +

CC-21.5st.WTi

-+--

CoZi.lat.%TI

’

‘*

10.5_ 3.0 1000 / T (K

.‘)

Fig. 8. Log Z’_& - l/T plots for (b-21.5 at.% Ti and C* 23.1 at.% Ti whose SP energies are 3 and 2 J, respectively.

0.39 eV for Co-21.5 at.% Ti. The values obtained for the activation energy are summarized in Table 1. The activation energies obtained from the crack depth measurement (Ed-cath.) are larger than those derived from the modeling of hydrogen transport by dislocation (Ed-disl.), especially when the crosscut method was adopted for a larger depth: x = 20 pm. Under the cathodic charging condition, the hydrogen fugacity at the surface easily reaches lo* atoms*’ and the resulting hydrogen concentration near the surface is considered to be high enough to form hydride. Kimura and Birnbaum** concluded in their Xray measurement of hydrogen charged nickel that the additional reflections of each Ni peak corresponded to the Ni hydride, which had a larger fee lattice than Ni. In Co3Ti alloys, hydride, which has a similar structure with a 2.6% larger lattice factor than the matrix, appears to be formed at the surface by the cathodic charging. Although the thickness of the hydride layer was not measured, it would be less than several micrometers,‘8’19’22 which is smaller than the crack penetration depth (more than 20 pm). Grain boundaries with an amount of hydrogen which is enough to cause grain boundary embrittlement break in a brittle manner. The activation energies obtained from the measurements of crack depth after cathodic charging could involve the activation energies of hydrogen diffusion in CosTi hydride. It has been reported that the activation energy of hydrogen in Ni hydride is larger than in Ni by a factor of 1.3 (Ed = 0.56 eV in Ni hydride and Ed = 0.42 eV in Ni).** If this is the case also for CosTi and its hydride, the activation energy of hydrogen diffusion in CosTi alloys could be smaller than the obtained values. The smaller values of the activation energy for the small penetration depth (x = 10 pm) could be interpreted in terms of thinner hydride at the surface, -vhich does not contribute much to the hydrogen diffusion. Thus the activation energy of hydrogen diffusion in Co-21 *5 at.% Ti and Co-21 at.% Ti-3 at.% Fe are estimated to be O-4 eV or less. Since no other data have been available for the activation energy for hydrogen diffusion in Co3Ti alloys, it is difficult to determine whether the obtained value is of bulk or grain boundary diffusion. An estimation of hydrogen penetration distance from the surface (x), using the assumption x = 2flDt) with the estimated activation energy, results in x = lo-* m at 220 K for 1 h, which is too small to cause entire brittle fracture at this temperature during the testing period, namely, at most 1 h. The hydrogen embrittlement,

Modeling hydrogen transport causing environmental embrittlement in Co3Ti alloys

however, actually happens at 220 K at a deformation rate of 0.02 mm/min in Co-21.5 at.% Ti. Therefore, whatever the obtained values are of the bulk or grain boundary diffusion, both could not be supported as the hydrogen transport mechanism. Nishimura et a1.23 reported in their hydrogen permeation study that the activation energy of hydrogen diffusion in (Co, Fe)3 V alloy is 0.57 eV, which may not be much different from that in Co3Ti alloys. The relatively small values obtained in this work may be attributed to the enhanced grain boundary diffusion.‘8,‘9 However, the values obtained for the activation energies of hydrogen bulk diffusion with eqn (3), which is derived from modeling dislocation transport, are quite consistent with those estimated from the cathodic charging experiments. This suggests that the obtained energies are of hydrogen bulk diffusion and that the dislocation transport mechanism works for the hydrogen transport in Co3Ti alloys. Modeling hydrogen-induced intergranular cracking It was pointed out that there was a critical hydrogen concentration to cause hydrogen-induced cracking in materials.24 The critical amount of hydrogen at the crack nucleation site (Cu*) which depends on the material, must be attained to bring about brittle fracture by hydrogen transportation during deformation. The amount of hydrogen at grain boundaries (CoGB) in Co3Ti alloys is expected to be less than Cu* before deformation, since CosTi alloys never show brittle fracture when they are deformed in a vacuum and/or at high deformation rates, where no hydrogen is allowed to be transported to grain boundaries from the environment and even in the bulk of the specimen. During deformation at slow strain rates, however, some amounts of hydrogen (CtGB) are transported to grain boundaries and the resulting hydrogen concentration at the grain boundaries (CoB) exceeds Cu*: CoGB < c,* < CGB = CoGB+ CtGB

(4)

then the intergranular cracking happens. Sources of hydrogen to be transported are in the specimen and in air: CtGB = Cou + CUEGB

(5)

where CtnGB is the amount of hydrogen transported from the bulk of the specimen, and &GB is that from the environment.

121

As has been clearly shown in our previous studies,7 the source of hydrogen which causes hydrogen embrittlement of Co3Ti alloy is considered to be the decomposed hydrogen from water vapor in the air, namely, CtnGB << ~~~~~ holds. The ductility, therefore, significantly depends on the environment. In the case of when CtnGB > Ct~GB, which can be attained in materials containing some amounts of hydrogen, their ductility will depend on both the hydrogen concentration of the materials and the environment. If the CoGBis larger than Cu* before deformation, brittle fracture will occur regardless of deformation conditions. Recent studies on the impact property of as-cast TiAl intermetallics25,26 have shown that this material breaks in a cleavage fracture mode in impact tests at ambient temperature, showing an extremely small value of absorption energy. Although this suggests that as-cast TiAl is brittle even without the effect of the environment, there is still a possibility that Cc,is larger than CH* in as-cast TiAl, and the brittleness is due to a small amount of hydrogen included in the material. It is also considered that cathodic charging before deformation raised CoGBup to a value above Cu* which resulted in the brittle intergranular cracking without any hydrogen transport during deformation. Processes of hydrogen transport from air to the grain boundaries of Co3Ti alloys are as follows: (1) adsorption of water vapor, (2) decomposition of water, (3) picking up of hydrogen by dislocations, (4) dragging by the dislocations, and (5) transferring to grain boundaries. Tien et al.*’ calculated the time for monolayer adsorption from ideal gas formulism to be of the order of 10m6sfor hydrogen at room temperature. Although it is uncertain that water vapor behaves as hydrogen gas, the time of adsorption for water vapor may not differ by much. As for (3) if dislocations generated near the surface pick up hydrogen, this process occurs by a short range diffusion of hydrogen and is considered to be not the rate determining process. This is the case also for (5) where hydrogen is transferred within the interaction distance between dislocation and its hydrogen cloud, which is expected to be about 30b (b is the Burgers vector). Although it may be possible to interpret the deformation rate dependence of fracture behavior in terms of the decomposition rate model, the behavior of decomposition of water vapor is rather uncertain. At present, the dislocation transport model, namely process (4), appears to be the most probable rate determining process of hydrogen-induced intergranular cracking in CosTi alloys.

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A. Kimura, H. Izumi, H. Saitoh, T. Misawa

Effects of composition A marked suppression of environmental embrittlement of Co3Ti alloys is attained by the addition of 3 at.% iron. One may think that the smaller activation energy makes hydrogen easier to be transported at lower temperatures and expands the temperature region of brittle fracture. However, the activation energy of iron-added alloy, which does not suffer from environmental embrittlement is smaller than that for the alloy without the addition of iron. This suggests that the mechanism of the suppression of environmental embrittlement does not involve hydrogen transport. It is expected that the addition of iron may suppress the decomposition of water vapor at the specimen surface.7 The difference in the susceptibility to environmental embrittlement between Q-23.1 at.% Ti and Co-21.5 at.% Ti is also considered to be not due to the difference in the amount of hydrogen transported during deformation, since the relative magnitude of the values obtained for the activation energies is again not in agreement with the extent of the embrittlement of those alloys. Although structural defects due to nonstoichiometry of these two alloys are considered to influence hydrogen behavior, hydrogen diffusivities evaluated in this work are not considerably different from each other, implying that the difference in the susceptibility to environmental embrittlement between these two alloys is not due to differences in hydrogen behavior, such as interaction with the structural defects, but due to differences in grain boundary strength. It is noteworthy that there is a significant difference in the fracture energy between these two alloys at 77 K at which hydrogen effects are negligible. This suggests that the ductility of these two alloys appears to be inherently different, which may be influenced by nonstoichiometric composition. It is, therefore, considered that Cn* may be smaller in Co-23.1 at.% Ti than that in Co-21.5 at.% Ti.

CONCLUSIONS 1. The dependence of the fracture energy of Co3Ti alloys on the deformation rate and temperature is interpreted in terms of a model involving hydrogen transport by dislocations. The model predicts the activation energy of hydrogen diffusion in CosTi to be O-4 eV, which is similar to that estimated from the measurement of penetration distance of hydrogen.

Measurements of depth of intergranular cracking following hydrogen charging revealed that the activation energy of hydrogen diffusion in CosTi alloy is not influenced significantly by the change in chemical composition. The activation energies of hydrogen diffusion in Co-21.5 at.% Ti, Co-23.1 at.% Ti and Co-21 at.% Ti-3 at.% Ti are estimated to be 0.4 eV or less. The compositional dependence of the susceptibility to environmental embrittlement is considered to be not due to the compositional dependence of the hydrogen transportation. The difference in the susceptibility to environmental embrittlement between Co21.5. at.% Ti and Co-21 at.% Ti-3 at.% Fe may be due to the difference in the decomposition rate of water vapor between the two alloys. The corresponding difference between Co21.5 at.% Ti and Co-23.1 at.% Ti may be due to the difference in the inherent ductility between the two alloys.

ACKNOWLEDGEMENTS We would like to express our appreciation to Dr T. Takasugi, Institute for Materials Research, Tohoku University, for his preparation the specimens, and to students, Miss Y. Igarashi and Mr T. Sugimoto, for their help in doing bulge tests. This work was partly supported by Grant-in-Aid for Fundamental Scientific Research from the Ministry of Education, Science and Culture, Japan.

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