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Hydrogen property in lattice associated with the embrittlement of Co3Ti alloys
Intermetallics 5 (1991) 443448 0 1997 Elsevier Science Limited Printed in Great Britain. All rights reserved PII:
SO966-9795(97)00015-O
0966-9795/97/%17.00
+ 0.00
ELSEVIER
Hydrogen property in lattice associated with the embrittlement of Co3Ti alloys T. Tatkasugi,” A. Kimura,” T. Sugimoto,b H. Saitohb & T. Misawab UInstitute,for MateriaIs Research, Tohoku University, Katahira 2-l-1, Aoba-ku 980-77. Japan hDepartment qf Materials Science and Engineering, Muroran Institute of Technology, Mizumoto 27-1, Muroran 050, Japan
(Received 23 October 1996; accepted 27 January 1997)
Some observations relating to solubility, diffusion and trap site of hydrogen were perlformed on Co3Ti alloys in association with their embrittlement. Hydrogen diffusivity at room temperature was evaluated to be low, similar to that of pure Ni. The desorption curve of hydrogen reached the maximum at -430 K and the corresponding activation energy was estimated to be 48 kJmol-‘. The fracture energy of the hydrogen-precharged Co3Ti alloys measured by a small punch (SP) test was dependent on the hydrogen content and the alloy stoichiometry; it decreases with increasing hydrogen content and Ti content. It is suggested that the environmental embrittlement of the Co3Ti alloy requires the hydrogen transport mode not via lattice diffusion. Also, the difference in the environmental embrittlement between the two tested alloy compositions is primarily due to the difference in the critical hydrogen content causing grain boundary fracture (i.e. intrinsic grain boundary cohesion). 0 1997 Elsevier Science Limited. Key words: A. intermetallics,
boundaries,
miscellaneous, B. hydrogen embrittlement, structure, F. mechanical testing, trace element analysis.
D. grain
Concerning the lattice property of hydrogen in Co,Ti alloys, only diffusion data at room temperature have been reported by measuring a circumferential brittle zone introduced in cathodically charged specimens. 9 On the other hand, the embrittlement due to ‘residual’ hydrogen has not yet been observed, although a number of observations have been performed on the environmental embrittlement of Co3Ti alloys.1,2,9-‘4 In this study, some observations relevant to solubility, diffusion and trap site of hydrogen in Co3Ti alloys are reported. Also, to learn the effect of ‘residual’ hydrogen on ductility, a tensile test is carried out in liquid nitrogen using specimens thermally charged in hydrogen gas (i.e. with a variety of concentrations of hydrogen). Finally, a few implications are discussed concerning the environmental embrittlement of Co3Ti alloys.
INTRODUCTION The initial discovery that ordered intermetallic is very susceptible to embrittlement due to hydrogen decomposed from moisture in air at ambient temperature was made using LIZ-type Co3Ti polycrystals.‘,* The specimen deformed in air showed considerably lower tensile elongation than that deformed in a vacuum. Since this finding, so-called environmental embrittlement has been shown to take place in many environmental media containing moisture (or aqueous solution) at ambient temperature in a wide range of deformation rates and also to occur in many types of crystal structures, alloy systems and microstructures.3-8 For a better understanding of the embrittlement of ordered intermetallics, as well as ordinary materials, the hydrogen property in lattice (such as solubility, diffusion and site occupancy) is important. Also, an understanding of the embrittlement due to ‘residual’ hydrogen (i.e. hydrogen contained in lattice prior to deformation) is necessary to clarify how important a role hydrogen decomposed from moisture plays in the environmental embrittlement.
EXPERIMENTAL Co3Ti alloys with two nominal compositions of 23 mol%Ti and 21 mol%Ti were used in this study. Starting raw materials were 99.9 wt% cobalt and 443
444
T. Takasugi
99.9 wt% sponge titanium. Alloy buttons were made by arc melting in argon atmosphere on a copper hearth and then homogenizing in a vacuum at 1323 K for 2 days. These alloy buttons were repeatedly rolled at 773 K in air and annealed in vacuum at 1273 K for 5 h. The final thickness was prepared to be about 1.Omm. Next, small punch (SP) specimens, and thermal absorption and desorption specimens with a dimension of 10 x 10 mm were prepared by a precision wheel cutter. All of these specimens were then annealed in vacuum at 1273 K for 5 h to prepare a well-ordered microstructure. The average grain size of all the specimens prepared was quite similar, i.e. ranging from 6 to 8 pm. Finally, the SP specimens, and the absorption and desorption specimens, were abraded into O-5 and 0.25mm thickness by SIC paper and alumina particle, respectively. Hydrogen was thermally charged in one atmospheric H2 gas at 473 K for desired times. The charged specimens were quenched in an oil bath, the temperature of which was held at room temperature, they were rinsed in an acetone solution by an ultrasonic cleaner, then kept at room temperature. Hydrogen concentration in the charged specimens was measured by an argon carrier fusion-thermal conductivity method. Also, X-ray diffraction measurement was performed on the hydrogen-charged and uncharged specimens using Cu Ka radiation. Next, a hydrogen thermal desorption test was performed on the specimens preceded by thermal charging in H2 gas at 473 K for 4 h. The hydrogen-charged specimen was placed in a quartz chamber and then evacuated to be below 10p6Pa at liquid nitrogen temperature using a turbo-molecular pump combined with an oil rotary pump. Three different heating rates (i.e. 0.5, 1.0 and 2.0 Kmin-‘) were chosen and the desorbed hydrogen gas was monitored in a temperature range between room temperature and 1200 K by quadrapole mass spectrometer (QMS) (in the condition of an emission current of 2.0mA and a sampling time of 60 s). Evaluation of ductility (and toughness) of Co3Ti alloys was done by SP test in liquid nitrogen (i.e. at 77K). This test in liquid nitrogen was carried out in order to exclude an entry of hydrogen decomposed from moisture into the materials (i.e. to observe only the effect of the ‘residual’ hydrogen). Details of the SP test, which is a kind of bulge test, were described in the previous articles.9~13~‘4Ductility was evaluated by SP fracture energy which is obtained by calculating an area beneath a loaddisplacement curve. It has been well accepted that
et al. there is a good correlation between the magnitude of the SP fracture energy and the degree of toughness. Fracture surfaces of the deformed specimens were observed by a scanning electron microscope (SEM).
RESULTS Figure 1 shows variations of hydrogen concentration with hydrogen gas exposure time for Co-23 mol%Ti and Co-21 mol%Ti alloys. Hydrogen concentration for both alloy compositions parabolically increased with increasing exposure time. Hydrogen concentration of the Co-23 mol%Ti alloy increased more rapidly than that of the Co21 mol%Ti alloy, perhaps indicating that hydrogen diffusion in the former alloy composition is higher than that in the latter. Here, it is noted that hydrogen concentration still increases within the exposure time measured (i.e. < 3 x lo4 s). Therefore, on the basis of an assumption x = fi, the hydrogen diffusivity at 473 K in the Co3Ti alloy is estimated to be less than 5.2 x lo-i3(m2 s-i). This value may be coherent to the hydrogen diffusivity (i.e. 1.1 x 10-‘4(m2s-‘) at 295 K) reported in the Co-21.5 mol%Ti alloy by measuring the circumferential brittle zone introduced in the cathodically charged specimen.9 Figure 2 shows the X-ray diffraction profiles from the hydrogen charged (Fig. 2(a)) and uncharged (Fig. 2(b)) specimens of the Co-23mol%Ti alloy. In both specimens, all reflection lines were indexed as due to an L12 structure. No additional reflection lines were observed. Thus, Co3Ti hydride
+ -D_
f
-0
5
10
15 Time
/
Co-23mol%Ti Co-21 mol%Ti
20
25
30
ks
Fig. 1. Variations of hydrogen concentration with hydrogen gas exposure time for Co-23mol%Ti and Co-21 mol%Ti alloys. Hydrogen gas exposure was carried out in one atmospheric H2 gas at 473 K.
Hydr#ogenproperty in lattice associated with the embrittlement of Co3Ti alloys
temperatures above about 600K, hydrogen desorption remained at a very low level. An increase in the heating rate resulted in an increase in the temperature (Tc) showing the maximum in the desorption vs temperature curve. Here, an activation energy, E, for the hydrogen desorption from trap site (or soluble position) can be derived on the basis of the following equation: I5
ia) Hydrogen-charged
(E,a/RTz)
*P (b) Uncharged
I 30
I
I
I
I
I
I
40
50
60
70
60
90
28
100
I deg.
Fig. 2. X-ray diffraction profiles of (a) hydrogen-charged (b) uncharged Ce23 mol%Ti alloys. Hydrogen charging done at 473 K for 4 h.
and was
was not detected in th[e hydrogen-charged specimen. However, it is apparent that the peak position of each reflection line shifted toward a lower angle and therefore the lattice parameter of Ll;? structure in the hydrogen-charged specimen increased by a few per cent. Quite the same result was observed in the Co-21 mol%Ti allo:y. Figure 3 shows the hydrogen desorption curve measured at three different heating rates in the Co-23 mol%Ti alloy. Hydrogen desorption measured from Hz ion current rapidly increased with increasing temperature and reached a maximum at about 430K, followed by a rapid decrease. At
.:. .-,
Q:
E F 5 0
A
10-g
:
I
445
= A exp( -E,/RT’)
(1)
where a! is the heating rate, R is the gas constant, T, is a maximum temperature, and A is constant. The Arrhenius plot of eqn (1) i.e. In (a/T:) vs l/T, is shown in Fig. 4. From this figure, the activation energy of hydrogen desorption from trap site (or soluble position) in the Co3Ti alloy was estimated to be about 48 kJmol-i. If one postulates that the binding energy of hydrogen is the difference in the activation energies between the desorption and lattice diffusion processes, approximately 13 kJmol_’ can be derived as the binding energy of hydrogen with trap site in the Co3Ti alloy (here, the value of 35 kJmol19 was used as the activation energy for lattice diffusion of hydrogen). A different interpretation for E, is also possible; this value is almost identical to the activation energy for lattice diffusion of hydrogen, therefore the desorption process is rate-controlled by the lattice diffusion of hydrogen itself. It is interesting to learn how the residual hydrogen affects the mechanical property of the CosTi alloy. Figure 5 shows the variation of SP fracture energy with hydrogen concentration for the Co-23 mol%Ti and Co21 mol%Ti alloys which were hydrogen-charged for a desired time (i.e. hydrogen concentration) at 473K. It is apparent that the SP fracture energy in the Co-21 mol%Ti
1o-6 \
Heating rate
.. -.
1\1 ‘.:..
---0SWmin P.OWmin l.OWmin
:r
10‘10
10-1'
5 I"
i:
10‘12
,o-13
-,
200
’
400
600
Temperature
’
600
I
’
1000
1200
1o-8 ‘. ’ . . . ** ’ ’ ’ 2.0 2.1
2.2
IIT,
/ K
Fig. 3. The effect of heating rate on variation of H2 ion current with temperature for Co-23 mol%Ti alloy.
..* ’ ’
Fig. 4. Arrhenius alloy.
’
- ’ ’ ’ ’ . ’ ’ . ’ ’ ’ ’ ’ ’ s’ * ’ ’
2.3
2.4
; 5
/ 1O-3 K-’
plot of a/Tc2 vs l/Tc for Co_23mol%Ti
T. Takasugi
446
et al. Co-23 mol%Ti alloy was moderately dependent on the hydrogen content, that is, it decreased with an increase in the hydrogen concentration. Measured SP fracture energy was reduced from 2.5 J (at 4mass ppm H) to 0.4 J (at 75 mass ppm H). Corresponding to this change, the brittle grain boundary fracture mode, accompanied with smooth facets, became more dominant in SEM fractography (Fig. 6).
20
40
Hydrogen
Concentration
0
60 /
80
DISCUSSION
mass ppm H
Fig. 5. Variation
of SP fracture energy with hydrogen concentration for Co-23 mol%Ti and Cc~21 mol%Ti alloys. Specimens were hydrogen-charged for desired time at 473 K. Tests were done in liquid nitrogen at a cross-head speed of 0.2mmmin-‘.
alloy was consistently very high (-6 J) in the prepared hydrogen content (or slightly decreased with increasing hydrogen concentration). SEM fractography showed mainly ductile transgranular fracture mode with dimple pattern (Fig. 6). On the other hand, the SP fracture energy in the
Fig. 6. SEM fractography of (a) Co-23mol%Ti and (b) C-21 mol%Ti alloys containing 70 mass ppm H and 75 mass ppm H, respectively. Tests were done in liquid nitrogen (at 77 K).
The hydrogen diffusivity (at room temperature) in Co3Ti alloys estimated in this study and also reported in a previous study9 is similar to the value measured in pure Ni (6.9 x 10-14m2s-‘)‘6 and the value extrapolated from high temperature data measured in (Co,Fe)3V alloy with Liz structure (1 .l x 10-14m2 s-‘)17 but smaller than that measured in Fe (7.2 x 10-9m2s-‘)‘8 and N&Al alloy with L12 structure (1.2 x 10-‘0m2s-‘).‘9 Here, data for hydrogen diffusion shown in the Co3Ti alloy reveal that hydrogen causing the environmental embrittlement is not transported via lattice but via short circuit path (e.g. dislocations or grain boundaries) to grain boundaries, or otherwise, by another transport mode such as dislocation sweeping. If 1.1 x 10-14m2 s-l is chosen as the lattice diffusivity of hydrogen in the CosTi alloy, hydrogen is transported only by 3.1 x lOA mm into the alloy interior. This distance is too short to cause whole brittle fracture during testing time, i.e. mostly less than 15 min. In this study, even though a large amount of hydrogen is charged into the Co3Ti alloys, hydride was not formed. This result indicates that embrittlement by residual hydrogen as well as penetrating hydrogen is caused by atomic hydrogen. Of course, an enrichment of a high level of hydrogen at grain boundaries is required to result in embrittlement, as will be discussed. Also, it was not determined whether hydrogen is trapped to preferential sites or exists in the soluble form in the lattices. In the former case, grain boundaries may be considered as the possible trap site of this alloy. In the latter case, octahedral sites (O-site) and tetragonal sites (T-sites)20-22 are considered. In iron and steel, the value of 58.6 kJ mol-’ has been estimated as the binding energies of hydrogen with dislocation and grain boundaries.23 Further investigations are needed to determine the preferential trap site or soluble position of hydrogen in this alloy.
Hydrogen
property
in lattice associated
The ductility (i.e. the toughness) of the Co23mol%Ti alloy was dependent on the ‘residual’ hydrogen content, as shown in Fig. 5. In the test in liquid nitrogen, absorption of hydrogen from air is impossible and in addition the residual hydrogen is immobile. These facts imply that the immobile hydrogen enriched to grain boundaries prior to deformation (in stead of penetrating hydrogen) certainly affects the grain boundary cohesion and the associated plasticity, resulting in the grain boundary separation. On the other hand, the ductility of the Co--21 mol%Ti alloy was independent of the ‘residu.al’ hydrogen content and very high even in a. high hydrogen content (N 70mass ppm H). This result means that the alloy stoichiometry plays an important role in the grain boundary cohesion and the associated grain boundary embrittlement by hydrogen. It has been suggested by a number of experimental observations24*25 and theoretical calculations26,27 in L12 structure that a grain boundary consisting of stoichiometric alloy composition has a more distorted structure (like atomic cavities) than that consisting of off-stoichiometric alloy composition, resulting in reduced grain boundary cohesion. Consequently, the critical hydrogen content,r4 which is defined as the lowest hydrogen content at grain boundary causing intergranular fracture, appears to be low and therefore very sensitive to the residual hydrogen content in lattice of near stoichiometric alloy composition (the Co-23mol%Ti alloy in this study). On the other hand, the critical hydrogen content in the Co-21 mol%Ti alloy appears to be high. Thus, the results indicating ‘,2,9-‘4 that the off-stoichiometric CosTi alloy (i.e. the Co-21 21.5 mol%Ti alloys) has been more resistant to environmental embrittlement are suggested to be attributable to higher critical hydrogen content (i.e intrinsically higher grain boundary cohesion) in addition to a lower decomposition rate of water vapor on the alloy surfac~e.28Finally, it is noted that the Co3Ti alloy has an intrinsically high fracture toughness if the residual1 hydrogen as well as the penetrating hydrogen are excluded from materials.
CONCLUSION Some observations relating to solubility, diffusion and trap site of hydrogen in Co3Ti alloy were carried out. Also, the effect of ‘residual’ hydrogen on ductility was observed by a small punch (SP) test in liquid nitrogen, using hydrogen-charged specimens. The following results were obtained by the present study:
with the embrittlement
of CojTi alloys
447
1. Hydrogen diffusivity at room temperature in the CosTi alloy was evaluated to be low as in pure Ni. 2. The desorption curve of hydrogen in the Co3Ti alloy reached the maximum at N 430 K and corresponding activation energy was estimated to be 48 kJmol-i. 3. SP fracture energy in the Co-23 mol%Ti alloy was dependent on the hydrogen content and decreased with increasing hydrogen content, while that in the Co-21 mol%Ti alloy was independent of the hydrogen content and showed consistently high values. 4. It is considered that the environmental embrittlement of the CosTi alloy requires a hydrogen transport mode other than via lattice diffusion. Also, the difference in the environmental embrittlement between the two prepared alloy compositions is primarily due to the difference in the critical hydrogen content causing grain boundary fracture (i.e. intrinsic grain boundary cohesion).
ACKNOWLEDGEMENTS This work was supported in part by the New Energy and Industrial Technology Development Organization (NEDO) in Japan. The authors are grateful to Nikko Inspection Service Co., Ltd. for hydrogen analysis.
REFERENCES I. Takasugi, T. and Izumi, 0, Scripta Metall., 1985, 19, 903. 2. Takasugi, T. and Izumi, 0, Acta Metall., 1986, 34, 607. 3. Takasugi, T. and Izumi, 0, Materials Forum, 1988, 12, 8. 4. Takasugi, T., High-Temperature Ordered Intermetallic Alloys IV, in MRS Symposium Proceedings Publication, Vol. 213, 1991, p. 403. 5. Takasugi, T., Critical Issues in the Development of High Temperature Structural Materials, ed. N. Stoloff, D. J. Duquette and A. F. Giamei. The Minerals, Metals and Materials Society, Warrendale. PA. 1993, p. 399. 6. Liu, C. T., 6th Int. symp. Intermetallic compounds-Structure and mechanicalproperties, ed. 0. Izumi. JIM, 1991, p. 703. 7. Liu, C. T., Ordered intermetallics-Physical metallurgy and mechanical behavior, ed. C. T. Liu et al. Kulwer Academic Publishers, 1992, p. 321 8. George, E. P. and Liu, C. T., High-temperature ordered intermetallic alloys VI, in MRS Symposium Proceedings Publication, Vol. 364. 1995. p. 1131. 9. Kimura, A., Izumi, H., Saitoh, H. and Misawa, T., Zntermetallics, 1995, 3, 115. 10. Liu, Y, Takasugi, T., Izumi, 0. and Yamada, T., Acta Metall., 1989, 37, 507.
11. Liu, Y., Takasugi, T., Izumi, 0. and Suenaga, H., Journal of Mater. Sci, 1989, 24, 4458.
T. Takasugi
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12. Takasugi, T., Takazawa, M. and Izumi, O., Journal of Mater. Sci., 1990, 25, 4239.
13. Izumi,
H., Saitoh,
H., Kimura,
A. and Misawa,
T.,
Journal of Japan Inst. Metals, 1993, 57, 1215. 14. Kimura, A., Izumi, H., Misawa, T. and Saitoh, H., Mater. transaction, JIM, 1994,35, 879. 15. Redhead, P. A., VACUUM, 1962,12,203. 16. Robertson, W. M., Z. Metallkde, 1973, 64, 436. 17. Nishimura, C., Komaki, M. and Amano, M., Proc. of the
5th Inter. Conf. on the Effect of Hydrogen on Material Behavior, TMS, in press. 18. Kiuchi, K. and McLellan, R. B., Actu Metall., 1983,31,961. 19. Yang, L. and McLellan, R. B., Journal of Phys. Chem. Solids, 1996, 57, 233. 20. Carstanjen, H. D., Phys. Status Solidi (a), 1980, 59, 11.
et al. 21. Yang, L. and McLellan, R. B., Acta metall. mater., 1994, 42, 3993. 22. Yang, L. and McLellan, R. B., Scripta metall. mater., 1994, 32, 779. 23. Hirth, J. P., Metall. Trans. A, 1980, 11, 861. 24. Takasugi, T., Masahashi, N. and Izumi, O., Acta Metall., 1987, 35, 381. 25. Liu, C. T., White, C. L. and Horton, J. A., Acta Metall., 1985, 33, 213. 26. Takasugi, T. and Izumi, O., Acta Metall., 1983, 31, 1187. 27. Kruisman, J. J., Vitek, V. and Hosson, J. Th., Acta Metall., 1988, 36, 2729. 28. Misawa, T., Saitoh, H., Yamamura, Y., Tanabe, H. and
Takasugi, T., Proc. of the Japan International Symposium., 1993, p. 1438.
SAMPE
SO966-9795(97)00015-O
0966-9795/97/%17.00
+ 0.00
ELSEVIER
Hydrogen property in lattice associated with the embrittlement of Co3Ti alloys T. Tatkasugi,” A. Kimura,” T. Sugimoto,b H. Saitohb & T. Misawab UInstitute,for MateriaIs Research, Tohoku University, Katahira 2-l-1, Aoba-ku 980-77. Japan hDepartment qf Materials Science and Engineering, Muroran Institute of Technology, Mizumoto 27-1, Muroran 050, Japan
(Received 23 October 1996; accepted 27 January 1997)
Some observations relating to solubility, diffusion and trap site of hydrogen were perlformed on Co3Ti alloys in association with their embrittlement. Hydrogen diffusivity at room temperature was evaluated to be low, similar to that of pure Ni. The desorption curve of hydrogen reached the maximum at -430 K and the corresponding activation energy was estimated to be 48 kJmol-‘. The fracture energy of the hydrogen-precharged Co3Ti alloys measured by a small punch (SP) test was dependent on the hydrogen content and the alloy stoichiometry; it decreases with increasing hydrogen content and Ti content. It is suggested that the environmental embrittlement of the Co3Ti alloy requires the hydrogen transport mode not via lattice diffusion. Also, the difference in the environmental embrittlement between the two tested alloy compositions is primarily due to the difference in the critical hydrogen content causing grain boundary fracture (i.e. intrinsic grain boundary cohesion). 0 1997 Elsevier Science Limited. Key words: A. intermetallics,
boundaries,
miscellaneous, B. hydrogen embrittlement, structure, F. mechanical testing, trace element analysis.
D. grain
Concerning the lattice property of hydrogen in Co,Ti alloys, only diffusion data at room temperature have been reported by measuring a circumferential brittle zone introduced in cathodically charged specimens. 9 On the other hand, the embrittlement due to ‘residual’ hydrogen has not yet been observed, although a number of observations have been performed on the environmental embrittlement of Co3Ti alloys.1,2,9-‘4 In this study, some observations relevant to solubility, diffusion and trap site of hydrogen in Co3Ti alloys are reported. Also, to learn the effect of ‘residual’ hydrogen on ductility, a tensile test is carried out in liquid nitrogen using specimens thermally charged in hydrogen gas (i.e. with a variety of concentrations of hydrogen). Finally, a few implications are discussed concerning the environmental embrittlement of Co3Ti alloys.
INTRODUCTION The initial discovery that ordered intermetallic is very susceptible to embrittlement due to hydrogen decomposed from moisture in air at ambient temperature was made using LIZ-type Co3Ti polycrystals.‘,* The specimen deformed in air showed considerably lower tensile elongation than that deformed in a vacuum. Since this finding, so-called environmental embrittlement has been shown to take place in many environmental media containing moisture (or aqueous solution) at ambient temperature in a wide range of deformation rates and also to occur in many types of crystal structures, alloy systems and microstructures.3-8 For a better understanding of the embrittlement of ordered intermetallics, as well as ordinary materials, the hydrogen property in lattice (such as solubility, diffusion and site occupancy) is important. Also, an understanding of the embrittlement due to ‘residual’ hydrogen (i.e. hydrogen contained in lattice prior to deformation) is necessary to clarify how important a role hydrogen decomposed from moisture plays in the environmental embrittlement.
EXPERIMENTAL Co3Ti alloys with two nominal compositions of 23 mol%Ti and 21 mol%Ti were used in this study. Starting raw materials were 99.9 wt% cobalt and 443
444
T. Takasugi
99.9 wt% sponge titanium. Alloy buttons were made by arc melting in argon atmosphere on a copper hearth and then homogenizing in a vacuum at 1323 K for 2 days. These alloy buttons were repeatedly rolled at 773 K in air and annealed in vacuum at 1273 K for 5 h. The final thickness was prepared to be about 1.Omm. Next, small punch (SP) specimens, and thermal absorption and desorption specimens with a dimension of 10 x 10 mm were prepared by a precision wheel cutter. All of these specimens were then annealed in vacuum at 1273 K for 5 h to prepare a well-ordered microstructure. The average grain size of all the specimens prepared was quite similar, i.e. ranging from 6 to 8 pm. Finally, the SP specimens, and the absorption and desorption specimens, were abraded into O-5 and 0.25mm thickness by SIC paper and alumina particle, respectively. Hydrogen was thermally charged in one atmospheric H2 gas at 473 K for desired times. The charged specimens were quenched in an oil bath, the temperature of which was held at room temperature, they were rinsed in an acetone solution by an ultrasonic cleaner, then kept at room temperature. Hydrogen concentration in the charged specimens was measured by an argon carrier fusion-thermal conductivity method. Also, X-ray diffraction measurement was performed on the hydrogen-charged and uncharged specimens using Cu Ka radiation. Next, a hydrogen thermal desorption test was performed on the specimens preceded by thermal charging in H2 gas at 473 K for 4 h. The hydrogen-charged specimen was placed in a quartz chamber and then evacuated to be below 10p6Pa at liquid nitrogen temperature using a turbo-molecular pump combined with an oil rotary pump. Three different heating rates (i.e. 0.5, 1.0 and 2.0 Kmin-‘) were chosen and the desorbed hydrogen gas was monitored in a temperature range between room temperature and 1200 K by quadrapole mass spectrometer (QMS) (in the condition of an emission current of 2.0mA and a sampling time of 60 s). Evaluation of ductility (and toughness) of Co3Ti alloys was done by SP test in liquid nitrogen (i.e. at 77K). This test in liquid nitrogen was carried out in order to exclude an entry of hydrogen decomposed from moisture into the materials (i.e. to observe only the effect of the ‘residual’ hydrogen). Details of the SP test, which is a kind of bulge test, were described in the previous articles.9~13~‘4Ductility was evaluated by SP fracture energy which is obtained by calculating an area beneath a loaddisplacement curve. It has been well accepted that
et al. there is a good correlation between the magnitude of the SP fracture energy and the degree of toughness. Fracture surfaces of the deformed specimens were observed by a scanning electron microscope (SEM).
RESULTS Figure 1 shows variations of hydrogen concentration with hydrogen gas exposure time for Co-23 mol%Ti and Co-21 mol%Ti alloys. Hydrogen concentration for both alloy compositions parabolically increased with increasing exposure time. Hydrogen concentration of the Co-23 mol%Ti alloy increased more rapidly than that of the Co21 mol%Ti alloy, perhaps indicating that hydrogen diffusion in the former alloy composition is higher than that in the latter. Here, it is noted that hydrogen concentration still increases within the exposure time measured (i.e. < 3 x lo4 s). Therefore, on the basis of an assumption x = fi, the hydrogen diffusivity at 473 K in the Co3Ti alloy is estimated to be less than 5.2 x lo-i3(m2 s-i). This value may be coherent to the hydrogen diffusivity (i.e. 1.1 x 10-‘4(m2s-‘) at 295 K) reported in the Co-21.5 mol%Ti alloy by measuring the circumferential brittle zone introduced in the cathodically charged specimen.9 Figure 2 shows the X-ray diffraction profiles from the hydrogen charged (Fig. 2(a)) and uncharged (Fig. 2(b)) specimens of the Co-23mol%Ti alloy. In both specimens, all reflection lines were indexed as due to an L12 structure. No additional reflection lines were observed. Thus, Co3Ti hydride
+ -D_
f
-0
5
10
15 Time
/
Co-23mol%Ti Co-21 mol%Ti
20
25
30
ks
Fig. 1. Variations of hydrogen concentration with hydrogen gas exposure time for Co-23mol%Ti and Co-21 mol%Ti alloys. Hydrogen gas exposure was carried out in one atmospheric H2 gas at 473 K.
Hydr#ogenproperty in lattice associated with the embrittlement of Co3Ti alloys
temperatures above about 600K, hydrogen desorption remained at a very low level. An increase in the heating rate resulted in an increase in the temperature (Tc) showing the maximum in the desorption vs temperature curve. Here, an activation energy, E, for the hydrogen desorption from trap site (or soluble position) can be derived on the basis of the following equation: I5
ia) Hydrogen-charged
(E,a/RTz)
*P (b) Uncharged
I 30
I
I
I
I
I
I
40
50
60
70
60
90
28
100
I deg.
Fig. 2. X-ray diffraction profiles of (a) hydrogen-charged (b) uncharged Ce23 mol%Ti alloys. Hydrogen charging done at 473 K for 4 h.
and was
was not detected in th[e hydrogen-charged specimen. However, it is apparent that the peak position of each reflection line shifted toward a lower angle and therefore the lattice parameter of Ll;? structure in the hydrogen-charged specimen increased by a few per cent. Quite the same result was observed in the Co-21 mol%Ti allo:y. Figure 3 shows the hydrogen desorption curve measured at three different heating rates in the Co-23 mol%Ti alloy. Hydrogen desorption measured from Hz ion current rapidly increased with increasing temperature and reached a maximum at about 430K, followed by a rapid decrease. At
.:. .-,
Q:
E F 5 0
A
10-g
:
I
445
= A exp( -E,/RT’)
(1)
where a! is the heating rate, R is the gas constant, T, is a maximum temperature, and A is constant. The Arrhenius plot of eqn (1) i.e. In (a/T:) vs l/T, is shown in Fig. 4. From this figure, the activation energy of hydrogen desorption from trap site (or soluble position) in the Co3Ti alloy was estimated to be about 48 kJmol-i. If one postulates that the binding energy of hydrogen is the difference in the activation energies between the desorption and lattice diffusion processes, approximately 13 kJmol_’ can be derived as the binding energy of hydrogen with trap site in the Co3Ti alloy (here, the value of 35 kJmol19 was used as the activation energy for lattice diffusion of hydrogen). A different interpretation for E, is also possible; this value is almost identical to the activation energy for lattice diffusion of hydrogen, therefore the desorption process is rate-controlled by the lattice diffusion of hydrogen itself. It is interesting to learn how the residual hydrogen affects the mechanical property of the CosTi alloy. Figure 5 shows the variation of SP fracture energy with hydrogen concentration for the Co-23 mol%Ti and Co21 mol%Ti alloys which were hydrogen-charged for a desired time (i.e. hydrogen concentration) at 473K. It is apparent that the SP fracture energy in the Co-21 mol%Ti
1o-6 \
Heating rate
.. -.
1\1 ‘.:..
---0SWmin P.OWmin l.OWmin
:r
10‘10
10-1'
5 I"
i:
10‘12
,o-13
-,
200
’
400
600
Temperature
’
600
I
’
1000
1200
1o-8 ‘. ’ . . . ** ’ ’ ’ 2.0 2.1
2.2
IIT,
/ K
Fig. 3. The effect of heating rate on variation of H2 ion current with temperature for Co-23 mol%Ti alloy.
..* ’ ’
Fig. 4. Arrhenius alloy.
’
- ’ ’ ’ ’ . ’ ’ . ’ ’ ’ ’ ’ ’ s’ * ’ ’
2.3
2.4
; 5
/ 1O-3 K-’
plot of a/Tc2 vs l/Tc for Co_23mol%Ti
T. Takasugi
446
et al. Co-23 mol%Ti alloy was moderately dependent on the hydrogen content, that is, it decreased with an increase in the hydrogen concentration. Measured SP fracture energy was reduced from 2.5 J (at 4mass ppm H) to 0.4 J (at 75 mass ppm H). Corresponding to this change, the brittle grain boundary fracture mode, accompanied with smooth facets, became more dominant in SEM fractography (Fig. 6).
20
40
Hydrogen
Concentration
0
60 /
80
DISCUSSION
mass ppm H
Fig. 5. Variation
of SP fracture energy with hydrogen concentration for Co-23 mol%Ti and Cc~21 mol%Ti alloys. Specimens were hydrogen-charged for desired time at 473 K. Tests were done in liquid nitrogen at a cross-head speed of 0.2mmmin-‘.
alloy was consistently very high (-6 J) in the prepared hydrogen content (or slightly decreased with increasing hydrogen concentration). SEM fractography showed mainly ductile transgranular fracture mode with dimple pattern (Fig. 6). On the other hand, the SP fracture energy in the
Fig. 6. SEM fractography of (a) Co-23mol%Ti and (b) C-21 mol%Ti alloys containing 70 mass ppm H and 75 mass ppm H, respectively. Tests were done in liquid nitrogen (at 77 K).
The hydrogen diffusivity (at room temperature) in Co3Ti alloys estimated in this study and also reported in a previous study9 is similar to the value measured in pure Ni (6.9 x 10-14m2s-‘)‘6 and the value extrapolated from high temperature data measured in (Co,Fe)3V alloy with Liz structure (1 .l x 10-14m2 s-‘)17 but smaller than that measured in Fe (7.2 x 10-9m2s-‘)‘8 and N&Al alloy with L12 structure (1.2 x 10-‘0m2s-‘).‘9 Here, data for hydrogen diffusion shown in the Co3Ti alloy reveal that hydrogen causing the environmental embrittlement is not transported via lattice but via short circuit path (e.g. dislocations or grain boundaries) to grain boundaries, or otherwise, by another transport mode such as dislocation sweeping. If 1.1 x 10-14m2 s-l is chosen as the lattice diffusivity of hydrogen in the CosTi alloy, hydrogen is transported only by 3.1 x lOA mm into the alloy interior. This distance is too short to cause whole brittle fracture during testing time, i.e. mostly less than 15 min. In this study, even though a large amount of hydrogen is charged into the Co3Ti alloys, hydride was not formed. This result indicates that embrittlement by residual hydrogen as well as penetrating hydrogen is caused by atomic hydrogen. Of course, an enrichment of a high level of hydrogen at grain boundaries is required to result in embrittlement, as will be discussed. Also, it was not determined whether hydrogen is trapped to preferential sites or exists in the soluble form in the lattices. In the former case, grain boundaries may be considered as the possible trap site of this alloy. In the latter case, octahedral sites (O-site) and tetragonal sites (T-sites)20-22 are considered. In iron and steel, the value of 58.6 kJ mol-’ has been estimated as the binding energies of hydrogen with dislocation and grain boundaries.23 Further investigations are needed to determine the preferential trap site or soluble position of hydrogen in this alloy.
Hydrogen
property
in lattice associated
The ductility (i.e. the toughness) of the Co23mol%Ti alloy was dependent on the ‘residual’ hydrogen content, as shown in Fig. 5. In the test in liquid nitrogen, absorption of hydrogen from air is impossible and in addition the residual hydrogen is immobile. These facts imply that the immobile hydrogen enriched to grain boundaries prior to deformation (in stead of penetrating hydrogen) certainly affects the grain boundary cohesion and the associated plasticity, resulting in the grain boundary separation. On the other hand, the ductility of the Co--21 mol%Ti alloy was independent of the ‘residu.al’ hydrogen content and very high even in a. high hydrogen content (N 70mass ppm H). This result means that the alloy stoichiometry plays an important role in the grain boundary cohesion and the associated grain boundary embrittlement by hydrogen. It has been suggested by a number of experimental observations24*25 and theoretical calculations26,27 in L12 structure that a grain boundary consisting of stoichiometric alloy composition has a more distorted structure (like atomic cavities) than that consisting of off-stoichiometric alloy composition, resulting in reduced grain boundary cohesion. Consequently, the critical hydrogen content,r4 which is defined as the lowest hydrogen content at grain boundary causing intergranular fracture, appears to be low and therefore very sensitive to the residual hydrogen content in lattice of near stoichiometric alloy composition (the Co-23mol%Ti alloy in this study). On the other hand, the critical hydrogen content in the Co-21 mol%Ti alloy appears to be high. Thus, the results indicating ‘,2,9-‘4 that the off-stoichiometric CosTi alloy (i.e. the Co-21 21.5 mol%Ti alloys) has been more resistant to environmental embrittlement are suggested to be attributable to higher critical hydrogen content (i.e intrinsically higher grain boundary cohesion) in addition to a lower decomposition rate of water vapor on the alloy surfac~e.28Finally, it is noted that the Co3Ti alloy has an intrinsically high fracture toughness if the residual1 hydrogen as well as the penetrating hydrogen are excluded from materials.
CONCLUSION Some observations relating to solubility, diffusion and trap site of hydrogen in Co3Ti alloy were carried out. Also, the effect of ‘residual’ hydrogen on ductility was observed by a small punch (SP) test in liquid nitrogen, using hydrogen-charged specimens. The following results were obtained by the present study:
with the embrittlement
of CojTi alloys
447
1. Hydrogen diffusivity at room temperature in the CosTi alloy was evaluated to be low as in pure Ni. 2. The desorption curve of hydrogen in the Co3Ti alloy reached the maximum at N 430 K and corresponding activation energy was estimated to be 48 kJmol-i. 3. SP fracture energy in the Co-23 mol%Ti alloy was dependent on the hydrogen content and decreased with increasing hydrogen content, while that in the Co-21 mol%Ti alloy was independent of the hydrogen content and showed consistently high values. 4. It is considered that the environmental embrittlement of the CosTi alloy requires a hydrogen transport mode other than via lattice diffusion. Also, the difference in the environmental embrittlement between the two prepared alloy compositions is primarily due to the difference in the critical hydrogen content causing grain boundary fracture (i.e. intrinsic grain boundary cohesion).
ACKNOWLEDGEMENTS This work was supported in part by the New Energy and Industrial Technology Development Organization (NEDO) in Japan. The authors are grateful to Nikko Inspection Service Co., Ltd. for hydrogen analysis.
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