- Email: [email protected]
Hydrogen embrittlement of MFR candidate vanadium alloys
Journal of Nuclear North-Holland
Materials
179-181
179
(1991) 779-782
Hydrogen em~~ttl~ment
of MFR candidate vanadium alloys
S. Yano, M. Tada and H. Math Institute for Materials Research, Tohoku University, 2-l-I Karahira, Aoba-ku, Sendai 980, Japan
Pure vanadium. V-3Ti-1Si and V-lSCr-5Ti alloys were K. The fracture surfaces of specimens were examined by strengthened, but had good ductility even at 77 K. The ductility loss at 150 K for V-3Ti-lSi, at 200 K for pure V
resistance to hydrogen embrittlement
charged with 0.3 at% H, and tested in tension between 77 and 520 SEM. Both non-hydrogenated alloys were appreciably solutionspecimen hydrogenated with 0.3 at% H exhibited a precipitous
and at 300 K for V-lSCr-5Ti. The V-3Ti-1Si alloy has a greater than pure vanadium. On the other hand, the V-15Cr-5Ti alloy was most severely
deteriorated by hydrogen and always fractured in a brittle in both alloys was discussed in terms of the inhomogeneity coefficient and solubility of hydrogen.
1. Introduction
Vanadium alloys have several advantages over other alloys for use as structural materials in fusion reactors [l-3]. While their low degree of activation, good radiation damage resistance and high strength at high temperatures make these alloys attractive, their sensitivity to interstitial impurities causes a major concern. Hydrogen emb~ttlement is an especially important ‘problem because the first wall or blanket materials are always exposed to hydrogen isotopes in the coolant or breeding materials. Vanadium and its binary alloys, e.g. V-Cr [4] or V-Ti [S] have been extensively studied with respect to this subject, but no detailed investigation has been carried out on candidate alloys such as V-3Ti-1Si and V-15Cr-5Ti alloys. The latter alloy has a larger susceptibility to hydrogen embrittlement than the former [6]; the ductility of V-15Cr-5Ti alloy was severely deteriorated by a small amount of hydrogen picked up during cutting or polishing with water, but the reason for the severe embrittlement of this alloy was not clear. The purpose of the present work is to investigate the effect of hydrogen on the mechanical properties of the V-3Ti-1Si and V-lSCr-5Ti alloys and to clarify the different behaviors of embrittlement caused by hydrogen.,
mode at temperatures lower than 300 K. Hydrogen emb~ttiement of deformation on yielding and the solute effect on the diffusion
gauge size of 1.2 mm x 5 mm. These samples were electro-polished, wrapped with tantalum and zirconium foils, and annealed in an evacuated quartz capsule for 2 h at 1327 K for the V-3Ti-1Si alloy, and at 1373 K for pure V and the V-15Cr-5Ti alloy. The average grain diameter is 100, 30, and 5 pm in pure V, V-3Ti-1Si and V-lSCr-STi, respectively. Hydrogenation of specimens was carried out by Sievert’s method. The quartz tube containing samples was evacuated up to a vacuum of lo-’ Pa and isolated from the pumping system. Hydrogen from the thermal decomposition of titanium hydride was admitted at predetermined pressures, and the quartz tube was held at 1123 K for 2 h and then allowed to furnace cool. Hydrogen content was determined to be (0.3 + 0.03) at% from the pressure change before and after hydrogenation. Tensile tests were performed on a screw driven tensile testing machine at a constant strain rate of 1.67 X 10e3 s-’ in a temperature range between 77 and 520 K. Test temperatures were obtained by immersing the specimens in suitable constant temperature baths ( f 3 K).
a
_
a
U Vanadium 0
Wl?lSi
2. Experimental procedures A pure vanadium ingot was prepared from 99.9% pure dendritic vanadium by electron beam melting. This ingot was also used to obtain a V-3Ti-1% alloy ingot by argon arc melting. These ingots were reduced to a finai thickness of 0.25 mm by cold rolling. V-lSCr-STi alloy was kindly provided as a 0.75 mm thick sheet by Dr. D.L. Smith of Argonne National Laboratory. This alloy sheet was further rolled to 0.25 mm in thickness. These sheets were punched into tensile specimens with a 0022-3115/91/$03.50
0 1991 - Elsevier
Science Publishers
200
TEMPERATURE
K
Fig. 1. Temperature dependence of yield stress for non-hydrogenated pure V and its alloys.
B.V. (North-Holland)
780
S. Yano et al. / H~vdrogen emhrittlemenr
3. Results
of MFR vanadium
alloys
30 1
The temperature dependence of yield stress and elongation of pure V, and of V-3TiLlSi and V-ISCr5Ti alloys are shown in figs. 1 and 2. Both alloys are appreciably strengthened by solid solution hardening and grain size effect. The small values of elongation for pure V at 77 and 113 K are due to local necking by plastic instability. so that all the unhydrogenated samples remained ductile down to the lowest temperature in the present experiment. The temperature dependence of elongation of the hydrogenated specimens is shown in fig. 3. The specimen hydrogenated with 0.3 at’% H exhibited a precipitous ductility loss below 150, 200 and 300 K for V-3TiLlSi, pure V and V-15Cr-5Ti. respectively. At still lower temperatures, however, pure V as well as the V-3Ti-1Si alloy showed a ductility return. When the specimen was tested at the temperature of minimum ductility, fracture occurred in a cleavage mode for pure V and V-lSCr-5Ti alloy, but in a dull fibrous tear mode with a small amount of cleavage facet for the V- 3Ti-1Si alloy. 4. Discussion Hydrogen embrittlement of pure V and its alloys has been understood in terms of the stress-induced formation of hydride ahead of advancing crack tips [7-91. Fig. 4 schematically shows the shift of the solvus temperature by stress and the temperature dependence of the corresponding fracture strain (r,) and fracture mode [8]. It should be noted that the stress level is usually not large enough to appreciably influence the solvus temperature other than at stress concentrated sites such as the crack tip. Most metals generally have no pre-induced crack. and brittle fracture without macroscopic strain can occur only just above the stress-free solvus temperature. Actually. pure V with 0.3 at% H fractured in a brittle mode near the solvus temperature (indicated
TEMPERATURE
K
Fig. 3. Temperature dependence of elongation genated specimens of pure V (square), V-3TiLlSi VP1S<‘rm5Ti (triangle).
for hydro(circle) and
by the arrow in fig. 3) which is estimated from Owen and Scott’s solubility data [lo]. The fracture mode can also be cleavage at temperatures far higher than the solvus. provided that a large amount of deformation is accumulated before fracture occurs [II]. For example, pure V with 0.3 at% H tested at 250 K fractured in a cleavage mode at cf of 20% (fig. 3). Hydrogen solubility data for the V-15Cr-5Ti alloy is not currently available. An internal friction measurement on the same specimen containing 0.3 at% H has been conducted. Although the result is still not conclusive due to some scatter in the data, the solvus temperature determined as the onset point of the precipitation peak of hydride was not very different from that of pure vanadium. Solvus temperature may also be estimated from the irregularity in the temperature dependence of yield stress. Fig. 5 shows the difference in yield stress between the hydrogenated and non-hydrogenated specimens as a function of temperature. The strengthening by hydrogen is known [4.5] to be prominent near the
Ductile
Increasing stress b
4
,//l//’
I’ ,’ /’ ,I /I ,’ /’ ,/ ,I ,I /’ ,’ /’ I /’ .;
/I
8Q
100
200
300
TEMPERATURE
400
600
600
K
Fig. 2. Temperature dependence of elongation genated pure V and its alloys.
for non-hydro-
Cleavage
..$J & ? 2 t E” F
/
Ductile
L 0
7 1,
iCH
HN
-;-Ef
Fig. 4. Illustrations showing stress effects on the solvus temperature and the temperature dependence of the corresponding fracture strain (z, ) and fracture mode.
S. Yano et al. / Hydrogen embrittlement
ofMFR
781
vanadium alloys
CI Vanadium
200
$ 150
z
0
V-37%1Si
A
V-i 5Cr-51
b
a ioo
50
0
100 100
200 300 TEMPE?IATURE K
400
Fig. 5. Temperature dependence of the difference in yield stress between hydrogenated and non-hydrogenated specimens; the arrow showing the real. value is larger than the plotted one, since the hydrogenated specimen brittle fractured below the proportional limit.
0
tom
Fig. 6. Stress-strain curves for the V-15Cr-5Ti selected temperatures; (a) non-hydrogenated, genated.
G
alloy tested at (b) hydro-
40
69
so
AQL
100
120
M?a
Fig. 7. Temperature dependence of Luders strain and stress concentration formed at the front of Luders band in V-lSCr5Ti alloy; solid symbols; hydrogenated, open symbols; non-hydrogenated.
I 150K
20
Au,
solvus temperature. This is in agreement with the present observation in pure V. The strengthening of the V-lSCr-5Ti is slightly enhanced below 150 K. This suggests the existence of the solvus temperature around 150 K, which is much lower than the solvus estimated from internal friction. Anyhow, it can be concluded that the solvus temperature is located far below the ductileto-brittle transition temperature of this alloy, i.e. around 300 K. This requires further discussion since the stress level is considered to be too small to induce this much shift of solvus temperature. In V-lSCr-STi alloy, even a small amount of deformation immediately after yielding may produce a large stress concentration because of the large yield strength and the inhomogenity of deformation. Whether the alloy was hydrogenated or not, the yielding proceeded by the growth of Luders bands as shown in fig. 6. In Luders deformation, strain is confined in Luders bands
200K
A A
500
where work hardening of Au, occurs after a strain of et is accumulated. Such a yielding is generally observed when the grain size effect on yield strength, i.e. the unlocking stress for dislocations, is large and the deformation can only propagate by the help of an additional stress concentration of Aa, at the front of the Luders bands. Fig. 7 shows the temperature dependence of the values et_ and Au,. This Au, is considered to be the measure of the stress required to unlock the dislocations at the Luders band front. In other words, some multiples of Au,, i.e. KAcT~ is the net stress required to unlock the dislocations in the undeformed grain, where the multiplication factor K is 2 at most. As shown in fig. 7, the value Au, increases by decreasing temperature, while it is only 80 MPa at most. In our previous study [12], we estimated the change in the solvus temperature in pure V to be less than 3 K by the application of stress of the order of the yield stress (150 MPa). This estimation shows that even the stress concentration at Luders fronts is not large enough to shift the solvus appreciably. There seems to be three possible explanations for the embrittlement of V-lSCr-STi alloy at temperatures far higher than the solvns. Firstly, the stress concentration may be much larger than the value estimated here caused by, e.g., dislocation pile-up. Secondly, the solvus shift may be much larger than the estimation here. Although the large solvus shift by stress observed in vanadium in the previous study [12] is not yet fully understood, the unexpectedly large shift may explain the high DBTT observed in the V-15Cr-5Ti alloy. Thirdly, hydride precipitation may not always be necessary for the embrittlement; some hydrogen clustering, for example, may be sufficient to cause embrittlement and the clustering is possible at temperatures much
782
S. Yano et al. / Hydrogen embrittlement
higher than the solvus for hydride precipitation. Anyway, it is very likely that the severe susceptibility of V-15Cr-5Ti alloy to hydrogen embrittlement is at least partly related to the inhomogeneous deformation of this alloy. In pure V, which deformed homogeneously on yielding, the abrupt ductility loss occurred near the solvus temperature. The abrupt ductility loss of the V-3Ti-1Si alloy occurred at a temperature lower than that of pure V. The better ductility of V-3Ti-1Si is caused by a lowering of the solvus temperature due to Ti solute, as has been reported in V-Ti binary alloys [13]. Besides, this alloy held a considerable elongation in the temperature range of minimum ductility. This is contrasted with pure V, which fractured in a brittle mode near 200 K. The mitigation of the embrittlement in the alloy is also attributed to the effect of Ti solute; it decreases the diffusion coefficient of hydrogen and therefore it is difficult for hydrogen to diffuse toward a crack tip [12]. 5. Conclusion Mechanical properties of pure vanadium, V-3Ti-lSi, and V-15Cr-5Ti containing 0.3 at% hydrogen have been studied in a temperature range between 77 and 520 K. (1) Both non-hydrogenated alloys are appreciably strengthened by solution hardening and gram size effect, while they remain ductile even at 77 K. (2) The hydrogenated specimen exhibited ductile-tobrittle transition at 150 K for V-3Ti-lSi, at 200 K for pure V and at 300 K for V-15Cr-5Ti. (3) At still lower temperatures, pure V as well as V3Ti-1Si show a ductility return.
of MFR
vanadium alloys
alloy held a considerable elongation in (4) V-3Ti-1Si the temperature range of minimum ductility. This greater resistance to hydrogen embrittlement is attributed to the effect of Ti in reducing the solvus temperature and the diffusivity of hydrogen. alloy is most severely embrittled by (5) V-15Cr-5Ti hydrogen. The reason for this is tentatively attributed to the stress concentration at the advancing Luders front in assisting hydride formation.
References
111 D.L. Smith, J. Nucl. Mater. 122 & 123 (1984) 51. [21 D.L. Smith, B.A. Loomis and D.R. Diercks, J. Nucl. Mater. 135 (1985) 125. [31 D.R. Diercks and B.A. Loomis, J. Nucl. Mater. 141-143 (1986) 1117. [41 C.V. Owen, W.A. Spitzig and 0. Buck, Metall. Trans. 18A (1987) 1593. C.V. Owen, 0. Buck and T.J. Rowland, J. iSI W.A.Spitzig, Less-Con-m. Met. 115 (1985) 45. 161 B.A. Loomis, R.H. Lee, D.L. Smith and J.R. Peterson. J. Nucl. Mater. 155-157 (1988) 631. [71 D.G. Westlake, Trans. ASM 62 (1969) 1000. and H.K. Bimbaum, Acta Metall. 25 PI M.L. Grossbeck (1977) 135. [91 S. Takano and T. Suzuki, Acta Metall. 22 (1974) 265. WV C.V. Owen and T.E. Scott, Metall. Trans. 3 (1972) 1715. and H.K. Bimbaum, Acta Met. [ill S. Gahr, M.L. Grossbeck 25 (1977) 125. P21 H. Matsui, 0. Kubota and M. Koiwa, Acta Metall. 34 (1986) 295. u31 S. Tanaka and H. Kimura, Trans. JIMIS 21 (1980) 513.
Materials
179-181
179
(1991) 779-782
Hydrogen em~~ttl~ment
of MFR candidate vanadium alloys
S. Yano, M. Tada and H. Math Institute for Materials Research, Tohoku University, 2-l-I Karahira, Aoba-ku, Sendai 980, Japan
Pure vanadium. V-3Ti-1Si and V-lSCr-5Ti alloys were K. The fracture surfaces of specimens were examined by strengthened, but had good ductility even at 77 K. The ductility loss at 150 K for V-3Ti-lSi, at 200 K for pure V
resistance to hydrogen embrittlement
charged with 0.3 at% H, and tested in tension between 77 and 520 SEM. Both non-hydrogenated alloys were appreciably solutionspecimen hydrogenated with 0.3 at% H exhibited a precipitous
and at 300 K for V-lSCr-5Ti. The V-3Ti-1Si alloy has a greater than pure vanadium. On the other hand, the V-15Cr-5Ti alloy was most severely
deteriorated by hydrogen and always fractured in a brittle in both alloys was discussed in terms of the inhomogeneity coefficient and solubility of hydrogen.
1. Introduction
Vanadium alloys have several advantages over other alloys for use as structural materials in fusion reactors [l-3]. While their low degree of activation, good radiation damage resistance and high strength at high temperatures make these alloys attractive, their sensitivity to interstitial impurities causes a major concern. Hydrogen emb~ttlement is an especially important ‘problem because the first wall or blanket materials are always exposed to hydrogen isotopes in the coolant or breeding materials. Vanadium and its binary alloys, e.g. V-Cr [4] or V-Ti [S] have been extensively studied with respect to this subject, but no detailed investigation has been carried out on candidate alloys such as V-3Ti-1Si and V-15Cr-5Ti alloys. The latter alloy has a larger susceptibility to hydrogen embrittlement than the former [6]; the ductility of V-15Cr-5Ti alloy was severely deteriorated by a small amount of hydrogen picked up during cutting or polishing with water, but the reason for the severe embrittlement of this alloy was not clear. The purpose of the present work is to investigate the effect of hydrogen on the mechanical properties of the V-3Ti-1Si and V-lSCr-5Ti alloys and to clarify the different behaviors of embrittlement caused by hydrogen.,
mode at temperatures lower than 300 K. Hydrogen emb~ttiement of deformation on yielding and the solute effect on the diffusion
gauge size of 1.2 mm x 5 mm. These samples were electro-polished, wrapped with tantalum and zirconium foils, and annealed in an evacuated quartz capsule for 2 h at 1327 K for the V-3Ti-1Si alloy, and at 1373 K for pure V and the V-15Cr-5Ti alloy. The average grain diameter is 100, 30, and 5 pm in pure V, V-3Ti-1Si and V-lSCr-STi, respectively. Hydrogenation of specimens was carried out by Sievert’s method. The quartz tube containing samples was evacuated up to a vacuum of lo-’ Pa and isolated from the pumping system. Hydrogen from the thermal decomposition of titanium hydride was admitted at predetermined pressures, and the quartz tube was held at 1123 K for 2 h and then allowed to furnace cool. Hydrogen content was determined to be (0.3 + 0.03) at% from the pressure change before and after hydrogenation. Tensile tests were performed on a screw driven tensile testing machine at a constant strain rate of 1.67 X 10e3 s-’ in a temperature range between 77 and 520 K. Test temperatures were obtained by immersing the specimens in suitable constant temperature baths ( f 3 K).
a
_
a
U Vanadium 0
Wl?lSi
2. Experimental procedures A pure vanadium ingot was prepared from 99.9% pure dendritic vanadium by electron beam melting. This ingot was also used to obtain a V-3Ti-1% alloy ingot by argon arc melting. These ingots were reduced to a finai thickness of 0.25 mm by cold rolling. V-lSCr-STi alloy was kindly provided as a 0.75 mm thick sheet by Dr. D.L. Smith of Argonne National Laboratory. This alloy sheet was further rolled to 0.25 mm in thickness. These sheets were punched into tensile specimens with a 0022-3115/91/$03.50
0 1991 - Elsevier
Science Publishers
200
TEMPERATURE
K
Fig. 1. Temperature dependence of yield stress for non-hydrogenated pure V and its alloys.
B.V. (North-Holland)
780
S. Yano et al. / H~vdrogen emhrittlemenr
3. Results
of MFR vanadium
alloys
30 1
The temperature dependence of yield stress and elongation of pure V, and of V-3TiLlSi and V-ISCr5Ti alloys are shown in figs. 1 and 2. Both alloys are appreciably strengthened by solid solution hardening and grain size effect. The small values of elongation for pure V at 77 and 113 K are due to local necking by plastic instability. so that all the unhydrogenated samples remained ductile down to the lowest temperature in the present experiment. The temperature dependence of elongation of the hydrogenated specimens is shown in fig. 3. The specimen hydrogenated with 0.3 at’% H exhibited a precipitous ductility loss below 150, 200 and 300 K for V-3TiLlSi, pure V and V-15Cr-5Ti. respectively. At still lower temperatures, however, pure V as well as the V-3Ti-1Si alloy showed a ductility return. When the specimen was tested at the temperature of minimum ductility, fracture occurred in a cleavage mode for pure V and V-lSCr-5Ti alloy, but in a dull fibrous tear mode with a small amount of cleavage facet for the V- 3Ti-1Si alloy. 4. Discussion Hydrogen embrittlement of pure V and its alloys has been understood in terms of the stress-induced formation of hydride ahead of advancing crack tips [7-91. Fig. 4 schematically shows the shift of the solvus temperature by stress and the temperature dependence of the corresponding fracture strain (r,) and fracture mode [8]. It should be noted that the stress level is usually not large enough to appreciably influence the solvus temperature other than at stress concentrated sites such as the crack tip. Most metals generally have no pre-induced crack. and brittle fracture without macroscopic strain can occur only just above the stress-free solvus temperature. Actually. pure V with 0.3 at% H fractured in a brittle mode near the solvus temperature (indicated
TEMPERATURE
K
Fig. 3. Temperature dependence of elongation genated specimens of pure V (square), V-3TiLlSi VP1S<‘rm5Ti (triangle).
for hydro(circle) and
by the arrow in fig. 3) which is estimated from Owen and Scott’s solubility data [lo]. The fracture mode can also be cleavage at temperatures far higher than the solvus. provided that a large amount of deformation is accumulated before fracture occurs [II]. For example, pure V with 0.3 at% H tested at 250 K fractured in a cleavage mode at cf of 20% (fig. 3). Hydrogen solubility data for the V-15Cr-5Ti alloy is not currently available. An internal friction measurement on the same specimen containing 0.3 at% H has been conducted. Although the result is still not conclusive due to some scatter in the data, the solvus temperature determined as the onset point of the precipitation peak of hydride was not very different from that of pure vanadium. Solvus temperature may also be estimated from the irregularity in the temperature dependence of yield stress. Fig. 5 shows the difference in yield stress between the hydrogenated and non-hydrogenated specimens as a function of temperature. The strengthening by hydrogen is known [4.5] to be prominent near the
Ductile
Increasing stress b
4
,//l//’
I’ ,’ /’ ,I /I ,’ /’ ,/ ,I ,I /’ ,’ /’ I /’ .;
/I
8Q
100
200
300
TEMPERATURE
400
600
600
K
Fig. 2. Temperature dependence of elongation genated pure V and its alloys.
for non-hydro-
Cleavage
..$J & ? 2 t E” F
/
Ductile
L 0
7 1,
iCH
HN
-;-Ef
Fig. 4. Illustrations showing stress effects on the solvus temperature and the temperature dependence of the corresponding fracture strain (z, ) and fracture mode.
S. Yano et al. / Hydrogen embrittlement
ofMFR
781
vanadium alloys
CI Vanadium
200
$ 150
z
0
V-37%1Si
A
V-i 5Cr-51
b
a ioo
50
0
100 100
200 300 TEMPE?IATURE K
400
Fig. 5. Temperature dependence of the difference in yield stress between hydrogenated and non-hydrogenated specimens; the arrow showing the real. value is larger than the plotted one, since the hydrogenated specimen brittle fractured below the proportional limit.
0
tom
Fig. 6. Stress-strain curves for the V-15Cr-5Ti selected temperatures; (a) non-hydrogenated, genated.
G
alloy tested at (b) hydro-
40
69
so
AQL
100
120
M?a
Fig. 7. Temperature dependence of Luders strain and stress concentration formed at the front of Luders band in V-lSCr5Ti alloy; solid symbols; hydrogenated, open symbols; non-hydrogenated.
I 150K
20
Au,
solvus temperature. This is in agreement with the present observation in pure V. The strengthening of the V-lSCr-5Ti is slightly enhanced below 150 K. This suggests the existence of the solvus temperature around 150 K, which is much lower than the solvus estimated from internal friction. Anyhow, it can be concluded that the solvus temperature is located far below the ductileto-brittle transition temperature of this alloy, i.e. around 300 K. This requires further discussion since the stress level is considered to be too small to induce this much shift of solvus temperature. In V-lSCr-STi alloy, even a small amount of deformation immediately after yielding may produce a large stress concentration because of the large yield strength and the inhomogenity of deformation. Whether the alloy was hydrogenated or not, the yielding proceeded by the growth of Luders bands as shown in fig. 6. In Luders deformation, strain is confined in Luders bands
200K
A A
500
where work hardening of Au, occurs after a strain of et is accumulated. Such a yielding is generally observed when the grain size effect on yield strength, i.e. the unlocking stress for dislocations, is large and the deformation can only propagate by the help of an additional stress concentration of Aa, at the front of the Luders bands. Fig. 7 shows the temperature dependence of the values et_ and Au,. This Au, is considered to be the measure of the stress required to unlock the dislocations at the Luders band front. In other words, some multiples of Au,, i.e. KAcT~ is the net stress required to unlock the dislocations in the undeformed grain, where the multiplication factor K is 2 at most. As shown in fig. 7, the value Au, increases by decreasing temperature, while it is only 80 MPa at most. In our previous study [12], we estimated the change in the solvus temperature in pure V to be less than 3 K by the application of stress of the order of the yield stress (150 MPa). This estimation shows that even the stress concentration at Luders fronts is not large enough to shift the solvus appreciably. There seems to be three possible explanations for the embrittlement of V-lSCr-STi alloy at temperatures far higher than the solvns. Firstly, the stress concentration may be much larger than the value estimated here caused by, e.g., dislocation pile-up. Secondly, the solvus shift may be much larger than the estimation here. Although the large solvus shift by stress observed in vanadium in the previous study [12] is not yet fully understood, the unexpectedly large shift may explain the high DBTT observed in the V-15Cr-5Ti alloy. Thirdly, hydride precipitation may not always be necessary for the embrittlement; some hydrogen clustering, for example, may be sufficient to cause embrittlement and the clustering is possible at temperatures much
782
S. Yano et al. / Hydrogen embrittlement
higher than the solvus for hydride precipitation. Anyway, it is very likely that the severe susceptibility of V-15Cr-5Ti alloy to hydrogen embrittlement is at least partly related to the inhomogeneous deformation of this alloy. In pure V, which deformed homogeneously on yielding, the abrupt ductility loss occurred near the solvus temperature. The abrupt ductility loss of the V-3Ti-1Si alloy occurred at a temperature lower than that of pure V. The better ductility of V-3Ti-1Si is caused by a lowering of the solvus temperature due to Ti solute, as has been reported in V-Ti binary alloys [13]. Besides, this alloy held a considerable elongation in the temperature range of minimum ductility. This is contrasted with pure V, which fractured in a brittle mode near 200 K. The mitigation of the embrittlement in the alloy is also attributed to the effect of Ti solute; it decreases the diffusion coefficient of hydrogen and therefore it is difficult for hydrogen to diffuse toward a crack tip [12]. 5. Conclusion Mechanical properties of pure vanadium, V-3Ti-lSi, and V-15Cr-5Ti containing 0.3 at% hydrogen have been studied in a temperature range between 77 and 520 K. (1) Both non-hydrogenated alloys are appreciably strengthened by solution hardening and gram size effect, while they remain ductile even at 77 K. (2) The hydrogenated specimen exhibited ductile-tobrittle transition at 150 K for V-3Ti-lSi, at 200 K for pure V and at 300 K for V-15Cr-5Ti. (3) At still lower temperatures, pure V as well as V3Ti-1Si show a ductility return.
of MFR
vanadium alloys
alloy held a considerable elongation in (4) V-3Ti-1Si the temperature range of minimum ductility. This greater resistance to hydrogen embrittlement is attributed to the effect of Ti in reducing the solvus temperature and the diffusivity of hydrogen. alloy is most severely embrittled by (5) V-15Cr-5Ti hydrogen. The reason for this is tentatively attributed to the stress concentration at the advancing Luders front in assisting hydride formation.
References
111 D.L. Smith, J. Nucl. Mater. 122 & 123 (1984) 51. [21 D.L. Smith, B.A. Loomis and D.R. Diercks, J. Nucl. Mater. 135 (1985) 125. [31 D.R. Diercks and B.A. Loomis, J. Nucl. Mater. 141-143 (1986) 1117. [41 C.V. Owen, W.A. Spitzig and 0. Buck, Metall. Trans. 18A (1987) 1593. C.V. Owen, 0. Buck and T.J. Rowland, J. iSI W.A.Spitzig, Less-Con-m. Met. 115 (1985) 45. 161 B.A. Loomis, R.H. Lee, D.L. Smith and J.R. Peterson. J. Nucl. Mater. 155-157 (1988) 631. [71 D.G. Westlake, Trans. ASM 62 (1969) 1000. and H.K. Bimbaum, Acta Metall. 25 PI M.L. Grossbeck (1977) 135. [91 S. Takano and T. Suzuki, Acta Metall. 22 (1974) 265. WV C.V. Owen and T.E. Scott, Metall. Trans. 3 (1972) 1715. and H.K. Bimbaum, Acta Met. [ill S. Gahr, M.L. Grossbeck 25 (1977) 125. P21 H. Matsui, 0. Kubota and M. Koiwa, Acta Metall. 34 (1986) 295. u31 S. Tanaka and H. Kimura, Trans. JIMIS 21 (1980) 513.