- Email: [email protected]
Microscopic observations of fracture behavior in a Ni55Pd35P10 metallic glass
Materials Science and Engineering, 23 (1976) 261 - 265
261
© Elsevier Sequoia S.A., Lausanne -- Printed in the Netherlands
Microscopic Observations of Fracture Behavior in a NissPdssP10 Metallic Glass*
S. TAKAYAMA Materials Research Center, Allied Chemical Corporation, Morristown, NJ 07960 (U.S.A.)
R. MADDIN Department of Metallurgy and Materials Science, University of Pennsylvania, Philadelphia, PA 19174 (U.S.A.)
SUMMARY
EXPERIMENTAL
Deformation and fracture of a NissPdssP10 metallic glass have been examined in a tensile stage in a transmission electron microscope. It is observed that a plastic zone appears at the crack tip in advance of crack propagation. The microscopic fracture profile shows two characteristic features: relatively sharp offsets (dislocation-like flow) and tearing behavior (viscous-type flow).
A powder metallurgy m e t h o d was used to prepare the Ni55PdssP1o ternary alloy. Ribbon filaments were obtained by quenching the molten alloy on to the wall of a rotating metallic drum [15]. These ribbon filaments were carefully examined by X-ray and electron diffraction techniques to determine whether or not they were amorphous. For transmission electron microscopy (TEM), the filaments were thinned by an ion-milling instrument. The in situ observations of the fracture behavior were made using a tensile stage m o u n t e d within the electron microscope.
INTRODUCTION Following the early studies of the mechanical properties of PdsoSi20 by Masumoto and Maddin [1] there have now been numerous reports on various amorphous alloys [2 - 17]. It has been shown that amorphous alloys deform in a localized manner, i.e., inhomogeneously, and that their tensile fracture surfaces are usually characterized by two distinct regions [1], a relatively featureless part [2] (shear offset) and a "vein" or "ridge" pattern. The formation of this latter pattern has been interpreted in terms of an adiabatic shear [2]. We present, herein, electron microscopic observations of the tensile mechanics of a Ni55PdssP1o metallic glass in order to contribute additional details to the understanding of the deformation and failure behavior.
*Paper presented at the Second International Conference on Rapidly Quenched Metals, held at the Massachusetts Institute of Technology, Cambridge, Mass., November 17 - 19, 1975.
PROCEDURE
RESULTS AND DISCUSSION Typical stress-strain curves for Ni55PdssPlo are presented in Fig. 1, (a) and (b), for room and liquid nitrogen temperatures, respectively. These curves show the same characteristics as other amorphous alloys, i.e., they show no well-defined yield point, but rather a small deviation from linearity before fracture. Some fracture surfaces are shown in Fig. 2 for a nominal strain rate of 5 × 10 -4 sec -1. The shear offset and the size of the vein cells decrease as the temperature decreases. Figure 2(c) shows that the vein morphology is curved, going from b o t t o m left to top right of the photos, i.e. along the applied torsion direction. A fracture surface morphology similar to amorphous alloys is often observed in crystalline materials. The processes producing the observed morphology have been explained via a dislocation model [19, 20]. Thus, Fig. 3 shows typical fracture surfaces in some crystalline materials ((a) shows a typical river pattern for quasi-brittle crystalline material [21],
262
:5xlO-4sec q
200 (a)
(b)
i
1 50
'E 1 O0
E
m 50
I
I I 1.0 STRAIN (,~L/L x 10 "2 )
Fig. 1. Stress-strain curves of Ni55Pd~5P10 metallic glasses, testing strain rate e = 5 × 10- sec-1.
while (b) shows ripple pattern for ductile failure [22] ). It is noteworthy that, compared with amorphous fracture surfaces, the fracture surface morphology of a quasi-brittle material is similar to a ductile glass (compare Fig. 2(a) with Fig. 3(a)), while that of a ductile crystalline material can look like that of a brittle amorphous one. Examination of the deformation of metallic glasses reveals that a small number of localized deformation bands appear prior to fracture. Moreover, these deformation bands are often observed near the fracture surface. It, therefore, appears that localized plastic flow contributed to produce those fracture patterns. Figure 4, (a)-(h), shows the crack propagat-ion sequence when a thin specimen of Ni55PdasP10 amorphous alloy was pulled to failure on the deformation stage of a transmission electron microscope. The crack propagates in order from (a) to (h) in Fig. 4 with increasing applied tensile stress. Thus, initially, crack growth proceeds slowly in a stable manner with increasing stress after the plastic zones are formed ahead of the crack (as shown in (a)-(c)). In (d) another void appears ahead of the crack. (The arrow indicates the part connecting the void with the crack.) The void, and the crack as well, grow slowly in a stable
Fig. 2. Scanning electron micrographs of fracture surfaces in Ni55Pd35Plo metal glasses tested at various conditions: (a) room temperature, (b) at 77 °K, and (c) under torsion added to tension at room temperature.
manner, and join together to allow the overall crack front to advance. The arrow in photo (f) shows the neck of the connecting part of the void through photos (d)-(e). It is noteworthy that this neck morphology is not of the "necked tearing type". Finally, the specimen fails catastrophically on a small incremental increase of stress. Accordingly, the crack becomes large enough and the applied stress high enough, to permit unstable propagation. This is schematically illustrated in Fig. 5, where CF and af are the critical crack length and failure stress respectively. These observations reveal that the fracture of an amorphous alloy proceeds quasi-statically with increasing applied stress, and that true fracture stresses of amorphous alloys could be higher than values previously reported. The paired segments after fracture are shown in Fig. 6. The edges of the fracture specimen in photos (a) and (b) appear relatively sharp.
263
Ip
Fig. 3. F r a c t u r e surface of s o m e crystalline materials: (a) r e p l i c a t e d f r a c t u r e surface in F e - 3 % Si alloys at 77 °K; (b) s c a n n i n g e l e c t r o n m i c r o g r a p h s o f t h e fract u r e surfaces at r o o m t e m p e r a t u r e in as-roiled 304 stainless steel foils recrystallized in v a c u o [ 22 ].
After sudden unstable fracture occurs, many large plastic necks are observed on the edges, as shown in (c) and (d) of Fig. 6, where photo (c) indicates the start of the unstable crack. Clearly, in this type of unstable crack propagation a large number of voids form rapidly at the crack tip and the material between the voids necks down. It is worth noting here that these plastic neck morphologies represent a two-dimensional ridge pattern on the fracture surface of the amorphous alloy. Accordingly, the vein or ridge patterns in the above scanning electron micrographs indicate a threedimension ridge pattern. One large neck is shown magnified in Fig. 7, where the left photo is the bright field image (BF) and the right is the dark fiel d image (DF). The DF of the torn pattern reveals m a n y small coherent scattering spots (~ .15 A). Figure 7 shows a smooth tearing morphology and hence a viscous-like deformation. It is interesting to compare Fig. 7 or Fig. 6(d) with the part indicated by the arrow in Fig. 4(f). This comparison indicates that even if a void is formed and coalesces with a crack (as shown in photos (d)-(e) of Fig. 4) a necking pattern such as Fig. 7 does not result. We, therefore, suggest
Fig. 4. T h e m i c r o s c o p i c o b s e r v a t i o n s w i t h increasing t i m e as a c r a c k p r o p a g a t e s in Nis5Pd35P10 m e t a l glass in t h e tensile stage in t h e TEM; (c) t h e a r r o w indicates a m i c r o c r a c k or void w h i c h f o r m s a h e a d o f t h e c r a c k (d) - (f); a r r o w s s h o w t h e same p a r t c o n n e c t i n g a void w i t h t h e crack.
264
OD
~:
U
CF
.
.
.
.
.
.
.
~u co
.
.
.
.
.
.
.
.
i
.
.
I
.
7~
Tensaon
. . . .
slow crock growth Ji...
IS/OWl
no-go.- ~ go~ :go(unstable) I
I
,
.
J
STRESS Fig. 5. Schematic illustration of the crack propagation for an amorphous alloy.
that the adiabatic heat induced by high strain rate deformation plays an important role in producing the necking pattern, i.e., a viscouslike flow. On the other hand, the relatively sharp offsets would call for a dislocation-like flow. It, therefore, appears that stable crack propagation corresponds to a dislocation-like flow while the unstable cracks involve some viscous flow. Gilman discussed the general aspects of deformation in terms of correlated molecular events [23]. His analysis suggests that there is a dislocation-flow mode and a viscous mode depending on the relaxation time for the molecular or atomic rearrangement. To estimate, roughly, the critical energy release rate Gc on initiating unstable fracture, we use the Irwin equation:
Fig. 6. The transmission electron micrographs Of the paired fracture segments; the large arrow exhibits the direction of the fracture propagation. Four small arrows from top to bottom are associated with photos (a)-(d) in the magnified scale.
G c = ?rCFOF2/E
Fig. 7. Transmission electron micrographs of the torn patterns on the fracture segments of Ni55Pd35P10 amorphous alloys; the left photo shows the BF and the right the DF.
265 w h e r e GF, E a n d CF d e n o t e f r a c t u r e stress, Y o u n g ' s m o d u l u s a n d critical c r a c k length respectively [24]. From the observation of f r a c t u r e s e g m e n t s in Fig. 6, t h e critical c r a c k length, CF, is m e a s u r e d as 25 ~m. S u b s t i t u t i n g ~F = 1 4 9 k g / m m 2 a n d E = 1.1 X 104 k g / m m 2 for Ni55PdssPlo metallic glass [11, 12] i n t o the a b o v e e q u a t i o n , Gc is c a l c u l a t e d to be 1.6 X 106 e r g / c m 2, i.e. ~ t w o o r d e r s o f magnit u d e l o w e r t h a n the d a t a p r e v i o u s l y r e p o r t e d [ 2 5 ] . Even t a k i n g a c c o u n t o f u n d e r e s t i m a t i o n o f t h e f r a c t u r e stress, this d i s c r e p a n c y is still large and, h e n c e , m a y be d u e to t h e d i f f e r e n c e s in t h i c k n e s s o f samples.
ACKNOWLEDGEMENTS This w o r k is s u p p o r t e d b y t h e office o f Naval K e s e a r c h a n d t h e N a t i o n a l Science Foundation. The authors gratefully acknowledge Dr. S. N a k a h a r a f o r c o l l a b o r a t i o n on t h e TEM a n d Dr. L. A. Davis f o r his valuable comments.
REFERENCES 1 T. Masumoto and R. Maddin, Acta Met., 19 (1971) 725. 2 H. J. Leamy, H. S. Chen and T. T. Wang, Met. Trans., 3 (1972) 699. 3 C. A. Pampillo and A. C. Reimschuessel, J. Mater.
Sci., 9 (1974) 718. 4 D. E. Polk and D. Turnbull, Acta Met., 20 (1972) 493. 5 S. Takayama and R. Maddin, Acta Met., 23 (1975) 943. 6 C. A. Pampillo, Scripta Met., 6 (1972) 915. 7 H. S. Chen, Scripta Met., 7 (1973) 931. 8 C. A. Pampillo and H. S. Chen, Mater. Sci. Eng., 13 (1974) 181. 9 S. Takayama, AIME Spring Meeting, 1973. 10 S. Takayama and R. Maddin, Cryst. Growth Intern. Conf., Tokyo, 1973. 11 S. Takayama, Ph.D. Thesis, Univ. Pennsylvania, 1974 12 S. Takayama and R. Maddin, Phil. Mag., 32 (1975) 457. 13 L. A. Davis, Scripta Met., 9 (1975) 339. 14 L. A. Davis and S. Kavesh, J. Mater. Sci., 10 (1975)453. 15 T. Masumoto and R. Maddin, Mater. Sci. Eng., 19 (1975) 1. 16 L. A. Davis, J. Mater. Sci., to be published. 17 L. A. Davis, Scripta Met., 9 (1975) 431. 18 R. Pond, Jr. and R. Maddin, Trans. AIME, 245 (1969) 2476. 19 J. J. Gilman, J. Appl. Phys., 27 (1968) 1968. 20 J. Friedel, Fracture, Wiley, New York, 1959. 21 J. R. Low, in B. Averbath, D. Felbeck, G. Hahn and D. Thomas (eds.), Fracture, Technology Press, 1959, p. 69. 22 R. W. Bauer and H. G. F. Wilsdorf, Scripta Met., 7 (1963) 1213. 23 J. J. Gilman, in A. R. Rosenfield et al. (eds.), Dislocation Dynamics, McGraw-Hill, p. 3 ; J. Appl. Phys., 44 (1973) 675. 24 J. F. Knott, Fundamentals of Fracture Mechanics, Wiley, New York, 1973, p. 117. 25 H. Kimura and T. Masumoto, Scripta Met., 9 (1975) 211.
261
© Elsevier Sequoia S.A., Lausanne -- Printed in the Netherlands
Microscopic Observations of Fracture Behavior in a NissPdssP10 Metallic Glass*
S. TAKAYAMA Materials Research Center, Allied Chemical Corporation, Morristown, NJ 07960 (U.S.A.)
R. MADDIN Department of Metallurgy and Materials Science, University of Pennsylvania, Philadelphia, PA 19174 (U.S.A.)
SUMMARY
EXPERIMENTAL
Deformation and fracture of a NissPdssP10 metallic glass have been examined in a tensile stage in a transmission electron microscope. It is observed that a plastic zone appears at the crack tip in advance of crack propagation. The microscopic fracture profile shows two characteristic features: relatively sharp offsets (dislocation-like flow) and tearing behavior (viscous-type flow).
A powder metallurgy m e t h o d was used to prepare the Ni55PdssP1o ternary alloy. Ribbon filaments were obtained by quenching the molten alloy on to the wall of a rotating metallic drum [15]. These ribbon filaments were carefully examined by X-ray and electron diffraction techniques to determine whether or not they were amorphous. For transmission electron microscopy (TEM), the filaments were thinned by an ion-milling instrument. The in situ observations of the fracture behavior were made using a tensile stage m o u n t e d within the electron microscope.
INTRODUCTION Following the early studies of the mechanical properties of PdsoSi20 by Masumoto and Maddin [1] there have now been numerous reports on various amorphous alloys [2 - 17]. It has been shown that amorphous alloys deform in a localized manner, i.e., inhomogeneously, and that their tensile fracture surfaces are usually characterized by two distinct regions [1], a relatively featureless part [2] (shear offset) and a "vein" or "ridge" pattern. The formation of this latter pattern has been interpreted in terms of an adiabatic shear [2]. We present, herein, electron microscopic observations of the tensile mechanics of a Ni55PdssP1o metallic glass in order to contribute additional details to the understanding of the deformation and failure behavior.
*Paper presented at the Second International Conference on Rapidly Quenched Metals, held at the Massachusetts Institute of Technology, Cambridge, Mass., November 17 - 19, 1975.
PROCEDURE
RESULTS AND DISCUSSION Typical stress-strain curves for Ni55PdssPlo are presented in Fig. 1, (a) and (b), for room and liquid nitrogen temperatures, respectively. These curves show the same characteristics as other amorphous alloys, i.e., they show no well-defined yield point, but rather a small deviation from linearity before fracture. Some fracture surfaces are shown in Fig. 2 for a nominal strain rate of 5 × 10 -4 sec -1. The shear offset and the size of the vein cells decrease as the temperature decreases. Figure 2(c) shows that the vein morphology is curved, going from b o t t o m left to top right of the photos, i.e. along the applied torsion direction. A fracture surface morphology similar to amorphous alloys is often observed in crystalline materials. The processes producing the observed morphology have been explained via a dislocation model [19, 20]. Thus, Fig. 3 shows typical fracture surfaces in some crystalline materials ((a) shows a typical river pattern for quasi-brittle crystalline material [21],
262
:5xlO-4sec q
200 (a)
(b)
i
1 50
'E 1 O0
E
m 50
I
I I 1.0 STRAIN (,~L/L x 10 "2 )
Fig. 1. Stress-strain curves of Ni55Pd~5P10 metallic glasses, testing strain rate e = 5 × 10- sec-1.
while (b) shows ripple pattern for ductile failure [22] ). It is noteworthy that, compared with amorphous fracture surfaces, the fracture surface morphology of a quasi-brittle material is similar to a ductile glass (compare Fig. 2(a) with Fig. 3(a)), while that of a ductile crystalline material can look like that of a brittle amorphous one. Examination of the deformation of metallic glasses reveals that a small number of localized deformation bands appear prior to fracture. Moreover, these deformation bands are often observed near the fracture surface. It, therefore, appears that localized plastic flow contributed to produce those fracture patterns. Figure 4, (a)-(h), shows the crack propagat-ion sequence when a thin specimen of Ni55PdasP10 amorphous alloy was pulled to failure on the deformation stage of a transmission electron microscope. The crack propagates in order from (a) to (h) in Fig. 4 with increasing applied tensile stress. Thus, initially, crack growth proceeds slowly in a stable manner with increasing stress after the plastic zones are formed ahead of the crack (as shown in (a)-(c)). In (d) another void appears ahead of the crack. (The arrow indicates the part connecting the void with the crack.) The void, and the crack as well, grow slowly in a stable
Fig. 2. Scanning electron micrographs of fracture surfaces in Ni55Pd35Plo metal glasses tested at various conditions: (a) room temperature, (b) at 77 °K, and (c) under torsion added to tension at room temperature.
manner, and join together to allow the overall crack front to advance. The arrow in photo (f) shows the neck of the connecting part of the void through photos (d)-(e). It is noteworthy that this neck morphology is not of the "necked tearing type". Finally, the specimen fails catastrophically on a small incremental increase of stress. Accordingly, the crack becomes large enough and the applied stress high enough, to permit unstable propagation. This is schematically illustrated in Fig. 5, where CF and af are the critical crack length and failure stress respectively. These observations reveal that the fracture of an amorphous alloy proceeds quasi-statically with increasing applied stress, and that true fracture stresses of amorphous alloys could be higher than values previously reported. The paired segments after fracture are shown in Fig. 6. The edges of the fracture specimen in photos (a) and (b) appear relatively sharp.
263
Ip
Fig. 3. F r a c t u r e surface of s o m e crystalline materials: (a) r e p l i c a t e d f r a c t u r e surface in F e - 3 % Si alloys at 77 °K; (b) s c a n n i n g e l e c t r o n m i c r o g r a p h s o f t h e fract u r e surfaces at r o o m t e m p e r a t u r e in as-roiled 304 stainless steel foils recrystallized in v a c u o [ 22 ].
After sudden unstable fracture occurs, many large plastic necks are observed on the edges, as shown in (c) and (d) of Fig. 6, where photo (c) indicates the start of the unstable crack. Clearly, in this type of unstable crack propagation a large number of voids form rapidly at the crack tip and the material between the voids necks down. It is worth noting here that these plastic neck morphologies represent a two-dimensional ridge pattern on the fracture surface of the amorphous alloy. Accordingly, the vein or ridge patterns in the above scanning electron micrographs indicate a threedimension ridge pattern. One large neck is shown magnified in Fig. 7, where the left photo is the bright field image (BF) and the right is the dark fiel d image (DF). The DF of the torn pattern reveals m a n y small coherent scattering spots (~ .15 A). Figure 7 shows a smooth tearing morphology and hence a viscous-like deformation. It is interesting to compare Fig. 7 or Fig. 6(d) with the part indicated by the arrow in Fig. 4(f). This comparison indicates that even if a void is formed and coalesces with a crack (as shown in photos (d)-(e) of Fig. 4) a necking pattern such as Fig. 7 does not result. We, therefore, suggest
Fig. 4. T h e m i c r o s c o p i c o b s e r v a t i o n s w i t h increasing t i m e as a c r a c k p r o p a g a t e s in Nis5Pd35P10 m e t a l glass in t h e tensile stage in t h e TEM; (c) t h e a r r o w indicates a m i c r o c r a c k or void w h i c h f o r m s a h e a d o f t h e c r a c k (d) - (f); a r r o w s s h o w t h e same p a r t c o n n e c t i n g a void w i t h t h e crack.
264
OD
~:
U
CF
.
.
.
.
.
.
.
~u co
.
.
.
.
.
.
.
.
i
.
.
I
.
7~
Tensaon
. . . .
slow crock growth Ji...
IS/OWl
no-go.- ~ go~ :go(unstable) I
I
,
.
J
STRESS Fig. 5. Schematic illustration of the crack propagation for an amorphous alloy.
that the adiabatic heat induced by high strain rate deformation plays an important role in producing the necking pattern, i.e., a viscouslike flow. On the other hand, the relatively sharp offsets would call for a dislocation-like flow. It, therefore, appears that stable crack propagation corresponds to a dislocation-like flow while the unstable cracks involve some viscous flow. Gilman discussed the general aspects of deformation in terms of correlated molecular events [23]. His analysis suggests that there is a dislocation-flow mode and a viscous mode depending on the relaxation time for the molecular or atomic rearrangement. To estimate, roughly, the critical energy release rate Gc on initiating unstable fracture, we use the Irwin equation:
Fig. 6. The transmission electron micrographs Of the paired fracture segments; the large arrow exhibits the direction of the fracture propagation. Four small arrows from top to bottom are associated with photos (a)-(d) in the magnified scale.
G c = ?rCFOF2/E
Fig. 7. Transmission electron micrographs of the torn patterns on the fracture segments of Ni55Pd35P10 amorphous alloys; the left photo shows the BF and the right the DF.
265 w h e r e GF, E a n d CF d e n o t e f r a c t u r e stress, Y o u n g ' s m o d u l u s a n d critical c r a c k length respectively [24]. From the observation of f r a c t u r e s e g m e n t s in Fig. 6, t h e critical c r a c k length, CF, is m e a s u r e d as 25 ~m. S u b s t i t u t i n g ~F = 1 4 9 k g / m m 2 a n d E = 1.1 X 104 k g / m m 2 for Ni55PdssPlo metallic glass [11, 12] i n t o the a b o v e e q u a t i o n , Gc is c a l c u l a t e d to be 1.6 X 106 e r g / c m 2, i.e. ~ t w o o r d e r s o f magnit u d e l o w e r t h a n the d a t a p r e v i o u s l y r e p o r t e d [ 2 5 ] . Even t a k i n g a c c o u n t o f u n d e r e s t i m a t i o n o f t h e f r a c t u r e stress, this d i s c r e p a n c y is still large and, h e n c e , m a y be d u e to t h e d i f f e r e n c e s in t h i c k n e s s o f samples.
ACKNOWLEDGEMENTS This w o r k is s u p p o r t e d b y t h e office o f Naval K e s e a r c h a n d t h e N a t i o n a l Science Foundation. The authors gratefully acknowledge Dr. S. N a k a h a r a f o r c o l l a b o r a t i o n on t h e TEM a n d Dr. L. A. Davis f o r his valuable comments.
REFERENCES 1 T. Masumoto and R. Maddin, Acta Met., 19 (1971) 725. 2 H. J. Leamy, H. S. Chen and T. T. Wang, Met. Trans., 3 (1972) 699. 3 C. A. Pampillo and A. C. Reimschuessel, J. Mater.
Sci., 9 (1974) 718. 4 D. E. Polk and D. Turnbull, Acta Met., 20 (1972) 493. 5 S. Takayama and R. Maddin, Acta Met., 23 (1975) 943. 6 C. A. Pampillo, Scripta Met., 6 (1972) 915. 7 H. S. Chen, Scripta Met., 7 (1973) 931. 8 C. A. Pampillo and H. S. Chen, Mater. Sci. Eng., 13 (1974) 181. 9 S. Takayama, AIME Spring Meeting, 1973. 10 S. Takayama and R. Maddin, Cryst. Growth Intern. Conf., Tokyo, 1973. 11 S. Takayama, Ph.D. Thesis, Univ. Pennsylvania, 1974 12 S. Takayama and R. Maddin, Phil. Mag., 32 (1975) 457. 13 L. A. Davis, Scripta Met., 9 (1975) 339. 14 L. A. Davis and S. Kavesh, J. Mater. Sci., 10 (1975)453. 15 T. Masumoto and R. Maddin, Mater. Sci. Eng., 19 (1975) 1. 16 L. A. Davis, J. Mater. Sci., to be published. 17 L. A. Davis, Scripta Met., 9 (1975) 431. 18 R. Pond, Jr. and R. Maddin, Trans. AIME, 245 (1969) 2476. 19 J. J. Gilman, J. Appl. Phys., 27 (1968) 1968. 20 J. Friedel, Fracture, Wiley, New York, 1959. 21 J. R. Low, in B. Averbath, D. Felbeck, G. Hahn and D. Thomas (eds.), Fracture, Technology Press, 1959, p. 69. 22 R. W. Bauer and H. G. F. Wilsdorf, Scripta Met., 7 (1963) 1213. 23 J. J. Gilman, in A. R. Rosenfield et al. (eds.), Dislocation Dynamics, McGraw-Hill, p. 3 ; J. Appl. Phys., 44 (1973) 675. 24 J. F. Knott, Fundamentals of Fracture Mechanics, Wiley, New York, 1973, p. 117. 25 H. Kimura and T. Masumoto, Scripta Met., 9 (1975) 211.