Fatigue-Strength Enhancement of Cast Zr 50 Cu 40 Al 10 Glassy Alloys

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1 Materials Transactions, Vol. 47, No. 5 (26) pp to 1293 #26 The Japan Institute of Metals -Strength Enhancement of Cast Cu 4 Glassy Alloys Yoshihiko Yokoyama 1, Peter K. Liaw 2, Masahiko Nishijima 3, Kenji Hiraga 3, Raymond A. Buchanan 2 and Akihisa Inoue 1 1 Advanced Research Center of Metallic Glasses, Institute for Materials Research, Tohoku University, Sendai , Japan 2 Faculty of Engineering, The University of Tennessee, 434 Dougherty Engineering Building, Knoxville, Tennessee , USA 3 Nanotechnology Support Project of the Ministry of Education, Institute for Materials Research, Tohoku University, Sendai , Japan In order to improve the fatigue strength of Cu 4 BGAs during cycles, we tried to use the small additive element. As a result, Pd, Ag, Pt, and Au are effective addition elements to enhance the fatigue strength. Especially, the additive Pd element has a beneficial effect on the fatigue-strength enhancement. The additive Pd element promotes the glass-structure expansion, which can be recognized from the volume change due to the structural relaxation. Consequently, we found the linear relationship between the fatigue limit and volume change in Cu 4 X [X: 7 atomic percent (at%)] glassy alloys. Furthermore, we conclude that the origin of the unique hardness distribution on the fatigue-fractured surface is probably related to hydrogen hardening. The Cu 37 Pd 3 glassy alloy, which exhibits the highest fatigue limit of 1,5 MPa, shows a superior resistance force against the hydrogen hardening. Thus, maintaining the balance between the hardening and embrittlement around the fatigue crack tip by hydrogen is the significant factor to enhance the fatigue strengths. [doi:1.232/matertrans ] (Received November 22, 25; Accepted March 7, 26; Published May 15, 26) Keywords: cast Cu 4 X (X: 7 at%) glassy alloys, volume change, Wöhler curves, fatigue limits, hydrogen effect 1. Introduction features of bulk-glassy alloys are important factors for the industrial applications of these materials. In order to obtain a reliability, fatigue testing for glassy alloys is an important study, whereas the fatigue strength is not usually so high with respect to its static high strength. 1) Furthermore, the Zr-Be based glassy alloy exhibits a loading rate dependence 2) of fracture toughness due to the viscoelastic feature. Unlike a glassy polymer, since the Zr-based glassy alloy exhibit a high glass-transition temperature, it is thought that the loadingrate dependability of the fracture-toughness value resulting from a viscoelasticity does not become a fatal drawback at room temperature. This is, the mechanical feature of glassy alloys is one of special interests, because the atomic bonding is a ductile metallic bonding but the structure has no periodicity. The mechanical features of crystalline plastic alloys can provide a superior plasticity by systematically moving dislocations. There are no operable dislocations in glassy alloys having an aperiodic structure. The deformation mechanism of glassy alloys is characterized by the unique adiabatic shear-band operation. 3) The existence of vein patterns on the fractured surface of ductile glassy alloy implies the pseudo-melting state in the adiabatic shear band. Therefore, one-slip movement brings about a final fracture with little uniform plastic deformation. The lack of the uniform deformability of glassy alloys is simply considered as limiting the toughness. 4) However, the glassy material, whose high strength (about 2 GPa) and high toughness (about 5 MPam :5 ) can be realized only by casting, never seen in ordinary crystalline material. Cast Zr-based bulk glassy alloys (BGAs) are characterized by these superior mechanical properties. Bulk-shaped glassy alloys are promoted by those superior glass-forming abilities, 5) which enable the production of bulk-shape samples by using a conventional metallic mold cast method. Bulk-shaped glassy alloys yield new application areas as specified structural materials 6,7) and also have new problems, 8,9) such as cast defects in bulk glassy alloys. The most important problem is the existence of crystalline inclusions, which act as crack-initiation sites and enhance crack propagation during the fatigue test. Therefore, the mechanical properties of cast BGAs sometimes depend on the quality of the cast structure which can be controlled by the cast method. 1) In this paper, we used an arc-cast furnace with a tilt cast mechanism 1) to prepare crystalline inclusions free BGA specimens. Since glassy alloys exhibit little plastic tensile elongation and no work-hardening phenomena, the fatigue features of glassy alloys are different from those of ordinary crystalline materials. However, it was reported 11) that some glassy alloys show distinct knees in the Wöhler curves of Pd-, Ni-, and Zrbased glassy alloys. Furthermore, some fatigue-fractured surfaces demonstrate striation-like marks, 12) which might result from the fatigue-crack-tip branching due to the yielding around the crack tip. The existence of the knee implies the strain-aging effect 13) at the fatigue crack tip, whereas glassy alloys usually do not show the work-hardening feature after the yielding. Especially, this fact suggests that the difference of the fatigue-crack-propagation mechanisms between the ordinary crystalline and glassy alloys. This paper will present the Wöhler curves of Cu 4 X - (X: 7 at%) BGAs, and discuss the effective factor, which influences the fatigue limit. We will also attempt to clarify the fatigue-crack-propagation mechanism in these BGAs with the high fatigue strength. 2. Experimental Procedure In this study, we examined ternary and quaternary Cu 4 X (X: 7 at%) and quaternary Zr 97 X Y - Cu X Al Y Pd 3 (X ¼ 27{42 at%, Y ¼ 7:5{12:5 at%) glassy al-

2 -Strength Enhancement of Cast Cu 4 Glassy Alloys 1287 loys. The master alloy ingots were prepared by arc melting mixtures of pure Zr, Cu, Al, and Pd metals in an argon atmosphere. To maintain the low oxygen concentration of the master alloys, we used a special Zr crystal rod (< :5 at% oxygen). The oxygen concentration of the bulk glassy alloys was measured using a fusion-in-helium gas-infrared-absorption method (give a reference). We prepared the cast rod sample (8 6 mm) with the tilt cast technique, 14) which is effective in eliminating crystalline inclusions in BMGs. The rod-shaped cast samples were converted into sandglassshaped fatigue specimens using a diamond-gliding machine to avoid the surface deterioration. We used rotating-beam fatigue testing machines to examine the Wöhler curves of BGAs. The frequency of the applied cyclic stress in this study is 5 Hz. We examined the cast structure and fatiguefractured surface by the optical microscopy (OM) and scanning-electron microscopy (SEM) and measured the hardness on the free surface and fatigue-fractured surface by a micro Vickers testing machine by applying a 2.9 N load for 15 s. We also observe the cast structure of Cu 4 X - (X: 7 at%) glassy alloys before and after the fatigue testing by high-resolution transmission electron microscope (HRTEM) JEOL 21 F. After the fatigue tests, the TEM specimen were cut from just under the fatigue fractured surface and then mechanical polishing were performed to obtain thin thickness less than 3 mm. TEM specimens were prepared by electro chemical etching using nitric acid ethanol at about 24 K. Hydrogen was charged by an electrolytic process with the BMG disk cathode in a 1 N sulfuric-acid solution with a constant current density of 1. ka/m 2. The hydrogen concentration of the hydrogencharged glassy alloys was also measured using a fusion-inargon gas-thermal conductivity method. 3. Results and Discussions The Wöhler curves of Cu 4 and Cu 3 Ni 1 glassy alloys are shown in Fig. 1. It reveals that the Zr-TM- Al BGAs shows the same tendency for the fatigue-strength degradation during cycles. Details about the abatement of the fatigue strength are given in Ref 13). In this paper, the consideration about an origin and measurement of the fatigue-strength degradation are described. The toughness of glassy alloys is influenced by the structural relaxation, which can be estimated by the volume change with the heat Amplitude Stress, σ a / MPa Cu 3Ni 1 Cu 4 treatment below the glass-transition temperature (T g ). By assuming that the glass structure is simply composed of clusters and voids, the volume change is caused by the disappearance of some voids in the glass structure. Since the glassy alloy maintains an amorphous structure after the structural relaxation, disappeared voids are corresponding to the frozen in the free volume during quenching, called as the excessive free volume. Therefore, we can estimate the difference of the excessive free volume in glassy alloys by the volume change due to structural relaxation. An excessive free volume actually means the crevice between atoms in the glass structure, and an atom is made easy to move and it is considered to raise the stress-relaxation ability. Since a fatigue phenomenon is a local yielding by the cyclic applied load, it is expected that the high stress relaxation ability is probably effective in the improvement of the fatigue strength. Therefore, in this research, a Pd element was added to the basic Cu 4 glassy alloy to control the volume change. 15) X-ray diffraction patterns of Cu 4 X (X: 7 at%) glassy alloys are shown in Fig. 2. A distinct Bragg peak, which is corresponding to the existence of second crystalline phases, was not seen in X-ray spectra of the tiltcast Cu 4 X (X: 7 at%) glassy alloys. In order to clarify the nano-structure of the tilt-cast Cu 4 X - (X: 7 at%) glassy alloys, the high-resolution TEM (HRTEM) observation was also performed. The HRTEM specimens were cat off the behind of the fatigue-fractured surface. A few nano-meters crystallized regions, whose morphology seems to be like a network structure, were observed in the tilt-cast Cu 37 Pd 3 glassy alloy as presented in Fig. 3. However, there is no large distinct crystalline particle over 1 nano-meters in HRTEM images. Besides, no distinct crystallized regions were seen in HRTEM images of Cu 37 Pd 3 glassy alloys before fatigue test. Static mechanical properties of the tilt-cast Cu 4 X (X: 7 at%) glassy alloys were examined as shown in Fig. 4. The tensile strength and Young s Intensity, I (a.u.) Cu-K α X=at% X=1at% X=2at% X=3at% X=5at% X=7at% Cycle Number to Fracture, Nf θ Fig. 1 Wöhler curves of tilt cast Cu 3 and Cu 3 Ni 1 BGAs. Fig. 2 X ray diffraction patterns of tilt cast Cu 4 X (X: 7 at%) BGAs.

3 1288 Y. Yokoyama et al. 2 nm Fig. 3 High-resolution transmission electron micrograph image of tilt cast Cu 37 Pd 3 BGA after fatigue test ( a ¼ 12 MPa, N f ¼ 1:8 1 4 ). White lines point out the nano-crystallized region with lattice fringes. Tensile Strength, σ B / MPa Cu 4 X Pd content (at%) Fig. 4 Vickers hardness, Young s modulus and tensile strength of tilt cast Cu 4 X (X: 7 at%) BGAs. modulus exhibit linearly increase with the Pd concentration, whereas the Vickers hardness takes a minimum value at Pd ¼ 3 at%. In general, tensile strength has close linier relationship with hardness in ordinary engineering alloys; Young s Modulus, E / GPa Volume Change (%) X(at%) Fig. 5 Volume change from as cast state to fully relaxed state of tilt cast Cu 4 X (X: 7 at%) BGAs. however, Zr-Cu-Al-Pd BMGs does not show the same tendency. Although the reason for softening at Pd ¼ 3 at% is not clear, it is suggested that there might be a significant difference in the glass structure. In order to estimate the glass structural difference, we use the volume change by the structural relaxation. Because the determination of the absolute value of the free volume is difficult in Zr-based BMGs, whereas the volume change by the structural relaxation is related to the excessive free volume. Figure 5 shows the relationship between the volume change and Pd

4 -Strength Enhancement of Cast Cu 4 Glassy Alloys 1289 Stress Range, σ / MPa X = X = 1 X = 2 X = 3 X = 5 X = Limit, σ W /MPa Zr-Cu-Al-Pd system (as cast state) r 2 = Volume Change (%) 1 Fig Cycle Number to Fracture, N f Wöhler curves of tilt cast Cu 4 X (X: 7 at%) BGAs. amo amo+cry cry.46~.41%.4~.35%.34~.29%.28~.23%.22~.17%.16~.11% No Composition Zr5Cu37Al1Pd3 Zr5Cu35Al1Pd5 Zr5Cu38Al1Pd2 Zr5Cu39Al1Pd1 Zr5Cu34.5Al12.5Pd3 Zr47.5Cu39.5Al1Pd3 Volume Change (%) σ L(MPa) No Composition Zr5Cu33Al1Pd7 Zr5Cu39.5Al7.5Pd3 Zr52.5Cu37Al7.5Pd3 Zr52.5Cu34.5Al1Pd3 Zr55Cu32Al1Pd3 Zr47.5Cu37Al12.5Pd3 Volume Change (%) σ L(MPa) Fig. 8 Relationship between the volume change and fatigue limit of tilt cast Cu 4 X (X: 7 at%) and Zr 97 X Y Cu X Al Y Pd 3 (X: at%, Y: at%) BGAs. r is the correlation factor Al (at%) Cu (at%) Fig. 7 Compositional dependence of volume change of tilt cast Zr 97 X Y Cu X Al Y Pd 3 (X: at%, Y: at%) BGAs. concentration of Cu 4 X (X: 7 at%) glassy alloys. The volume change by the structural relaxation takes a maximum value at Pd ¼ 3 at%, whose alloy composition shows a softening phenomenon as shown in Fig. 4. The larger value of the volume change means the higher volume expansion effect of the amorphous structure. Therefore, the volume change indicates the existence of thermally unstable volume below T g in the amorphous structure, when we assume that the amorphous structure is composed of clusters and voids. Thermally unstable voids are probably effective to relax the localized stress, because the atoms surrounded the void move easier than others. The superior ability of the stress relaxation has advantages on fracture and fatigue properties. Figure 6 shows the Wöhler curves of - Cu 4 X (X: 7 at%) glassy alloys. Pd ¼ at% BGA shows the lowest fatigue limit about 25 MPa. The fatigue limit increases with the Pd content and shows the maximum value at Pd ¼ 3 at% as 1,5 MPa, and then the fatigue limit decreases with the Pd content. The trend of the fatigue limit vs. the Pd content is similar to that of the volume change vs. Pd content, as shown in Fig. 5. Therefore, we tried to determine the optimum alloy composition with the largest value of the volume change. Figure 7 shows the compositional dependence of the volume change in Zr 97 X Y - Cu X Al Y Pd 3 (X: at%, Y: at%) glassy alloys. The phase characterization of the tilt-cast alloys in this figure was performed by the X-ray diffractometry. This figure points out that the glassy Cu 37 Pd 3 alloy, whose composition is considered as the pseudo-ternary eutectic composition, exhibits the maximum volume change about.5%. Accordingly, we summarize the relationship between the fatigue limit and volume change of Cu 4 X (X: 7 at%) and Zr 97 X Y Cu X Al Y Pd 3 (X: at%, Y: at%) glassy alloys in Fig. 8. All of the tilt-cast Zr-Cu- Al-Pd glassy alloys were studied by the fatigue-fractured surface observation to select the high-quality glassy samples, whose fatigue-fractured surface exhibits no crystalline inclusions. As the result, the volume change and fatigue limit shows a linier relationship. This fact means the glass structural control with increasing the free volume, which can be roughly estimated by the volume change, is important to obtain the ductile structural glassy alloy with high fatigue limit. The fractured surface of the glassy sample after fatigue tests shows thumb nail shape fatigue-fractured regions as shown in Figs. 9(a) and with SEM observation. In Fig. 9, a fatigue crack initiated from the surface of the sample with shear slips, and the fatigue crack propagation region shows striation-like marks and the final fracture with vein patterns starts from the interface like a thumb-nail shape. Striations-like marks are usually composed of shear bands. Therefore, the period of striation-like marks is not related to the stress-intensity factor, whereas the morphology of striation-like marks change near the final fracture to exhibit large striation marks whose period is related to the stressintensity factor. Figure 1 shows the schematic illustration of the fatigue crack propagation (a) and striation-like image with shear-band and striation marks. By using the striation marks on the fatigue-fractured surface as shown in Fig. 1, we tried to measure the crack-propagation rate. Figure 11 shows the crack-propagation rates of Cu 4 X - (X: 7 at%) glassy alloys. These fatigue-crackpropagation rates were fit to Paris s equation: ðda=dnþ ¼ AK m, where A and m are constant. Especially, the value of

5 129 Y. Yokoyama et al. (a) 1 µ m µ m 1 µ m µ m 1 µ m Fig. 9 -fractured surface of tilt cast Cu 37 Pd 3 BGA with applied cyclic stress amplitude of 12 MPa and lifetime of 1:8 1 4 cycles. (a) Slip plane Crack Propagation 5 µ m Fig. 1 Schematic illustration of fatigue crack propagation with amount of shear bands operation (a) and actual fatigue-fractured surface with large striation-like mark and small striation-like mark. The former is originated to the blanching and the latter is caused by the shear band. m has a close relationship with the work-hardening coefficient on the cyclic stress strain curve. 16) Assuming that the origin of the fatigue-limit increase is mechanical features change of Cu 37 Pd 3 glassy alloy, the value of m should change by the Pd concentration. However, there is no significant difference in the value of m as shown in Fig. 11. This fact means that the fatigue-limit enhancement is probably caused by other factors like an environment. We have already reported 17) micro Vickers hardness changes on the fatigue-fractured surface of Zr-Cu-Al and Zr-Cu-Ni-Al BGAs. The hardness takes the maximum value at the interface between the fatigue and final fractured regions. Therefore, the well-grown fatigue crack propagation accompanied with the hardening mechanism around the fatigue crack tip to enhance the fatigue-crack-propagation stoppage. The hardening mechanism is a significant problem

6 -Strength Enhancement of Cast Cu 4 Glassy Alloys 1291 Crack Growth Rate, da/dn / µm cycle Cu 4-X m = 3.76 m = 3.81 m = 3.45 m = 3.58 m = 3.61 m = 3.68 X = X = 1 X = 2 X = 3 X = 5 X = Stress Intensity Factor Range, K / MPa m.5 Fig. 11 Relationship between crack fatigue propagation rate (da=dn) and stress intensity factor (K) of tilt cast Cu 4 X (X: 7 at%) BGAs. The m value means the exponent factor of Paris s equation. to control the fatigue features of BGAs. In case of Pd 4 Cu 3 Ni 1 P 2 BGAs, 18) they also show the hardness change and the fatigue-crack-initiation site exhibits the maximum value of the hardness distribution, because no fatigue-crack-propagation marks were seen on the fatiguefractured surface. Since the BGAs do not show the workhardening phenomenon, which is usually considered as an origin of the strain aging, the hardening mechanism of BGAs is significant factor to control the fatigue features in BGAs. Figure 12 shows micro-vickers hardness change on the fatigue-fractured surface of Cu 4 X (X: 7 at%) glassy alloys. Cu 4 BGA shows a significant increase of the hardness value around the interface. With increasing the additive Pd concentration, the maximum value of hardness at the interface decreases, and then, there is no distinct hardness change on the fatigue-fractured surface of Cu 37 Pd 3 BGA. With an additive Pd concentration over 3 at%, the hardness value takes the maximum value at the interface. Consequently, the Cu 37 Pd 3 BGA has a unique mechanism to restrict the hardness increase around the fatigue-crack tip during the fatigue testing. In order to clarify the origin of the hardness increase, we observed the fatigue-fractured surface by the TEM method; however, we could not find any crystallized region leading to the hardness increase. The structural relaxation is also an insufficient phenomenon to explain the hardness increase on the fatiguefractured surface in these Zr-based BMGs. 19) Consequently, there is not distinct phase/structural change by fatigue testing. Except for the structural change to increase hardness, hydrogen is the only element to increase the hardness and keep the glass structure. Figure 13 shows the relationship between the fatigue-fractured toughness and the maximum value of the micro Vickers hardness on the fatigue-fractured surface of Cu 4 X (X: 7 at%) BGAs. In this figure, a broken line shows the relationship between the fracture toughness (K IC ) and Vickers hardness of electronically hydrogen-charged Zr-based BGAs. 2) Both relationships are in good agreement. This fact means that the hardness increase at the fatigue-crack tip is probably caused by the hydrogen concentration. The hydrogen concentration at the fatigue-crack tip is probably caused by the hardness increase to enhance the fatigue-crack-propagation stoppage, and, then, more concentrated hydrogen will cause the embrittlement around the fatigue-crack tip to decrease the K IC value. Assuming that the hydrogen effect as described above is true, we can recognize the significant cycle dependence of the fatigue-fractured toughness and the opposite phenomenon of fatigue lifetime decrease in vacuum 21) of Zr-based BMGs. In other words, glassy alloys have an advantage of the hydrogen effect on fatigue features. Because engineering crystalline alloys and intermetallic compounds have fatal drawbacks regarding hydrogen effect on fatigue features. In order to consider the fatigue-strength enhancement of Zr-based BMGs, hydrogen is a necessary element to stop the fatigue-crack propagation like a strainaging effect. Figure 14 shows the relationships between the Vickers hardness and hydrogen-charging time of - Cu 4 X (X: 7 at%) glassy alloys. This figure also points out the superior advantage of the resistance ability against the hydrogen hardening of Cu 37 Pd 3 BGA. This fact means that the hydrogen hardening is not restricted by the simple chemical effect of the Pd element but a unique glass structure of Cu 37 Pd 3 BGA. Because the - Cu 37 Pd 3 BGA has the largest volume of Cu 4 X - (X: 7 at%) glassy alloys, that promotes a large Poisson s ratio. Since the deformation of the glass structure is characterized by the localized group atoms movement in shear bands, the yielding criterion is defined by the strain at about 2%. This criterion is quite different from engineering crystalline materials, which can be estimated by the stress condition to start the dislocation movement. Therefore, the ductility of the crystalline materials is directly promoted by the easiness of dislocation movements. The dislocation is the line defect with the shear (screw dislocation) and tension/ compression (edge dislocation) stress field. Since the hydrogen usually is located at he dislocation to disturb the dislocation movement, crystalline materials exhibit the high embrittlement susceptibility on the hydrogen embrittlement. On the other hand, in the case of glassy materials, the homogeneity of the structure and excessive free volume can reduce the sensitivity to hydrogen embrittlement. We also checked the hydrogen concentration of 15 h hydrogencharged Cu 4 X (X: 7 at%) glassy alloys, Cu 37 Pd 3 BGA shows the lowest value about 12 ppm, and Cu 4 BGA shows the highest value about 1, ppm. 4. Summary We examined fatigue properties of Cu 4 X (X: 7 at%) BGAs with rotating-beam fatigue testing machines. In order to clarify the origin of the high fatigue strength of Cu 37 Pd 3 BGA, the micro Vickers hardness change on the fatigue-fractured surface and the hardness change by hydrogen charging were examined. The results obtained are summarized below. (1) The fatigue limits of Zr-Cu-Al-Pd BGAs show a good linear relationship with the volume change, which is probably corresponding to the excessive free volume. (2) The origin of the hardness increase on the fatiguefractured surface might be considered as hydrogen

7 1292 Y. Yokoyama et al. X= (a) Interfac e (σ a =357MPa N f = ) X=1 µ m µ m Distance from the Inter face, x / mm ( σ a =12MPa N f = ) Ditance from the, x/mm X=2 µ m ( σ a =15MPa N f = ) Ditance from the, x/mm X=3 µ m (σ a =12MPa N f = ) Distance from the, x/mm X= (σ a =12MPa N f = ) 1 µ m Distance from the, x/mm (k) X=7 (l) (σ a =15MPa N f = ) µ m Distance from the, x/mm Fig. 12 SEM images of fatigue-fractured surface of Cu 4 (a), Cu 39 Pd 1, Cu 38 Pd 2, Cu 37 Pd 3, Cu 35 Pd 5, and Cu 33 Pd 7 (k), and micro Vickers hardness distribution on fatigue-fractured surface of Cu 4, Cu 39 Pd 1, Cu 38 Pd 2, Cu 37 Pd 3, Cu 35 Pd 5, and Cu 33 Pd 7 (l). (x ¼ ) means the boundary between the fatigue and final fracture.

8 -Strength Enhancement of Cast Cu 4 Glassy Alloys 1293 Fracture Toughness, Kq / MPa m r 2 =.66 X = X = 1 X = 2 X = 3 X = 5 X = 7 Stanford (K1C)* Vickers Hardness on Fracture Surface, HV max Hydrogen Charging Time, t / ks X = X = 1 X = 2 X = 3 X = 5 X = 7 Fig. 13 Relationship between the fatigue-fractured toughness and maximum value of Vickers hardness on fatigue-fractured surface of tilt cast Cu 4 X (X: 7 at%) BGAs. Doted line is the data of hydrogen charging embrittlement. r is the correlation factor. hardening, which is probably considered as the substitution of strain aging effect. (3) The Cu 37 Pd 3 BGA, whose fatigue limit exhibits 15 MPa, shows no distinct hardness change on the fatigue-fractured surface. (4) The high fatigue strength of Cu 37 Pd 3 BGA probably results from the restriction feature of hardness increasing with the hydrogen charging. Acknowledgement This research is funded by NEDO [New Energy and Industrial Technology Development Organization] and authors are grateful to the Nanotechnology Support Project of the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan for technical support of this research. REFERENCES 1) A. Inoue: Bulk Amorphous Alloys (Trans Tech Publications Zurich 1998) ) R. O. Ritchie, C. J. Gilbert, V. Schroeder: Proc. Inter. Conf. on Advanced Technology in Experimental Mechanics 99, JSME- MMD, Vol. 1 (1999) ) H. S. Chen and T. T. Wang: J. Appl. Phys. 41 (197) ) C. C. Hays, C. P. Kim and W. L. Johnson: Phys. Rev. Lett. 84 (2) Fig. 14 Vickers hardness change with cathodic hydrogen charging of tilt cast Cu 4 X (X: 7 at%) BGAs ) A. Inoue: Bulk Amorphous Alloys (Trans Tech Publications Zurich 1998) ) H. Kakiuchi, A. Inoue, M. Onuki, Y. Takano and T. Yamaguchi: Mater. Trans. 42 (21) ) W. L. Johnson: JOM 3 (22) ) Y. Yokoyama, T. Shinohara, K. Fukaura and A. Inoue: Mater. Trans. 45 (24) ) Y. Yokoyama, A. Kobayashi, K. Fukaura and A. Inoue: Mater. Trans. 43 (22) ) Y. Yokoyama, K. Fukaura and A. Inoue: Intermetallics 1 (22) ) Y. Yokoyama, N. Nishiyama, K. Fukaura, H. Sunada and A. Inoue: Mater. Trans. 4 (1999) ) K. Fujita, A. Inoue and T. Zhang: Mater. Trans. JIM 41 (2) ) Y. Yokoyama, K. Fukaura and H. Sunada: Mater. Trans. 41 (2) ) Y. Yokoyama, K. Inoue and K. Fukaura: Mater. Trans. 43 (22) ) Y. Yokoyama, M. Nishijima, K. Hiraga, P. K. Liaw and A. Inoue: J. Metastable and Nanocrystalline Materials 24 (25) ) J. P. Hickerson and R. W. Hertzberg: Met. Trans. 3 (1972) ) Y. Yokoyama, K. Fukaura and A. Inoue: Mater. Trans. 45 (24) ) Y. Yokoyama, N. Nishiyama, K. Fukaura, H. Sunada and A. Inoue: Mater. Trans. 42 (21) ) Y. Yokoyama, Y. Akeno, T. Yamasaki, P. K. Liaw, R. A. Buchanan and A. Inoue: Mater. Trans. 46 (25) ) D. Suh and R. H. Dauskardt: Mater. Trans. 42 (21) ) W. H. Peter, P. K. Liaw, R. A. Buchanan, C. T. Liu, C. R. Brooks, J. A. Horton, Jr., C. A. Carmichael, Jr. and J. L. Wright: Intermetallics 1 (22)