Structure and properties of amorphous metal hydrides

Structure and properties of amorphous metal hydrides

Journal of the Less-Common STRUCTURE HYDRIDES* KENJI Metals, 89 (1983) AND PROPERTIES 183 183 - 195 OF AMORPHOUS METAL SUZUKI The Researc...

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Journal

of the Less-Common

STRUCTURE HYDRIDES*

KENJI

Metals,

89 (1983)

AND PROPERTIES

183

183

- 195

OF AMORPHOUS

METAL

SUZUKI

The Research (Japan) (Received

Institute

for Iron, Steel and Other Metals,

Tohoku

University,

Sendai 980

May 31,1982)

Summary Current progress in investigations of the structure and properties of including hydrogen absorption-desorption amorphous metal hydrides, characteristics, hydrogen-induced surface segregation, the static and dynamic density of states, environments around hydrogen atoms, the electronic positron annihilation and normal and superconducting properties, is reviewed.

1. Introduction The polyhedral unit structures existing in a dense random packing of soft spheres, which were used in the ideal structure model for amorphous pure metals developed by Finney and Wallace [ 11, are classified into two categories based on tetrahedra and octahedra. In amorphous structures some distortion exists in the polyhedral unit structures and there are many more tetrahedra than octahedra, in contrast with what is observed in crystalline structures. High resolution observation of the radial distribution functions has shown that chemical short-range order is preserved in both metalmetalloid and metal-metal amorphous alloys and is quite similar to that found in the corresponding crystalline alloys 121. Hydrogen and deuterium atoms are easily absorbed into the polyhedral sites in amorphous metal-metal alloys consisting of a combination of light and heavy transition metals. Therefore it is expected that hydrogen and deuterium atoms can be used as probes for investigating the atomic scale structure and dynamics of both amorphous and crystalline metals. In this paper current investigations of the hydrogen absorptiondesorption characteristics, atomic structure and electronic properties of hydrogenated amorphous alloys are reviewed. Particular emphasis is placed *Paper presented at the International of Metal Hydrides, Toba, Japan, May 30

0022-5088/83/0000-0000/$02.75

Symposium

on the Properties

and Applications

-June 4, 1982. @ Elsevier Sequoia/Printed

in The Netherlands

184

on the examination of the static and dynamic environments of hydrogen atoms by highly resolved pulsed neutron scattering measurements, the observation of electronic partial densities of states by soft X-ray spectroscopy, the identification of valence electrons by positron annihilation experiments and the investigation of the normal and superconducting properties of hydrogenated amorphous Zr-Ni and Zr-Pd alloys.

2. Hydrogen

absorption-desorption

characteristics

Amorphous alloys can be charged electrolytically [3] or from the gas phase [4, 51. Electrolytic hydrogenation provides a macroscopically inhomogeneous distribution of hydrogen atoms in the alloys such that they rupture during electrolysis before reaching the maximum hydrogen content [4] . Therefore the electrolytically hydrogenated amorphous alloys often contain regions of hydrided alloy together with regions of non-hydrided or partially hydrided alloy. When hydrogen atoms are introduced from the gas phase and sufficient time is allowed to ensure that equilibrium is reached, homogeneously hydrogenated amorphous alloys with a relatively high hydrogen content are obtained [4,5]. Whether amorphous alloys retain their amorphous structure or transform into crystalline phases during the hydrogenation process depends sensitively on the temperature, pressure and duration of the absorptiondesorption process. Amorphous alloys can usually be hydrogenated without crystallization when hydrogen atoms are introduced at low temperatures (below 100 “C) and under high hydrogen pressures of 10 - 30 atm. However, hydrides of amorphous alloys are decomposed and/or transformed into crystalline phases during hydrogenation processes at high temperatures and low hydrogen pressures as follows [4,6, 71: TiCu(a) + Hz -+ TiH,(a or c) + Cu(c)

(1)

Zr*Ni(a) + Hz --f ZrHz(c) + ZrzNi,(c)

(2)

ZrNi(a) + Hz -+ ZrNiH,(c) It should be noted that crystalline Zr,Ni(c)

(3) ZrzNi undergoes

+ Hz + ZrHz(c) + ZrNiH,(c)

the reaction

[6] (4)

rather than reaction (2) as is the case for amorphous ZrzNi. Amorphous alloys become brittle but do not spontaneously disintegrate even after 100 hydrogen absorption-desorption cycles [8]. No pressure plateaux have yet been reported for the composition-pressure isotherms of amorphous alloys consisting of light and heavy transition metals (Fig. 1) [ 5, 61. These features are characteristic of hydrogenated amorphous alloys and are different from the observations for the corresponding crystalline alloys. Since the local environment around hydrogen atoms absorbed in amorphous alloys is rather similar to that in crystalline alloys [ 10,111, the

185

. Amorphous

%NG3

%oNi50

0 Cryslalllne

iI

i

i i



0.2

04

06

0.8

1.0

1.2

0

(HIM)

Fig. 1. Cotiposition-pressure different temperatures [9].

01

02

0.3

04

05

0.6

xH

isotherms of amorphous

and crystalline

ZrNi alloys at

Fig. 2. Sieverts’ law plots for hydrogen absorption in amorphous ZrNit, ZrNi and ZrzNi 191.

two-phase region may not be absent in hydrogenated amorphous alloys. The critical temperature for phase separation in hydrogenated amorphous alloys is expected to be too low to allow equilibrium to be reached within a reasonable experimental time. Figure 1 shows that the maximum hydrogen content absorbed in the amorphous ZrNi alloy is about two-thirds of that in crystalline ZrNi under the same conditions of temperature and pressure [6] . The case of the ZrNi alloy is rather unusual. Amorphous alloys generally have larger hydrogen absorption capacities than the corresponding crystalline alloys because the severe restriction of hydrogen atomic sites by the crystal structure is relaxed in the amorphous ahoy. Aoki et al. [6] have observed that amorphous Zr-Ni alloys can store two hydrogen atoms per zirconium atom over the whole composition range. Furthermore hydrogen absorption in amorphous Zr-Ni alloys obeys Sieverts’ law in various concentration ranges as shown in Fig. 2. This may mean that some of the hydrogen atomic sites are selectively occupied depending on the hydrogen content in the amorphous alloys. Figure 1 shows that there is a large hysteresis between absorption and desorption in the composition-pressure isotherms of amorphous alloys. This hysteresis may originate not only from the kinetics of hydrogen absorption and desorption which are controlled by the surface conditions but also from the structural relaxation which occurs during hydrogen absorption and desorption processes. However, a proton magnetic resonance study carried out by Panissod and Mizoguchi [12] has indicated that there are supplementary paths with higher activation energies for hydrogen diffusion in amorphous Zr,.,, Pd0.35H, (x = 1.7, 1.1) although the primary mechanism

186

for hydrogen diffusion in the amorphous and crystalline alloys is essentially the same. Marked changes in the surface compositions of amorphous alloys are induced by repeated hydrogen absorption and desorption. Surface segregation in amorphous Zr,Ni was first observed by Spit and coworkers [7, 81 using Rutherford backscattering and Auger electron spectroscopy measurements. According to their observations a tetragonal ZrOz layer 100 - 150 A thick is formed on the surface of amorphous Zr,Ni after 100 cycles of hydrogen absorption and desorption. During this cycling a nickel-enriched layer appears immediately below the tetragonal ZrOz layer and a third layer containing ZrOz and nickel crystals is produced below the nickel-enriched layer by internal oxidation. Since hydrogen molecules easily permeate the tetragonal ZQ layer and dissociate catalytically into hydrogen atoms in the nickel-enriched layer, the hydrogen absorption rate is substantially enhanced by repeated hydrogen absorption and desorption. The surface of as-quenched amorphous Zr,Ni is covered with a monoclinic ZrO, layer 50 - 100 A thick which prevents hydrogen molecules from penetrating inside the amorphous Zr*Ni. However, it is well known that hydrogen is very quickly absorbed by amorphous Zr,Ni even at room temperature if the ZrOz scale is removed by abrading with sandpaper. In some cases annealing in an oxygen atmosphere at low temperature also accelerates the rate of hydrogen absorption. Spit et al. [8] have found that the oxidation of amorphous Zr,Ni at 590 K for 30 min results in the growth of a ZrOz layer about 150 A thick on the surface and the formation of an internally oxidized layer including ZrOz and nickel crystals between the ZrOz scale and the nickel-enriched layer. Surface segregation due to low temperature oxidation again results in an increase in the rate of hydrogen absorption in Zr,Ni. Oelhafen et al. [ 131 have carried out UV photoelectron spectroscopy, X-ray photoelectron spectroscopy and Auger electron spectroscopy measurements on amorphous Zr,OPd,,-, and its hydride containing about 20 at.% H and have concluded that there are no palladium atoms within the first few atomic layers at the surface of as-quenched amorphous Zr,,Pd,,. However, hydrogen-induced palladium segregation takes place at the hydride surface and zirconium atoms are preferentially sputtered from it.

3. Local environments

around hydrogen

atoms

Neutron scattering experiments have the unique advantage of being able to detect hydrogen and deuterium atoms sensitively in amorphous alloys since the incoherent neutron scattering cross section of the nucleus of the hydrogen atom has a relatively large value (79.7 b) and the coherent neutron scattering cross section of the nucleus of the deuterium atom has a medium value (5.6 b) [ 141.

187

The deuterium and metal atoms are quite close together and the vibrational energy of the hydrogen atoms is much higher in amorphous alloys because the sizes and masses of the hydrogen and deuterium atoms are extremely small. Therefore short wavelength incident neutrons with epithermal energy are required to obtain well-resolved experimental observations of the local environment around a deuterium atom and the localized hydrogen vibration in amorphous alloys. The use of high flux pulses of fast neutrons generated in an accelerator [15, 161 rather than steady state neutrons from a reactor is recommended. Figure 3 shows the total radial distribution functions (RDFs) 4nr2pg(r) of amorphous ZrO.,sPd 0.35Ds (3t = 0 - 1.25) obtained by Kai et al. [lo] using a spallation neutron source installed at the 500 MeV proton booster synchrotron in the National Laboratory for High Energy Physics, Japan. The first peak at r = 2 ,& corresponding to the D-M (M - Zr, Pd) correlation is clearly distinguished from the higher-order peaks of the M-M correlations. The coordination number of the metallic atoms around a deuterium atom was calculated as a function of the deuterium content from the first peak area, assuming that the local chemical composition of metallic atoms surrounding a deuterium atom is similar to the average chemical composition of metallic atoms in the amorphous alloy. This simple assumption provides reliable information on the topology of polyhedral unit structures around a deuterium atom because the magnitudes of the coherent neutron scattering length for zirconium (bz, = 0.694 X 10 I2 cm) and palladium (bPd = 0.603 X 10 l2 cm) nuclei are almost equal [ 141. We can conclude that the deuterium atoms in amorphous Zr0.6SPd0.35Dx with a low deuterium content initially prefer to occupy tetrahedral sites consisting of four M atoms. As the deuterium content increases to 3c z 0.42 the deuterium atoms are located at octahedral sites with an average of six M atoms around a deuterium atom. Finally the average local environment around each deuterium atom approaches a hexahedral site consisting of five M atoms at a deuterium content of x = 1.25. These types of polyhedrai site have been identified in crystalline transition metal hydrides such as PdH, [ 171, ZrH, [ 181 and ZrNiH, [ 191. It should be noted that the octahedral and hexahedral sites in amorphous Zr 0.6SPd0.35Ds may indicate Jahn-Teller splitting because the first peak in the RDF is split into two subpeaks corresponding to two different D-M separations. The RDF of amorphous ZrNiDi.s shows a marked separation between the D-Ni and D-Zr correlations, as illustrated in Fig. 4 [ 11, 201, because of the large difference in the atomic sizes of the nickel and zirconium atoms. The small peak located at r = 1.7 i$ is due to the D-Ni correlation, while the large peak at r = 2.1 A corresponds to the D-Zr correlation. The atomic separations of the D-Ni and D-Zr correlations in the hydrogenated amorphous alloy are almost equal to those of the H-Ni (r = 1.66 A) and H-Zr (r = 2.1 A) correlations at the tetrahedral site in crystalline ZrNiH, [19]. The coordination of the metallic atoms surrounding a deuterium atom is calculated from the areas under the two peaks to be one nickel and three

20.

15 -

Xr t.25

0 0

12

/

3 r

4 (A)

L a-ZrNiDl6

,

5

6

7

r

(A)

Fig. 3. Total RDFs of amorphous Zro.6$do.35D, (w = 0 - 1.25) [ 10 1. Fig. 4. Total RDFs of amorphous ZrNi and ZrNiD1.8 [ 111.

zirconium atoms. This result supports the fact that deuterium atoms in ZrNiD,,s prefer to occupy tetrahedral sites consisting on average of one nickel atom and three zirconium atoms, while hydrogen atoms in crystalline ZrNiH, occupy tetrahedral sites consisting of one nickel atom and three zirconium atoms and hexahedral sites consisting of two nickel atoms and three zirconium atoms [19] . The local env~onment around hydrogen atoms in hydrogenated amorphous alloys is not a statistical configuration of the constituent atoms but is the same as that in the corresponding crystalline compounds. The energy spectrum of the localized hydrogen vibration g(w) also provides information regarding the topological distortion and chemical ~u~tuation of the local env~onment around a hydrogen atom in ~o~hous alloys. Figure 5 shows g(w) for crystalline ZrNiHzs and amorphous ZrNiH,., measured by Kaneko et al. [II, 201 using a time focusing crystal analyser spectrometer installed at a (r,n) pulsed neutron source in the Tohoku University 300 MeV Electron Linac Laboratory. The peak centres of the g(o) are located at Bw = 130 + 3 meV in both the crystalline and amorphous hydrides. The fact that there is no variation in the position of the peak centre implies that the sites occupied predominantly by hydrogen atoms in the crystalline and amorphous states of the ZrNi alloy are equivalent. However, the width of the g(w) of amorphous ZrNiHr.s (full width at halfmaximum (FWHM), 70 + 5 meV) increases by about 40% compared with that of crystalline ZrNiHqs (F~M, 50 2 5 meV) although the hydrogen content of the amorphous alloy is less than that of the crystalline ahoy. The g(w) of amorphous ZrNiH1.s has broad tails at each extreme of the energy range, i.e. below fiw = 100 meV and above ho = 150 meV. The root mean squares of the hydrogen vibration amplitudes obtained from the Debye-

0 Energy

Transfer

I meV

1

.2

3

4

.5 hd

.6 (eV)

.7

8

91.0

Fig. 5. Energy spectra of localized hydrogen vibration in crystalline ZrNiI3z.s and amorphous ZrNiH,.s [ll, 211. Fig. 6. Energy spectra of the focalized hydrogen vibration in crystalline TiH2 and in crystalline and amorphous TizNiH 1.5 [ 22 1.

Waller factors for crystalline ZrNiHzSs ( (Ux 2)“2 = 0.192 A) and amorphous ZrNiH,_s ((~,~)r’~ = 0.205 A) are not significantly different [21]. We conclude from the observations discussed above that hydrogen and deuterium atoms absorbed in ~o~hous ZrNi alloys are preferenti~ly located at tetrahedral sites where the topological distortion is rather restricted and the chemical order in the tetrahedral unit structures fluctuates appreciably around an average composition of one nickel atom and three zirconium atoms. The tetrahedral site consisting of two nickel atoms and two zirconium atoms may enhance the contribution to g(w) from the lower energy vibration, while hydrogen atoms located in the tetrahedron constructed from four zirconium atoms shift the position of the centre of g(w) towards the higher energy region. The hydrogen vibration spectra of crystalline TiH2 and of crystalline and amorphous Ti2NiH1_,measured by Kai et al. [ 221 using a high resolution crystal analyser time~f-flight spectrometer installed at the spallation neutron source are compared in Fig. 6. It should be noted that higher harmonics in the localized hydrogen vibration are present even in the amorphous alloy. The low energy tail appearing below Rw = 100 meV in crystalline TizNiH,., is retained but becomes very broad owing to topological and chemical

190

fluctuations of the hydrogen atom sites in amorphous TizNiHlos. Hydrogen atoms in crystalline TiHz occupy only tetrahedral sites corresponding to an energy Ro of about 150 meV. The results of X-ray and neutron diffraction studies of crystalline Ti*NiH(D),,, [ 231 suggest that the low energy tail may be due to hydrogen atoms located in octahedral sites in both the crystalline and the amorphous hydrides. Therefore it is again concluded that hydrogen atoms in amorphous TizNiH t .5 remain in essentially the same environment as in the crystalline T&NiHi.s.

4. Electronic

states and transport

properties

Soft X-ray spectroscopy (SXS) is a particularly powerful technique for studying the electronic energy levels of complex disordered materials such as hydrogenated amorphous alloys since the partial structures of the electronic states in the vicinity of a specified atom can be selectively measured by tuning the energy to the absorption or emission edge. Figure 7 shows the X-ray Ls (sd + 2~s,~) emission spectra of zirconium atoms in amorphous Zr,,65Pd0.35 and Zr0.67Ni0.33 and in their amorphous hydrides with hydrogen-to-metal ratios of unity together with that of pure zirconium metal measured by Tanaka et al. [24]. The Zr L3 emission spectra are drastically modified by hydrogen or deuterium absorption. A striking observation is that a new low energy shoulder appears about 7 eV below the emission edge and the main peak contributed by the 4d state is suppressed and shifted towards a higher energy in both amorphous ZrO.ssPd,,.ssHl.O and of the Zr L3 spectra inamorphous Zr0.67Ni0.33D1.o. Such a modification duced by hydrogenation has been reported in crystalline ZrH, [25]. In contrast with the Zr L3 spectra, the addition of hydrogen or deuteriurn atoms to amorphous Zro.ss Pdo.3s and Zr0.67Ni0.33 has little effect on either the Pd L3 or the Zr L3 emission spectra as shown in Fig. 8. These results indicate that hydrogen and deuterium atoms in the amorphous alloys prefer to form bonding states with the zirconium atoms at the bottom of the main peak and consequently to rearrange the electronic states about the zirconium atoms near the Fermi level by effectively avoiding interactions with the palladium and nickel atoms. The Zr L, and Pd L3 emission spectra of crystalline Zro.ssPdo.35 and Zro.ssPdo.35H0.61 have been confirmed to be very similar to those of the amorphous alloys [26]. Therefore the characteristic behaviour of the soft X-ray emission spectra described above is not restricted to the amorphous alloys but is common in both the crystalline and the amorphous alloys. This conclusion is supported by investigations of the atomic structure of the hydrogenated amorphous alloys which show that the local environments around hydrogen atoms in the amorphous alloy and the corresponding crystalline alloy are similar. The formation of bonds between hydrogen and zirconium atoms in the hydrogenated amorphous alloys supports the neutron

I 2210

2215 2220 ENERGY/&

Pd

13

band

2225 ENERGY

I

eV

Fig. 7. X-ray Zr L3 emission spectra of amorphous Zr0.65Pd,,3s and Zr0.67Ni”.33, their hydrides and a pure zirconium crystal [ 24 1. Fig. 8. X-ray Pd LJ emission spectra of amorphous Zro.6sPdo.35, its hydride and a pure palladium crystal, and X-ray Ni L3 emission spectra of amorphous Zr0.67Ni0.33r its hydride and a pure nickel crystal [ 24 1.

scattering observation that even in amorphous alloys hydrogen atoms prefer to occupy polyhedral sites constructed mainly from zirconium atoms. A marked change in the behaviour of the conduction electrons in amorphous alloys takes place when hydrogen is absorbed. The conduction electron concentration in amorphous Zr o.65Pd0.35Hx is expected to increase proportionally with the hydrogen concentration if hydrogen atoms donate one conduction electron per atom. Figure 7 shows that the addition of hydrogen or deuterium atoms elevates the Fermi energies of amorphous Zr0.6sPdo.s5Hi.o and Zra.6&.&.o. Figure 9 shows the angular correlation curves of positron annihilation in amorphous ZroessPd0.s5Hr (x = 0 - 0.83) measured by Itoh et al. [27]. The central peak height around 0 mrad is lower in hydrogenated amorphous alloys when the angular correlation curves are normalized to the number of electrons present below 15 mrad. This result implies that the Fermi momentum is increased and that positron trapping is suppressed by the introduction of hydrogen atoms. However, the angular correlation curves are enhanced over the intermediate angle from 5 to 10 mrad with increasing hydrogen content. This observation corresponds to the formation of bonds between hydrogen and zirconium atoms detected by SXS measurements. The Fermi cut-off angles BFexp in amorphous Zr0.65Pd0.35Hx- obtained by least-squares fitting to a gaussian curve for the core electrons are plotted as a function of the hydrogen content in Fig. 10. The experimental values

.

ANGLE

( mrad )

Fig. 9. Angular correlation curves for positron annihilation in amorphous Zro.ssPdo.ssH, (x = 0 - 0.83) [27]. Fig. 10. Fermi momenta and densities of amorphous Zr0.s5Pd0_35Hx (x = 0 - 0.83) [27].

13~~ at low hydrogen contents (hydrogen-to-metal ratios less than or equal to 0.3) are in good agreement with the theoretical values eFCa’ calculated using the free-electron model and assuming the donation of one conduction electron per palladium atom, two conduction electrons per zirconium atom and one conduction electron per hydrogen atom. The disagreement between ratios greater or’= and Orcal at high hydrogen contents (hydrogen-to-metal than or equal to 0.4) can be ascribed to the formation of bonds in amorphous Zr0.6sPd0.35HX. Despite the increase in the Fermi momentum, the electrical conductivity of amorphous alloys decreases substantially on hydrogen absorption. For example the electrical resistivity of amorphous Zr0.65Pd0.35 is 220 1.ts1cm at room temperature, while that of amorphous Zr0.65Pd0.35H1.01 is 270 /.LJ cm [28]. The temperature coefficients of electrical resistivity (TCRs) of amorphous alloys of light and heavy transition metals are usually negative and quite small. After hydrogen absorption, the TCRs of amorphous Zr0.65Pd0.35H(D)x [29] and Zr,NiH, [ 301 are still negative but are much larger. As shown in Fig. 11 Kai et al. [ 281 have found that two TCR regions appear in the resistivity uersus T2 plot for amorphous Zr0.65Pd0.35H1.01 whereas amorphous Zr0.65Pd0.35 obeys the T2 law over the whole temperature range examined. This may be an indication that optical phonon scattering is well separated from acoustic phonon scattering in amorphous Zr 0.65Pdo.&I1.0r because there is a large difference between the atomic masses of the metal and hydrogen atoms.

193

0

40

80

140

200

300

T2

(lo* K*)

400

Fig. 11. Electrical resistivities of Zr o.&‘do.35 normalized at 300 K U.S.T2 [ 281.

and

Zro.&‘do.&1.o1

amorphous

allw

Ziman’s diffraction theory including the Debye-Waller factor [ 311 and an experimental structure factor [lo] was used to calculate the Debye temperatures of amorphous Zr0.65Pd0.35H1_01 which were found to be 6naC = 200 K from the low temperature resistivity data and Onop = 1100 K from the high temperature resistivity data (Fig. 11). The Debye temperature Bn of amorphous Zr,.,5Pd,,s, was found to be 270 K. The magnitude of the Debye temperature for the optical phonon branch obtained from the electrical resistivity data is in good agreement with the frequency of hydrogen vibrations measured using inelastic neutron scattering. Kai et al. [29] have observed that amorphous Zr0.65Pd0.35 and Zr 0.65Pd0.35Hx (X = 0 - 0.4) undergo a superconducting transition, the temperature of which decreases linearly with increasing hydrogen content. The transition temperature is fairly well defined in amorphous Zr0.65Pd0.35 but becomes much more diffuse on the addition of hydrogen. This means that inhomogeneities of dimensions ranging up to several hundred angstroms are induced in hydrogenated amorphous alloys. Similar results have been obtained for amorphous ZrzNi and ZrzNiH, (X < 0.37) by Babic et al. [30].

5. Conclusion Hydrogen and deuterium atoms are easily absorbed into amorphous alloys consisting of light and heavy transition metals. The compositionpressure isotherms of the hydrogenated amorphous alloys studied to date do not show a pressure plateau. Amorphous alloys usually absorb more hydrogen than the corresponding crystalline alloys. During repeated hydrogen absorption and desorption surface segregation is induced without spontaneous disintegration. Amorphous alloys expand almost linearly with increasing hydrogen content. The absorbed hydrogen atoms preferentially occupy tetrahedral, hexahedral and/or octahedral sites with topological distortion and chemical

194

fluctuation and then selectively form chemical bonds with light transition metal atoms even in amorphous alloys. Therefore large deviations from the free-electron model are obtained in amorphous alloys with a high hydrogen content. Superconducting transition temperatures are extremely sensitive to hydrogen absorption.

References 1 J. L. Finney and J. Wallace, J. Non-Cryst. Solids, 43 (1981) 165. 2 T. Fukunaga, K. Kai, M. Naka, N. Watanabe and K. Suzuki, in T. Masumoto and K. Suzuki (eds.), Proc. 4th Int. Conf. on Rapidly Quenched Metals, Sendai, August 1981, Vol. 1, Japan Institute of Metals, Sendai, 1982, p. 347. 3 A. J. Maeland, in A. F. Andresen and A. J. Maeland (eds.), Hydrides for Energy Storage, Pergamon, Oxford, 1978, pp. 447 - 462. 4 A. J. Maeland, L. E. Tanner and G. G. Libowitz, J. Less-Common Met., 74 (1980) 279. 5 F. H. M. Spit, J. W. Drijver and S. Radelaar, 2. Phys. Chem. N.F., 116 (1979) 225. 6 K. Aoki, A. Horata and T. Masumoto, in T. Masumoto and K. Suzuki (eds.), Proc. 4th Int. Conf. on Rapidly Quenched Metals, Sendai, August 1981, Vol. 2, Japan Institute of Metals, Sendai, 1982, p. 1649. 7 F. H. M. Spit, J. W. Drijver, W. C. Turkenburg and S. Radelaar, J. Phys. (Paris), Colloq. C8, 41 (1980) 890. 8 F. Spit, K. Blok, E. Hendriks, G. Winkels, W. Turkenburg, J. W. Drijver and S. Radelaar, in T. Masumoto and K. Suzuki (eds.), Proc. 4th Int. Conf. on Rapidly Quenched Metals, Sendai, August 1981, Vol. 2, Japan Institute of Metals, Sendai, 1982, p. 1635. 9 K. Aoki, M. Kamachi and T. Masumoto, to be published. 10 K. Kai, T. Fukunaga, T. Nomoto, N. Watanabe and K. Suzuki, in T. Masumoto and K. Suzuki (eds.), Proc. 4th Int. Conf. on Rapidly Quenched Metals, Sendai, August 1981, Vol. 2, Japan Institute of Metals, Sendai, 1982, p. 1609. 11 H. Kaneko, T. Kajitani, M. Hirabayashi, M. Ueno and K. Suzuki, in T. Masumoto and K. Suzuki (eds.), Proc. 4th Znt. Conf. on Rapidly Quenched Metals, Sendai, August 1981, Vol. 2, Japan Institute of Metals, Sendai, 1982, p. 1605. 12 P. Panissod and T. Mizoguchi, in T. Masumoto and K. Suzuki (eds.), Proc. 4th Int. Conf. on Rapidly Quenched Metals, Sendai, August 1981, Vol. 2, Japan Institute of Metals, Sendai, 1982, p. 1621. 13 P. Oelhafen, R. Lapka, U. Gubler, J. Krieg, A. DasGupta, H.-J. Giintherodt, T. Mizoguchi, C. Hague, J. Kiibler and J. R. Nagel, in T. Masumoto and K. Suzuki (eds.), Proc. 4th Int. Conf. on Rapidly Quenched Metals, Sendai, August 1981, Vol. 2, Japan Institute of Metals, Sendai, 1982, p. 1259. 14 B. T. M. Willis (ed.), Chemical Applications of Thermal Neutron Scattering, Oxford University Press, Oxford, 1973. 15 K. Suzuki, in T. Masumoto and K. Suzuki (eds.), Proc. 4th Znt. Conf. on Rapidly Quenched Metals, Sendai, August 1981, Vol. 1, Japan Institute of Metals, Sendai, 1982, p. 309. 16 K. Suzuki, M. Misawa, T. Fukunaga and N. Hayashi, in C. Hargitai, I. Bakonyi and T. Kemeny teds.), Proc. Conf on Metallic Glasses: Science and Technology, Vol. 1, Central Research Institute for Physics, Budapest, 1981, p. 327. 17 J. E. Worsham, M. K. Wilkinson and C. G. Shull, J. Phys. Chem. Solids, 3 (1957) 363. 18 E. L. Slaggie, J. Phys. Chem. Solids, 29 (1968) 923. 19 D. G. Westlake, J. Less-Common Met., 75 (1960) 177.

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25 26 27 28 29 30

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H. Kaneko, T. Kajitani, M. Hirabayashi, M. Ueno and K. Suzuki, J. Less-Common Met., 89 (1983) 4643. H. Kaneko, M.Sc. Thesis, Tohoku University, 1982. K. Kai, S. Ikeda, N. Watanabe and K. Suzuki, in Y. Ishikawa and N. Watanabe (eds.), KENS Rep. III, 1982 (National Laboratory for High Energy Physics, Tsukuba). H. Buchner, M. A. Gutijar, K.-D. Beccu and H. Saufferer, 2. Metallkd., 63 (1972) 497. K. Tanaka, M. Higatani, K. Kai and K. Suzuki, J. Less-Common Met., 88 (1982) 317. K. Tanaka, N. Hamasaka, M. Yasuda and Y. Fukai, Solid State Commun., 30 (1979) 173. K. Tanaka, M. Higatani, K. Kai and K. Suzuki, to be published. F. Itoh, K. Kai, T. Yoda and K. Suzuki, to be published. K. Kai, T. Nomoto and K. Suzuki, Sci. Rep. Res. Inst., Tohoku Univ., Ser. A, 29 (1981) 204. K. Kai, T. Nomoto, M. Ikebe and K. Suzuki, J. Less-Common Met., 89 (1983) 229. E. Babic, B. Leontic, J. Lukatela, M. Miljak and M. G. Scott, in T. Masumoto and K. Suzuki (eds.), Proc. 4th Int. Conf on Rapidly Quenched Metals, Sendai, August 1981, Vol. 2, Japan Institute of Metals, Sendai, 1982, p. 1617. P. J. Cote and L. V. Meisel, in H.-J. Giintherodt and H. Beck (eds.), Glassy Metals I, Springer, Heidelberg, 1981, p. 141.