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Hydrogen absorption in CuTi metallic glasses. I. X-ray diffraction measurements

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1985 J. Phys. F: Met. Phys. 15 2069

(http://iopscience.iop.org/0305-4608/15/10/005)

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J. Phys. F: Met. Phys. 15 (1985) 2069-2083. Printed in Great Britain

Hydrogen absorption in CuTi metallic glasses. I: x-ray diffraction measurements

B Gr ie ta t , K Din$, N Cowlam and H A Daviesg Department of Physics, University of Sheffield, Sheffield S3 7RH, UK 0 Department of Metallurgy, University of Sheffield, Sheffield S I 3JD, UK

Received 21 February 1985

Abstract. The effect of hydrogen absorption on the atomic structures of some CuTi glasses has been investigated by x-ray diffraction. The dependence on composition has been studied using three alloy compositions, Cuo.soTio.so, Cuo.5sTio.4z and Cuo,66Tio.,s, and the influence of loading increasing amounts of hydrogen investigated for the latter two compositions. It is confirmed that these glasses can accept significant quantities of hydrogen to form amorphous hydrides and that the process is limited by a dissociation reaction involving the formation of titanium hydride TiH2 and free copper. Quite characteristic changes in the structure factor are produced by hydrogen loading and these are reflected in definite changes in the radial distribution function, similar to those observed elsewhere for other metallic glass-hydrogen systems. The results obtained in this study are in broad agreement with the conclusions of independent bulk measurements on CuTi-H glasses.

1. Introduction

The reversible absorption or occlusion of hydrogen in palladium and titanium is well known and has led to the widespread use of these metals and their alloys as gettering agents and in gas-separation devices. Furthermore, it has been possible for some time now to produce certain alloys of these two metals, such as Pd-Si, Pd-Ge, Ti-Cu and Ti-Ni, in the form of metallic glasses by quenching directly from the melt and there is some likelihood that these newer materials may also be used in practical devices in the future. It is necessary to impose high cooling rates, typically 105-106 K s- I , in the production of these metallic glasses in order to access the glassy phase. This is usually achieved by bringing a puddle of the molten alloy into good thermal contact with a rapidly moving, well conducting, roll substrate. The process, known as chill-block melt-spinning (see e.g. Davies 1978) can be used to produce continuous lengths of thin metallic glass ribbon having a width typically w= 1-10 mm and thickness t= 10-30pm for roll surface velocities upwards of 25 m s - ' . Lengths of such ribbon have a high value of the ratio surface area/volume (of order 2 / t z 1000 cm-'), which is commensurate with a finely divided powder with grain size - 60 pm. This suggests that palladium- and titanium-based metallic glass ribbons may have potential application in gas absorption devices in which a large surface area of the active agent is required. Yet another factor is the evidence from density measurements that the atomic packing in the amorphous state is slightly less than in

t Permanent address: Ruder BoSkoviC Institute, 4 1001 Zagreb, BijeniEka 54, Croatia, Yugoslavia. $ Present address: Department of Physics, Dalhousie University, Halifax, Nova Scotia, Canada, B3H 3JS.

0305-4608/85/102069 + 15 $02.25 0 1985 The Institute of Physics 2069

2070 B Grieta et a1

the equivalent crystalline alloy, so that a metallic glass may therefore be able to accommodate a higher concentration of hydrogen than the crystalline counterpart.

Whilst these and other factors have been responsible for an increased technological interest in metallic glasses, there is evidence that hydrogen absorption may also provide information of a more fundamental kind. The hydrogen atom is a particularly appropriate structural probe, which should allow the geometries of the interstices of metallic glasses to be investigated on account of its small atomic size. Nuclear magnetic resonance measurements, for example, can provide information about the static and dynamic behaviour of absorbed hydrogen, including diffusion. In magnetism, hydrogen absorption offers an opportunity of altering the strengths of magnetic exchange interactions by controlled expansion of the host structure. This can allow the relative influences of interatomic distance and of electron concentration on the exchange interactions to be resolved more effectively than is sometimes the case in conventional alloying work.

We have made an investigation of the effects of hydrogen absorption on the structure of some CuTi glasses using x-ray diffraction, as an extension of our neutron and x-ray diffraction investigations of the as-quenched glasses (Sakata et a1 1980, 1982). Although there have been previous studies of the effects of hydrogen on metallic glass structure (e.g. Kai et a1 1982) there have been few systematic investigations, with the possible exception of the combined neutron elastic and inelastic scattering work of Suzuki et a1 (1984). We have therefore investigated the influence of hydrogen absorption on the structure of CuTi glasses of different composition and also the influence of different quantities of absorbed hydrogen on glasses of fixed composition. This paper gives fuller details of work already described in outline (Dini et a1 1982a). An experimental problem which can have an effect on the observations is also described in an appendix.

2. Sample preparation and the experimental methods

Master ingots having the compositions Cuo.soTio.so, Cuo,sI Ti0.42 and Cuo.66 Ti0,34, as well as Pdo,soSi0,20 were prepared by argon arc melting of spectrographically pure constituents with negligible weight loss. Material from these ingots was conventionally melt-spun in air on to a copper wheel, to give metallic glass ribbon of width typically 1 mm and thickness 25 pm. Specimens of these ribbons were measured in the as-quenched states, but the majority of the experiments were made on glasses loaded with hydrogen. An electrolytic cell was found to be the most convenient of several arrangements tested for the hydrogen loading. A small cubic vessel 5 cm x 5 cm x 5 cm was fitted with three platinum anode plates 5 cm x 3 cm, between which two metallic glass cathodes were placed. These were made by winding a length of - 2 m (0.3 g) of melt-spun ribbon on to flat open PVC formers 4.5 cm x 4 cm. A 0.05 M solution of H 2 S 0 4 in H,O was used as electrolyte under the conditions -0.3 V and 100-150 mA. A 48 h loading period was used initially, which produced significant changes in the glasses of all three compositions but also induced some crystallisation of the C U ~ , , , T ~ ~ , , ~ and Cuo,soTio,so. New samples of these were examined after shorter loading times (as specified in table 1) so that both the structural changes before any crystallisation and the details of the whole loading process could be investigated.

Hydrogen loading was found to cause embrittlement of the CuTi ribbon, severe buckling (probably associated with the increase in volume on taking up the gas) and some surface discolouration. It appeared unlikely that these macroscopic changes would be reversible and, to determine if this were also true for the atomic-scale structure, certain ribbon lengths were heat-treated to drive off the hydrogen. This was performed under a

Hydrogen absorption in CuTi metallic glasses 207 1

steady flow of argon to prevent oxidisation. Initially a treatment of 30 min at 150 "C (423 K)-well below the normal crystallisation temperatures of the as-quenched CuTi glasses, which are in the region of 420 "C (693 K) (Sakata et a1 1982)-was used, but this resulted in the appearance of Bragg peaks in the scattered intensity distribution. All three glasses were then treated for a similar period at 80 "C (353 K). Nuclear magnetic resonance experiments showed that the hydrogen had been evolved at this lower temperature (J M Titman, private communication).

The hydrogen concentrations in the samples were determined by chemical analysis kindly performed at the Department of Metallurgy and Materials Science, at the University of Cambridge. The values, expressed in the usual chemical equations are given in table 1. It will be seen that for Cu0.38 Ti,,, and Cu0,66Ti0,34 the hydrogen content increases regularly with time in the electrolytic cell, while the hydrogen concentrations after 48 h in each case increase with increasing titanium concentration. The hydrogen concentrations given in table 1 are comparable with, and slightly higher than, those obtained by loading with hydrogen gas reported by Rush et a1 (1980), Maeland (1978), and Bowman et a1 ( 1982)-who concentrated on the equiatomic and titanium-rich glasses.

In the literature there is some ambiguity, by a factor of two in the quoted hydrogen concentrations, arising from the use of chemical formulae for the metallic glass concentrations, e.g. TiCu = Ti, Cu, which is also Tio,~oCuo,so (Maeland 1978, p 453). It is logical to use chemical formulae for crystalline intermetallic compounds which usually have simple superlattice structures but we feel that their application to structurally disordered metallic glasses is less appropriate, and have used percentage alloy concentrations throughout this paper.

The main purpose of this work has been a direct investigation of the structural changes accompanying hydrogen absorption using x-ray diffraction. These measurements were made conventionally using a Philips PW 1050 vertical goniometer, with a graphite curved- crystal monochromator and molybdenum K, radiation, A = 0.7 1 1 A. In the standard experiment a range of scattering angles, 5" < 28 < 1 60°, corresponding to scattering

Table 1. This table shows the specimens used, the conditions of electrolysis using the cell described in the text and the hydrogen concentrations obtained.

Electrolysis time (h) at Comment;

Sample 0.3 V and 150 mA Hydrogen content further treatment

Cu0.50Ti0.50 - 48 48

Cu0.S8Ti0.42 - 12 I 2 12 24 48

Cu0.66Ti0.34 - 6 6 12 24 48

as-quenched

80 OC for 30 min

as-quenched

80 "C for 30 min 150 "C for 30 min

-

-

as-quenched

80 "C for 30 min -

-

2012 B Grjeta et a1

vectors, 0.8 < Q < 17.4 A-‘. was covered with a step length of 0.5” and a count time of roughly 3 min with a total scan time of 18 h. The x-ray specimens were made by sticking short ribbon lengths on to the usual flat aluminium holder, with spacer shims to ensure that the upper surface of the ribbon was tangential to the focusing circle of the instrument (Dini et a1 1982b). Great care was taken to ensure that buckled lengths of the hydrogen-loaded ribbon were as flat to the sample plate as possible, since we observed that the geometry of the specimen could have a profound effect on the scattered intensity distribution and hence on the structure factor S(Q) and the reduced RDF G(r) obtained. Details of these observations and their associated effects are given separately in the appendix. The data analysis followed conventional lines, with the subtraction of background scattering and appropriate corrections to the intensity distribution, which could then be normalised in a number of different ways to obtain the structure factor S(Q). A modified version of a program originally written by C N J Wagner was used in this analysis (Wagner 1968).

3. Structural changes in CuTi glasses caused by hydrogen loading

3.1. Changes in structure factor S(Q)

The structure factors S(Q) obtained directly from the experimental data are perhaps the best curves for comparison between experiment and simulation, and will be presented first. They depend to some extent on the resolution of the x-ray instrument, but many workers now use equipment having broadly similar resolution. The structure factors for the three CuTi glasses in the as-quenched state are shown in figure 1. The variations at high Q arise from statistical fluctuations in the data points obtained, which is the case for all the structure factors shown here. The S(Q) curves for these three glasses have the characteristic and well established form, having a sharp maximum with halfwidth AQ 2 0.42 A-‘ : a second maximum and shoulder having positions roughly

3 i 4 6 0 10 12 1 4 1 6 18 Scattering v e c t o r P [I-’)

Figure 1. The structure factors for CuTi glasses in the as-quenched condition: A. Cuo.~oTio.~o: B. Cu0.58Ti0.42: C, CUo.66Tio.34.

Hydrogen absorption in Cu Ti metallic glasses 2073

L

c U

0 -

2 2.0

c 2 1.0

01 L 3 -

c U

Vl

0 -

Q2/Q, z 1.7 and Q,,/Q, 2 1.9 relative to the first; and discernible oscillations in S(Q) out to 10-12 A- ' . The small pre-peak at Q z 1.9 A-' is just detectable (Sakata et a1 1982).

The form of S(Q) is extensively modified by hydrogen absorption. The structure factors for the same three compositions after hydrogen loading are shown. in figure 2. The periods of loading are different in the three cases, and are specified together with the hydrogen concentrations in the figure caption and table 1. These periods of loading were found to produce significant changes in the glassy state without inducing any crystallisation and the consequent appearance of crystalline peaks in S(Q) which the longer periods of electrolysis lead to. The most striking changes produced are first, a reduction in height of the first peak in S(Q) by 30-40%; a significant increase of the order of 75% in the halfwidth of the first peak, which has values AQ z 0.70 A- ', and a small movement of the peak to smaller Q values. The second peak also has a more rounded outline and the shoulder is lost, while the magnitudes of the third and fourth peaks are also reduced.

The changes in the structure factor S(Q) of the Cuo,66Tio34 glass caused when increasing amounts of hydrogen were absorbed are illustrated in figure 3. The periods of electrolysis varied between 6 h and 48 h leading to the hydrogen concentrations specified in table 1. It will be noted that there is a steady increase in the number and intensities of Bragg peaks superimposed on the S(Q) curve, as the hydrogen concentration increases. These peaks arise from the presence of free copper, the (1 11)(200) and (220)(311, 222) peaks falling on the first and second maxima in S(Q) respectively. The small peak at Q - 4 A - ' is possibly from a-Ti (HCP) or alternatively from a T i H , , phase having a tetragonal unit cell. The presence or otherwise of Bragg peaks in the S(Q) curves of all samples examined is specified in table 2. The second aspect of the structural changes is most effectively illustrated by overlaying the S(Q) curves; it becomes clear that the basic form of the distorted S(Q) remains essentially unchanged, as far as can be determined, as the period of electrolysis increases.

The first annealing treatment for the C U ~ * T ~ ~ ~ - H sample (30 min at 150 "C) also caused the appearance of Bragg peaks in S(Q). The majority of these could be identified again with FCC copper, while three less intense peaks were consistent with the presence of a TiH, hydride having an FCC cell with a. - 4.4 A. As described above, a lower temperature

- -

I 1 4.0 1

s [ ~1 3.0

I 1 0 2 4 6 8 10 12 14 1 6 18

S c a t t e r i n g v e c t o r Q [ I - ' )

Figure 2. The structure factors for CuTi glasses after the absorption of hydrogen without crystallisation: A. (Cuo.soTio.so)H0.,, : B. (Cuo.ssTi0.42)H0.59: c , (Cu0.66Ti0.34)H0.~,.

2014 B Grieta et a1

0 4.0 crl

c VI

1.0 - %dwv-

0 -

0 2 4 6 a 10 12 14 1 6 18 S c a t t e r i n g v e c t o r GI ( X - ' )

anneal (30 min at 80 "C) was given to all three samples, and figure 4 shows the S(Q) curve for the (Cu0,66Ti0,34)H0,47 sample, together with the equivalent curve after heat treatment. It is clear in this case that the changes produced in the structure were irreversible (this is also the case for the other two samples) and this is consistent with the external appearance

Table 2. The first-neighbour distances, and coordination numbers for CuTi glasses in as-quenched and hydrogen-loaded states as given. The presence of crystalline phases in the structure factor is also indicated.

First-neighbour distances (A) and coordination number

from RDF Crystalline

Sample rl r ; n l phases present

2.66 -

2.63 2.97 2.64 3.14

2.66 - 2.68 3.16 2.62 not resolved 2.62 -3.2 2.69 3.21

2.64 not resolved 2.62 not resolved 2.59 not resolved 2.62 not resolved 2.59 not resolved 2.60 not resolved

11.7 - 12 - 12

11.9 - 12 - 12 - 12 - 12

12.5 - 12 - 12 - 12 - 12 - 12

- c u c u

- c u c u Cu a-Ti/TiHo.5 -

Hjidrogen absorption in C u Ti metallic glasses 2075

I

D - L 2.0 c?

c Y - O -

1.0 2 c U -

3 L c

* 0 -

3.0 I I I I I I I I

-

-

I I I I I I 1 I 8 10 2 4 6 810 1012 4 6

S c a t t e r i n g v e c t o r Q ( X - ' )

Figure 4. The structure factor for Cu0.66Ti0.34 glass with hydrogen for 6 h ( C L I O , ~ ~ T ~ O ~ . , ) H O . ~ ~ (curve A) is not reversibly changed by driving off the hydrogen in a heat treatment of 30 min at 80 "C (curve B).

of the ribbon. However, we understand that the appearance of the ribbon can apparently be preserved, and a degree of reversibility achieved by loading directly with hydrogen gas under controlled conditions (R C Bowman Jr, private communication).

3.2. Changes in real-space structure

The curves which describe the real-space structures of the glasses-the RDF 4nr2p(r) and the reduced RDF G ( r ) = 4nr(p(r ) -Po)-are much more strongly influenced by resolution effects than the structure factors. This arises because of the finite limit of the Fourier transform. The Q,, value for our x-ray data - 17.4 A-' leads to a characteristic width Ar- 0.36 A. An interesting effect is nevertheless observed with these CuTi-H glasses. Comparison of figures 1 and 2 shows that most of the details in the S(Q) curves of the as- quenched state are lost on loading with hydrogen. Intuitively this might suggest that more disordered glassy structures are produced as S(Q) becomes more 'liquid-like'. However, far from this being the case, the first-neighbour peak in the RDF of the loaded glass is split into components, as will be shown below, and the corresponding neighbour distances r l and ri are if anything better resolved. Unfortunately the coordination numbers are less accessible in the loaded state as, owing to peak broadening, the first minimum in the RDF fills in. We suggest that the general form of these RDF curves for the loaded glasses provides valuable evidence of the efficiency with which the Fourier transform can retrieve structural information from S ( Q ) even with apparently poorly defined input material.

The RDFS for the as-quenched structures shown in figure 5 are consistent with the findings of our earlier x-ray and neutron work (Sakata et a1 1982), both in terms of radial distances and coordination numbers (see table 2). The first-neighbour peak in the figures is a combination of unresolved Cu-Cu, Cu-Ti and Ti-Ti correlations, and can be compared with a first-neighbour distance found from the Goldschmidt diameters of the constituents dcu = 2.55 A, dTi = 2.93 A. The average values can be calculated either from the weighting terms of the partial structure factors, or by taking into account the established tendency to chemical short-range order in these glasses (Sakata et a1 1980, 1982). The calculated first- neighbour distances 2.72 A, 2.70 A, 2.67 A for the 50/50, 58/42 and 66/34 glasses, respectively, are only some 3% greater than the observed values in table 2. The coordination numbers, which are subject to a k0.5 atom error, are close to 12 for the three glasses. The packing fraction v = $ n ( r , ) 3 p based on the first-neighbour distance ( r , ) from the RDFS is calculated to lie between 0.65 and 0.67, lower than for the BCC structure and is

2076 B Grieta et a1

0 2 4 6 8 10 R a d i a l d i s t a n c e r (11

Figure 5. The radial distribution functions of CuTi glasses in the as-quenched condition derived from the structure factors shown in figure 1: A, Cuo.soTioso; B, Cuo,ssTio,42; c. CUo.66Tio.34.

100

N

i !=

2 0

0 2 4 6 8 10 R a d i a l d i s t a n c e r (11

Figure 6. The radial distribution functions for CuTi glasses after absorption of hydrogen, derived from the structure factors shown in figure 2: A, (Cuo.soTio.so)Ho.,,; B, ( C ~ O . S E T ~ O . ~ Z ) H O . S ~ ; c , (CUo.66Tio.34)Ho.47.

Hydrogen absorption in Cu Ti metallic glasses 2077

probably a reasonable value for these glasses having a 15% size difference between the constituents.

The RDFS of the hydrogen-loaded glasses are very different from those of the as- quenched glasses. The most significant feature is a splitting and broadening of the first-neighbour peak, which is particularly pronounced for the ( C u , , Ti0.s )H0,77 and (Cuo.s8Ti0.42)H0.59 specimens-see figure 6. In addition the splitting of the second peak is altered and the shoulder enhanced, while there are also small contributions to the RDF at the foot of the first peak. The radial distances r , and r ; for the first peak are given in table 2. In terms of the Goldschmidt diameters, it is found that while the first component corresponds quite closely to the first peak in the as-quenched state, the second component, most clearly resolved for the (Cuo,,oTio,,o)Ho,77 glass, is at 2.97 A only 2% greater than the calculated Ti-Ti distance. The contributions to the RDF at small radial distances are at r - 1.4 A and r - 1.8 A. The latter distance is close to that expected for Ti-H correlations, but the weighting factors for all the hydrogen-metal correlations are very small, of order 5%. It is unlikely that these contributions relate to real interatomic correlations, especially since such features may arise artificially, as shown in the appendix.

The changes in real-space structure caused by loading increased quantities of hydrogen follow broadly the trends for the three alloys given above. Figure 7 shows the RDF curves for the Cuo,,8Tio,,8 glass with the loadings as specified in table 1. Figure 8 gives similar information in terms of the reduced RDF G(r) for the Cu0.66Ti0,34 glass, over a wider range of loading times, although for these specimens the first-peak splitting is less pronounced. The RDF curves show that the first-peak splitting and broadening remains, and possibly increases once established. The fact that the structure changes are long range and not confined to the first-neighbour shell is confirmed by the changes which occur in both the

0 2 4 6 8 10 R a d i a l d i s t a n c e r (1)

Figure 7. The changes in radial distribution functions for Cu0.~gTio,42 glass caused by the absorption of increasing amounts of hydrogen: A, C U O . ~ ~ T ~ O . ~ ~ ; B, (C~o,5gTi0.~2)H0.~~; c, (Cuo.ssTio.a)H0.67; D, ( C U O . S ~ T ~ O . ~ ~ ) H O . ~ ~ .

2018 B Grieta et ai

second and third peaks in the RDF and G(r) curves. Similar changes have also been observed for both PdZr-H glasses (Kai et a1 1982-who also first observed the split$rst peak in the RDF) and in ZrNi-D glasses (Suzuki el a1 1984). These changes appear to be quite irregular, except that there is a sharpening-up of the real-space structure at the highest hydrogen loadings achieved in our samples (curve D in figure 7 and curve E in figure 8).

4. Discussion

Absorbed hydrogen, even when present in metallic samples in considerable concentrations, is likely to remain undetected in most conventional x-ray diffraction experiments. Furthermore, for TM-TM glasses like CuTi where there is little contrast between the species when using x-rays, the structure factor contains information about only topological as opposed to chemical aspects of the glassy structure. Consequently, we are chiefly concerned with geometrical changes to the glass structures caused by hydrogen loading in this study. Within these constraints our results can be compared with both bulk

Hydrogen absorption in Cu Ti metallic glasses 2079

measurements on similar CuTi specimens and structural work on other metallic glass-hydrogen systems.

In parenthesis, neutron-diffraction experiments on these CuTi samples could, of course, provide important complementary information to the x-ray results. We have produced deuterium-loaded specimens using 0.05 M solutions of D2 SO4 in D2 0 but contaminations by small quantities of hydrogen have so far made the neutron experiments unworkable, because hydrogen nuclei have an exceptionally high cross section (8 1 b) for incoherent neutron scattering.

The most comprehensive bulk measurements on CuTi-H glasses, combining calorimetry, qualitative x-ray diffraction and nuclear magnetic resonance, have been made by Maeland and his co-workers (Maeland 1978, Maeland et a1 1980, Bowman and Maeland 1981, Bowman et a1 1982). Our x-ray observations are in broad agreement with the macroscopic models of hydrogen absorption developed by Maeland and colleagues. We confirm that TiCu glasses can absorb considerable quantities of hydrogen and we have demonstrated directly that the glassy structures remain intact to form amorphous hydrides. Maeland (1978) has proposed that moderate heating (e.g. 150-200 "C, as in the course of DTA experiments) causes the decomposition of these amorphous hydrides and the formation of crystalline TiH2 and free copper:

a-(TiCu)H + TiH, + Cu.

We observe this transition in the 150 "C annealed sample, and note that the characteristic temperatures for the appearance of these crystalline phases are significantly below the respective crystallisation temperatures of the as-quenched CuTi glasses.

Loading the CU0.66Ti0.34 glass with increasing amounts of hydrogen (figure 3) causes a similar gradual growth of Bragg peaks associated with metallic copper in the S(Q) curve. We suppose this behaviour to be due to a similar kind of dissociation mechanism, but one occurring at room temperature and possibly stimulated by the saturation of the interstices of the amorphous structure by hydrogen atoms. The Bragg peaks of the TiH, phase have not been detected in this case. Generally their visibility will be less than one-fifth the visibility of the copper peaks, because of the weaker x-ray scattering from the lighter elements. The single extra (i.e. non-copper) peak for the Cuo,6aTi0,34 sample at the highest loading (D in figure 3) is not at a position appropriate to the TiH2 structure. Whilst it is not feasible to identify the crystalline phase on the basis of a single peak, its position is consistent with the presence of a tetragonal intermediate hydride of the kind described by Jaffe (1 956).

It has been suggested (Maeland et a1 1980) that non-homogeneous distributions of hydrogen can occur when loading CuTi glasses electrolytically, with maximum amounts of absorbed hydrogen at the surface and unreacted glass at the centre of the ribbon. In such a case the copper peaks appearing in figure 3 would presumably originate from copper deposited at the surface of the ribbon from a reaction such as that given above, rather than from material distributed uniformly throughout the bulk. It is impossible to make any definite statement on this kind of inhomogeneity from the x-ray experiments in which several ribbon thicknesses, surface and bulk, are penetrated. (The ribbons are comparable in thickness to the filters normally used to remove Kg radiation from the beam.) On balance we favour a uniform distribution of hydrogen and therefore segregated copper for the following reasons. There is no significant difference in the ribbon surface for long or short electrolysis times-those which give rise to extra copper peaks and those which do not. A pronounced inhomogeneity of the hydrogen distribution does not seem to be

B Grieta et a1

consistent with the rapid hydrogen diffusion which has been observed in the glasses (Bowman et a1 1982). Thirdly the characteristic changes in S(Q) lead to well defined changes in the RDF, which can be specifically identified. This kind of behaviour is not what would be expected if a range of disordered structures existed through the ribbon sample, from surface to centre as a result of an inhomogeneous hydrogen distribution. In addition the structural changes observed in electrolytically loaded CuTi ribbons are in agreement with similar investigations of other metallic glass specimens loaded with hydrogen gas directly (Kai et a1 1982), and for which the experimental conditions and possible inhomogeneities are likely to be different. With hindsight, it appears that CuTi glasses are not the best candidates for the present kind of investigation on account of the dissociation reactions of the kind referred to above. Nevertheless the results of our study are in broad agreement with both bulk measurements on similar CuTi-H specimens and diffraction work on other hydrogen-loaded metallic glasses.

Finally, there remains the possibility that CuTi glasses may be unstable with respect to gases other than hydrogen. During the preparation of samples for the present experiments we have detected crystalline impurities in ribbons prepared some time ago ( t > lo7 s) and having compositions towards the edges of the glass-forming range rather than in the equiatomic region. The stability of any metallic glass alloy can be expressed as a lifetime of the glassy state at room temperature, which may in principle be estimated from the activation energy determined from crystallisation studies at elevated temperatures. For some transition-metal glasses changes are observed on a much shorter timescale than predicted and these changes probably occur as a result of the greater affinities for gaseous atoms shown by one or both constituents in this type of glass. CuZr glasses are the best known example, in which surface changes have an important influence on electrical (Gallagher and Greig 1982, Gallagher et a1 1983) and other properties. The changes are thought to be associated with oxide formation (ZrO,) and the consequent removal of zirconium atoms from the solid solution which destabilises the glass and leads to the deposition of free copper which discolours the ribbon surface (Calvayrac et a1 1980). Although some CuTi glasses can also show surface changes of this kind with age they are less marked. The crystalline phases detected in old ribbons appear to be based on CuTi intermetallic compounds. It is possible that some subtle surface decomposition of the glassy alloys can occur stimulated by gaseous absorption and/or oxidisation, but it is difficult to make an accurate assessment in view of the multitude of intermetallic compounds reported in the four known CuTi phase diagrams and the relative invisibility of oxide and hydride phases in x-ray diffraction.

5. Conclusions

Three copper-titanium metallic glasses in the equiatomic range, and including the glass CU66Ti34 in which chemical short-range order between the species Cu and Ti is most highly developed (Sakata et a1 1982), have been loaded with hydrogen electrolytically in a systematic way. In the first instance amorphous hydrides are produced as suggested by Maeland (1978), and at high loadings of hydrogen a dissociation reaction takes place giving rise to crystalline TiH2 and free copper. Characteristic changes are observed in the structure factors S(Q) of the glasses on hydrogen loading which causes long-range variations in the glassy structures. Specifically there is a broadening and splitting of the first-neighbour peak in the RDF. This finding confirms the similar observations made on PdZr-H (Kai et a1 1982) and ZrNi-D (Suzuki et a1 1984) metallic glasses by x-ray and

Hydrogen absorption in CuTi metallic glasses 208 1

neutron diffraction. A discussion of the hydrogen absorption in relation to atomic-scale models of metallic glass structures will be presented in a later paper.

Acknowledgments

This work has arisen from a programme of study on metallic alloy glasses supported in personnel, equipment and materials by the Science and Engineering Research Council. The authors would like to thank Dr M D Archer and also staff at the Department of Metallurgy and Materials Science, University of Cambridge for assistance with the chemical analysis.

Appendix. Apparent structural changes caused by systematic modifications to S(Q)

X-ray experiments on metallic glasses may be performed fairly routinely with the commercial equipment currently available, but there is nevertheless a need for careful attention to experimental detail. As an example we have discussed the use of spacer shims in conjunction with the conventional sample holders for x-ray diffractometers (Dini et a1 1982b). These shims ensure that the metallic glass ribbon is in the correct plane of the diffractometer and that the peaks in the scattered intensity distribution and hence in the RDF are in their correct positions. It was observed at the start of these studies that the electrolytically loaded metallic glass ribbon became buckled. We have tried to minimise the effects of this by careful preparation of the x-ray specimens, but it seems likely that two problems can arise. Firstly, the ribbon may not all be in the correct plane of the diffractometer and secondly, at low angles the geometry of the scattering may be changed giving rise to longer absorption paths and a consequent reduction of observed intensity.

One main effect of hydrogen loading is the reduction in height of the first peak in S(Q)-which is inevitably at a fairly small scattering angle. We have therefore sought to imitate the possible effects on excessive x-ray absorption at low angles by systematically

1 I I I , '7 L c n 1

-1

I I c c L I 1 1 1 I I I I

0 2 4 6 8 10 1 2 1 4 16 18

Scat te r ing v e c t o r U tX- ' )

Figure A.l. The structure factor S(Q) for a Pd80Si20 glass obtained using MO K, x-radiation is given. The first peak is shown systematically reduced to 90%, 80% and 70% of its actual value and the rest of the curve remains unchanged.

2082 B Grieta et a1

0 2 4 6 8

Radia l distance r 1x1

Figure A.2. The reduced RDFS G(r) for PdsoSiZo glass obtained from the Fourier transform of the experimental and the modified structure factors of figure A.l are given. The arrows show the direction of change in G(r) as the first peak in S(Q) is reduced.

reducing the height of the first peak in S(Q) to 90%, 80% and 70% of its actual value, while keeping the rest of the curve constant. Figure A.l shows this process for S(Q) of a Pd80Si20 glass. In Figure A.2 the equivalent G(r) curves resulting from the Fourier transform are shown. The curves are not labelled individually but the main result of decreasing the first peak in S(Q) is to cause a reduction in the amplitudes of the oscillations in G(r). However, in detail two small, but important, effects occur. Firstly the second-peak shoulder becomes relatively more pronounced as the amplitude of the whole peak reduces. Secondly a broad feature emerges at the foot of the first-neighbour peak, at an interatomic spacing - 1.8 A which would in fact be appropriate for Pd-H correlations! Both of these features, which have been produced artificially, are observed for real hydrogen-loaded specimens as discussed above. Their existence in the simulated data points to the need for continuous monitoring of such experimental uncertainties, the effects of which might be confused with the real effects under investigation.

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