The preparation of high hydrogen content yttrium silicide carbides with reversible storage potential

Journal of Alloys and Compounds 313 (2000) 95–103

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The preparation of high hydrogen content yttrium silicide carbides with reversible storage potential M.A. Hassen, I.J. McColm* Department of Industrial Technology, University of Bradford, Bradford BD7 1 DP, UK Received 29 May 2000; accepted 15 August 2000

Abstract Carbides of Y 5 Si 3 , of general formula Y 5 Si 3 C x , where x 5 0.05–0.7, have been prepared by arc melting and their reaction with hydrogen studied. Single-phase carbides can only be made melting pre-formed Y 5 Si 3 with carbon. An activation process, which earlier work did not find, has been developed to induce reaction of the carbide with hydrogen. It involves temperature cycling in gas pressures above 2 atm to higher than 5508C with the severity of these conditions depending on the carbon content. Higher T and P are needed to condition samples of increased carbon content. Conditioned samples, where x$0.3 in Y 5 Si 3 C x , will repeatedly cycle 2.5 H (fu)21 between room temperature and 5008C with between 0.9 and 1.2 H (fu)21 being desorbed or absorbed with fast kinetics within a very small temperature range. This behaviour is believed to arise from a polymorphic change between an a-form of the D8 8 structure and a supercell b-form when x,0.3 and between b-form and a b9-form when x.0.3. The first phase change is identified from the rapid expulsion of hydrogen to be close to 4388C and the second at 4468C. Hydrogen saturation of the samples is close to 7 H (fu)21 and is related to the carbon content. Y 5 Si 3 C 0.5 H 7.33 appears to be an upper limit. The large amount of hydrogen in the crystal structure does not result in loss of good crystallinity or loss of hydrogen reactivity that is always observed for samples without carbon, such as Y 5 Si 3 H ¯5 . The significantly enhanced crystal stability and the greatly improved cycle lifetime of the carbide is analysed in terms of a structural disorder caused by the presence of carbon in at least two crystallographic sites. An attempt to quantify the enhanced stability, compared to Y 5 Si 3 , is made by measuring reaction enthalpies using Van’t Hoff isopleths. To achieve this, rigorous conditions were imposed on the reaction of the binary alloy with hydrogen in order to obtain reliable and repeatable values for DHabs and DHdes .  2000 Elsevier Science B.V. All rights reserved. Keywords: Interstitial alloys; Hydrogen storage materials; Gas–solid reaction; Enthalpy; Entropy

1. Introduction An important characteristic of the D8 8 -silicide, M 5 Si 3 , structure is that small atoms such as C, B, O, N and H can be inserted into it without any apparent major structural effect. Of all these non-metal additives, carbon and hydrogen have the greatest effect in terms of structural change and as a result carbides and hydrides have been most studied [1–7]. Mayer and Shidlovsky [5] originally suggested that carbon simply dissolved into Y 5 Si 3 with little structural effect up to composition Y 5 Si 3 C 1.0 . Later, this was shown to be incorrect when Al-Shahery et al. [1] showed that carbon causes complex unit cell changes which became apparent through the appearance of weak super lattice lines on Guinier X-ray films. For yttrium silicide, it was shown

*Corresponding author.

that up to a carbon content of x50.3, there was a decrease in the unit cell volume. At carbon content x50.3, there was a sudden increase in the unit cell volume with the ] a-parameter larger by a factor of Œ3 and an unchanged c-parameter; this was called Type-1 superstructure. The superstructure diffraction lines became more obvious as the carbon content was increased up to x50.5. At carbon content x50.65 single crystal X-ray diffraction patterns showed a lack of order in the planar stacking in the c-axis direction. This effect produces an amorphous appearance in X-ray powder investigations. Further addition of carbon made the structure fully ordered again. A return to the basic structure at x50.9 preceded more dramatic change at x51.0, when a new superstructure, called Type-2, was ] encountered with a ss 5Œ3a, c ss 5 3c. At x51.8, a completely new structure appeared which was orthorhombic. Finally at carbon content x52.0, the structure changed to a so far unresolved variant. Because of these structural changes, the Y 5 Si 3 -carbides

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were felt to be a good system to use in a hydrogen sorption investigation for the following reasons. • Carbon occupies the larger sized octahedral sites and may direct hydrogen to other sites in the structure. • The structural changes associated with increasing carbon content, has in the past, been ascribed to electron donation from the carbon to a delocalised band system [8]. Hydrogen also does this in many binary alloys and so new possibilities for H-uptake may arise near to phase boundaries. • It is relatively easy to control the composition during preparation of the carbides since no gaseous reactants are involved compared to forming nitrides and oxides [9,10]. In an earlier study McColm and Button [4] briefly reported that several single phase Y 5 Si 3 C x alloys reacted with hydrogen at 6508C and 1 atm pressure. The reaction was very slow and reached a limiting composition of 1.7 H (fu)21 after 2 h for Y 5 Si 3 C 0.3 . Higher carbon content alloys absorbed less hydrogen over the same time. Reversibility was not reported. The reaction was not significant at temperatures below 6508C and above this temperature only slow uptake was reported, such that after several days compositions around 8 H (fu)21 were found but it was emphasised that X-ray films showed Y 2 O 3 to be present and so the hydrogen content was uncertain. All the carbohydrides had the Type-2, or as they named it, the b-superstructure. An interesting fact is that hydrogen is known to cause a composition-related structural change in Y 5 Si 3 similar to the carbon effect and this structural change can be reversed if the hydrogen is desorbed [6]. It may, despite the earlier report by Button [4], be possible to make hydrogen react reversibly with one of these carbide phases, by moving it from one structure-type to another if more care is taken in choosing the carbide starting composition. Hence it was decided to re-examine the hydrogen reactivity of the carbides to see if the earlier reported difficulties associated with the reaction could be overcome through some form of activation process. Cyclical applications of pressure and temperature are often used to condition intermetallic materials and so it was decided to investigate this approach. The work was focused on Y 5 Si 3 because of past experience with this material. As the work progressed two slightly different approaches to the preparation of the carbides appeared to have significant effects on the resultant products and their reaction with hydrogen. Carbide products synthesised from the three elements Y, Si and C, behaved differently compared to the product of the reaction of molten, preformed Y 5 Si 3 and carbon. The application of hydrogen gas pressure at temperatures above 6008C was found to be a way to condition the carbides and change their hydrogen reactivity in a significant way.

2. Experimental details Firstly, the carbides were made by arc melting pieces of the required elements in a small argon-arc furnace at a pressure above ambient. Rare Earth Products Ltd. supplied yttrium 99.99%, in sublimed ingot form. Pieces were sheared off using a sharp knife or chisel. Silicon, KochLight Ltd., 99.999%, was used as fractured pieces from a large block. If the weight loss after melting was greater than 0.5% then the sample was not subsequently used. Alloy carbide was prepared from the same components using small pieces from spec-pure carbon rod from Johnson Mathey Chemical Ltd. In this way preparation from powder was avoided in order to restrict possible oxidation. Since each preparation consisted of three or four small pieces, this minimised uptake of impurities from the surface of the materials. Some samples were made by re-melting previously prepared Y 5 Si 3 with small pieces of carbon. Homogeneity was checked by metallographic examina¨ tion and powder XRD analysis. A Hagg-Guinier Camera was used with Cu Ka 1 radiation to check phase composition and measure the lattice parameters. Silicon, lattice parameter a50.543088 nm, was used for calibration. Hydrogenation was monitored thermogravimetrically using two systems based on a Sartorius Model 4406 microbalance. Experimental details were identical with those described in Refs. [4,6,9] except that hydrogen was further purified using an Oxyclear column in the hydrogen entry train.

3. Results

3.1. General observations In this work the carbides were prepared by two slightly different methods which produced significantly different results with respect to phase purity and the resultant reaction with hydrogen. Carbides made by melting premade Y 5 Si 3 with carbon were essentially single phase Y 5 Si 3 Cx up to an x value of 0.7, see Table 1. However, despite efforts to homogenise the melts made by mixing the three elements, the samples were always multiphase, Table 1 Lattice parameters for D8 8 phases formed by the reaction Y 5 Si 3 1 xC→Y 5 Si 3 C x x in Y 5 Si 3 C x

a nm

c nm

0.0 0.1 0.2 0.3 0.5 0.7

0.8417 0.8444 0.8433 1.4637 1.4639 1.4581

0.6351 0.6376 0.6395 0.6419 0.6470 0.6473

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Table 2 Lattice parameters for D8 8 phases made from arc melting the elemental constituents

superstructure, b-phase, decreases as the carbon content increases. For example

x in Y 5 Si 3 C x

a nm

c nm

Other phases

Y 5 Si 3 H 2.43 →Y 5 Si 3 H 1.54 1 0.46H 2 a b

0.1

0.8438

0.6386

0.2

0.8432

0.6419

0.3

1.4581

0.6419

a-SiC, Unknown phase (possibly Y 5 Si 4 ) a-SiC, Y 15 C 19 , Y 5 Si 4 a-SiC, Y 15 C 19 , Y 5 Si 4

containing SiC, Y 15 C 19 , Y 5 Si 3 and Y 5 Si 3 C x 2 D on solidification, see Table 2. With respect to the reaction with hydrogen, there was a difference between multiphase and single phase carbide samples which is seen by comparing the temperatures at which reaction begins in 1 atmosphere of hydrogen after the sample has undergone the conditioning process described below. For samples with the same nominal composition some data are given in Table 3. In general, multiphase samples made from the elements begin to absorb hydrogen at lower temperatures than single-phase materials. Concentrating on the data from samples made from Y 5 Si 3 plus carbon, the single phase material, the effect of carbon on hydrogenation can be more readily seen and summarised as follows. • Reaction start temperature rises as the carbon content increases. • The amount of hydrogen needed to reach the Type-1

While for the carbohydride Y 5 Si 3 C 0.1 H 1.58 →Y 5 Si 3 C 0.1 H 0.93 1 0.32H 2 a b Information on other samples is given in Table 3, where it can also be seen that the amount of hydrogen rapidly desorbed during the first re-heating stage also decreases until the carbon content x50.3 is reached. • At a carbon content of x50.3 hydrogen is absorbed, but not into the basic a-Y 5 Si 3 structure because Y 5 Si 3 C 0.3 already has the Type-1 superstructure as Table 1 shows. Now, when the composition reaches Y 5 Si 3 C 0.3 H 2.84 more hydrogen, around 1.2 H (fu)21 , is rapidly desorbed under the experimental conditions of 1 atmosphere pressure of hydrogen at 4458C. The ability of the b-structure to rapidly cycle about 1 H (fu)21 without involving a phase transformation is now a notable feature. • For carbon contents of x up to 0.3, when a portion of the hydrogen is rapidly desorbed on heating and reabsorbed on cooling it occurs over quite a small temperature range of 442–4698C and is related to the combined content of carbon and hydrogen. The method of sample preparation has little affect on this temperature but does influence the quantity of hydrogen that is reversibly exchanging. An interpretation of this is that the a to b, Type-1 superstructure phase change, is

Table 3 Hydrogen absorption-desorption data for Y 5 Si 3 C x samples made from the elements [A] and from the reaction of pre-formed Y 5 Si 3 1x carbon [B] x in Y 5 Si 3 C x

Reaction start temperature 8C

H absorbed up to 5008C H (fu)21

H content at. room temp H (fu)21

Amount / T desorbed rapidly H (fu)21 / 8C

Amount / T absorbed rapidly H (fu)21 / 8C

A 0 0.1 0.2 0.3 0.4 0.5

381 498 482 458 505 523

1.41 0.67 0.79 0.79 0.68 0.40

2.43 1.32 1.23 1.59 0.88 0.41

0.91 / 469 0.52 / 443 0.47 / 461 0.42 / 445 n.o.a n.o.a

0.91 / 467 0.52 / 412 0.49 / 340 0.42 / 323 n.o.a n.o.a

B 0 0.05 0.1 0.2 0.3 b 0.5 c

381 452 493 501 503 –

1.41 1.24 1.00 1.27 1.37 –

2.43 1.81 1.58 1.58 2.84 2.5

0.91 / 469 0.77 / 447 0.65 / 442 0.31 / 446 1.20 / 446 0.9 / 446

0.91 / 467 0.73 / 447 0.64 / 442 0.28 / 442 1.10 / 446 0.9 / 446

a

n.o. is not observed This sample had a final composition of Y 5 Si 3 C 0.3 H 5.9 after 50 cycles of absorption / desorption. c This sample had a final composition of Y 5 Si 3 C 0.5 H 7.33 after 50 cycles of absorption / desorption. b

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responsible for the fast kinetics and this phase change temperature is only slightly affected by composition. • The effect that increasing the carbon content has on the total hydrogen content of the carbohydride is broadly that increasing the x value in Y 5 Si 3 C x causes a decrease in the amount of hydrogen ejected at the phase transition until x50.3. Then as suggested a different mechanism seems to operate. The atomic size of carbon is such that it should only occupy octahedral sites of which there is 1 (fu)21 . Thus for a composition, Y 5 Si 3 C 0.2 , it is expected that 0.8 H (fu)21 is still available for desorption from such sites. In practice only 0.31 H (fu)21 is desorbed. A suggestion arising from this observation is that the presence of carbon makes it more difficult to fill the active, non-octahedral, hydrogen sites as hydrogen diffuses into the crystal. A series of experiments was done to test this idea and they produced results completely unexpected from previous experience with hydrogenation of the parent Y 5 Si 3 alloy, namely these carbides are stable in hydrogen at high temperature and high pressure in a way that the alloy is not. Exposure of the binary Y 5 Si 3 alloy to pressures of hydrogen above 1 atmosphere and temperatures around 5008C rapidly drives the sample to hydrogen saturation when it becomes completely amorphous. However when Y 5 Si 3 C x phases are heated to 6508C in hydrogen at 2 atm pressure, cooled to room temperature and then reheated to 5008C, sample degradation and saturation does not occur. Even more remarkably, it is found that increased reactivity with hydrogen under milder conditions is induced. In short, the carbides have undergone a conditioning process. This enabled the hydrogen reactivity of two samples to be studied in detail via a repeated absorption–desorption cycle programme. The results from these experiments are described in the next section.

3.2. Absorption–desorption of hydrogen in Y5 Si3 C0.3 – 0.5 samples Heating conditioned Y 5 Si 3 C 0.3 , that is a sample previously taken-up to 6008C in 2 atm pressure of hydrogen, from ambient in 1 atm pressure of hydrogen produced a weight gain at 3688C and by 4358C the sample had a composition Y 5 Si 3 C 0.3 H 0.33 . As the temperature increased to 4388C hydrogen began to desorb. In all previous research involving Y 5 Si 3 , samples containing so little hydrogen would not lose any at this stage. The sample when cooled from 4408C started to absorb hydrogen in three stages rapid

Y 5 Si 3 C 0.3 H 0.3 1 0.085H 2 → Y 5 Si 3 C 0.3 H 0.47 v. fast

Y 5 Si 3 C 0.3 H 0.47 1 0.45H 2 → Y 5 Si 3 C 0.3 H 1.37 slow

Y 5 Si 3 C 0.3 H 1.37 1 0.72H 2 →Y 5 Si 3 C 0.3 H 2.84

(1) (2) (3)

On re-heating there was a slow loss of 0.12 H (fu)21 between 197 and 2458C: slow

Y 5 Si 3 C 0.3 H 2.84 →Y 5 Si 3 C 0.3 H 2.72 1 0.06H 2 Then a very rapid expulsion of hydrogen at 4468C: v. fast

Y 5 Si 3 C 0.3 H 2.72 → Y 5 Si 3 C 0.3 H 1.52 1 0.6H 2 This pattern of absorption and desorption between RT– 5008C–RT was repeated over many cycles during which a gradual build-up of the total hydrogen content occurred and a final composition of Y 5 Si 3 C 0.3 H 5.9 was reached. At this limiting composition the total reversible hydrogen in the sample was constant at 2.5 H (fu)21 in the range RT–5008C. It was still a process in which a rapid desorption of between 0.9 and 1.2 H (fu)21 occurred at 4468C but this was superimposed on a 1.6 H (fu)21 gradual gain or loss over the temperature interval. Similar behaviour was observed for the conditioned Y 5 Si 3 C 0.5 samples but a more severe conditioning programme was needed to activate them, namely heating to 7008C at 2 atm hydrogen pressure before cooling and repeating the process. Once conditioned, Y 5 Si 3 C 0.5 absorbed / desorbed hydrogen between RT and 5008C like the Y 5 Si 3 C 0.3 sample but the final hydrogen content was higher at Y 5 Si 3 C 0.5 H 7.33 . 21 The sample still exchanged between 0.9 and 1.2 H (fu) rapidly at 4468C and this was superimposed on 1.6 H (fu)21 exchanging throughout the temperature interval. To remove more hydrogen than 2.5 H (fu)21 from this and other carbohydrides requires an increase in temperature. However when the temperature was raised to 9008C a significant amount of hydrogen was still present in the solid. Thus a re-melt was tried to see if all hydrogen could be removed this way. Arc melting was easy and the resultant beads had an X-ray powder pattern identical to the parent carbide, Y 5 Si 3 C 0.3 , as Fig. 1 shows, or Y 5 Si 3 C 0.5 , indicating that all the hydrogen had been removed. These samples could be re-conditioned and made to re-absorb hydrogen. From previous experience with D8 8 -structure binary alloys an unexpected feature of the carbohydrides is that samples containing large amounts of hydrogen produce extremely clear X-ray powder films with very sharp diffraction lines. Unlike the binary alloys they have not become amorphous. This is shown in Fig. 1 for Y 5 Si 3 C 0.5 H 7.33 where all the lines are clear and sharp and can be accounted for on a strong Type-1 superstructure unit cell, with a51.4578 nm, c50.6486 nm, plus a few very faint lines, possibly from SiC and two weak broad lines that could be from YH 2 . The observed lattice parameters for Y 5 Si 3 C 0.3 H 5.9 are, a51.4586 nm, c50.6424 nm, which show differences from the alloy carbide in Table 2 where it can be seen that there is an a-parameter contraction and c-parameter expan-

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99

Fig. 1. Guinier focusing X-ray powder photographs of (a) Y 5 Si 3 C 0.5 H 7.33 , (b) parent Y 5 Si 3 C 0.5 and re-melted Y 5 Si 3 C 0.5 H 7.33 .

sion. Both these hydrides are Type-1 superstructure compounds, which excludes an a⇔b phase change in the fast desorption–absorption reaction. However it is too difficult to detect a Type-2 superstructure from anything but single crystal X-ray work which was not available. Hence there is the possibility that the mechanism involves a b⇔b9 transformation.

3.3. Reaction enthalpies The effect of carbon on hydrogen-sorption of Y 5 Si 3 can be quantified by experiments that measure the equilibrium hydrogen pressure as a function of temperature. In order to do this comparatively, the enthalpy of the reaction in Eq. (4) needs to be determined: Y 5 Si 3 H 1.52 1 0.45H 2 ⇔Y 5 Si 3 H 2.43

(4)

However this enthalpy is somewhat uncertain for the following reasons. • In the past, samples with different values for x in Y 5 Si 3 H x have been used to make the initial measurements. • In experiments where the same sample was used for tests at more than two temperatures, the hydrogen content was considerably different as the slow hydrogen up-take occurred. • Different total hydrogen content means that the lattice

parameters are different and this leads to different strain energies operating during the sorption process. These problems mean that DHf for Eq. (4) needed to be established by experiments that come as near as possible to the following idealised requirements. • All experiments should be carried out on hydrides where the hydrogen composition is the same. • The same amount of hydrogen has to be exchanged and that should be close to 0.9 H (fu)21 . • Values obtained for DH by desorption and absorption experiments should be considered separately. In order to come near to meeting these requirements some strict conditions were set when considering the results from a large number of experiments of the Van’t Hoff-type. 1. If uDHabsorption u . uDHdesorption u the data is suspect and was discarded. 2. If x total .4.0, where the crystallinity becomes poor, the results were discarded. 3. If x abs 2 x des . 0.2 the result is discarded. 4. If the temperature at which the sample begins to absorb hydrogen is below 3908C it probably contains free yttrium metal and the result is discarded. 5. If the maximum temperature for the fast desorption .5508C, it is discarded.

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6508C but they did not use any gas pressure above 1 atmosphere. This is consistent with the observations made in this research which demonstrates the need for a conditioning process involving hydrogen gas pressures above 1 atm. Once conditioned, and reversibly reacting with hydrogen, the narrow temperature range over which a part of the hydrogen is rapidly desorbed and re-absorbed, and the speed of the process, is in accord with a crystal phase change playing an important part in the process. When the b→a change occurs the annihilation of some hydrogenfilled sites is the reason for the fast desorption step. The observation that the Type-1 superstructure phase, Y 5 Si 3 C 0.3 , absorbs hydrogen but does not readily desorb it, is positive evidence that the transition between the basic a-D8 8 -structure, and the b, Type-1 superstructure phase, is implicated in the fast desorption of a portion of the hydrogen contained in the structure. Hence for these materials an important part of the mechanism of rapid hydrogen sorption is: a-carbide1yH 2 →b-carbohydride→a-carbohydride1xH 2 T1

Fig. 2. Van’t Hoff plot for Y 5 Si 3 C 0.2 1xH 2 ⇔Y 5 Si 3 C 0.2 H 2x .

The experiments were of two types, one where a single sample was used and equilibrated at a number of temperatures. In this type of experiment x in Y 5 Si 3 H x steadily increased. The second type used a new Y 5 Si 3 sample for each temperature. Good linear plots of ln Peq vs. 1 /T were obtained when data from experiments conforming to the above criteria were considered. On applying these constraints to more than 50 data sets the average values of DHabs 5 243.862.4 kJ mol 21 and 21 DHdes 550.561.9 kJ mol for the rapid exchange of 0.9 21 H (fu) are the most reliable so far obtained for Y 5 Si 3 and it is against these that the effect of carbon in the structure is measured. Fig. 2 shows that carbon in the lattice of Y 5 Si 3 has a significant effect on the appearance of the ln Peq –1 /T plots. Many of the isopleths are curved, particularly for the desorption process. The curvature can be rationalised as two linear regions as Table 4 shows for all compositions in the range of carbon 0.05 to 0.3. At Y 5 Si 3 C 0.5 both absorption and desorption data produce single linear plots like the parent Y 5 Si 3 , as Fig. 3 shows. The enthalpy values obtained from the slopes of all the lines drawn are given in Table 4.

4. Discussion In earlier studies, Button and McColm [4,6] reported little or no reactivity for yttrium silicide carbides up to

T2

where x , y and T 2 . T 1 . This is true at least for carbides up to x50.3 and above x50.3 there is an improvement in the hydriding characteristics with the retention of a fast hydrogen in / out capability of more than 1 H (fu)21 at 4468C. Although it could not be confirmed, a b⇔b9 transformation might also be involved. It is significant that single lines are obtained for ln Peq vs. 1 /T plots for Y 5 Si 3 and Y 5 Si 3 C 0.5 compositions which represent the most stable a and the fully ordered b-phase. Assuming that the fast hydrogen uptake is a function of: a 1 H 2 → b or b 1 H 2 → b9 then the reaction enthalpies for the x50 and x50.5 compositions also involve these phase transformations. The enthalpy values in between these two compositions, which are generally less than 243.8 kJ (mol H 2 )21 for absorption, are presumably influenced by the way the carbon atoms make it easier to achieve the b-structure. From the enthalpy values it can be summarised: a 1 H2 → b

DH 5 2 43.8 kJ (mol H 2 )21

b 1 H 2 → b9

DH 5 2 50.8 kJ (mol H 2 )21

The first of the above reactions represents the a-phase reacting with hydrogen, which is exothermic, and the a-phase transforming to the b-phase which will be an endothermic process. The second reaction is an exothermic reaction of H 2 with the b-structure combined with an endothermic phase change. Energies associated with phase changes are usually small, and so assuming they are of similar magnitude, the difference of 27 kJ (mol H 2 )21

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Table 4 Reaction enthalpy data from Van’t Hoff isopleths for the reaction of hydrogen with Y 5 Si 3 C x samples made by reacting pre-formed Y 5 Si 3 with carbon x in Y 5 Si 3 C x

Shape of Van’t Hoff plots

0

Linear for both (abs) and (des) Curved; two lines intersecting at 5528C (abs)

0.05

DHabs kJ (mol H 2 )21 243.8

50.5

255.9 a

79.2 a

25.1 b 277.3 c

5408C (des) 0.1

215.3 a

Curved; two lines intersecting at 5308C (abs)

0.3

Linear for abs Curved; two lines intersecting at 5218C (des)

211.0

12.3 b 44.1 c 33.5 a

Very curved; three lines intersecting at 4838C (abs) 4768C (abs)

231.6 a

14.2 b 104.3 c 44.7 a

a

Value Value c Value d Value b

223.0 b 2149.7 d 221.4 c

4838C (des) 4738C (des) Linear for both (abs) and (des)

0.5 calculated calculated calculated calculated

for for for for

69.8 b 76.6 c 27.0 a

211.0 b 215.9 c

5338C (des) 0.2

DHdes kJ (mol H 2 )21

250.8

17.2 b 142.2 d 39.5

the best straight line fit using all data points. linear part at high temperatures above stated temperature. linear part below the stated temperature. linear part between the two temperatures.

represents the increased reactivity of the b-polymorph towards hydrogen compared to the a-structure. A similar estimate indicates that the b9 polymorph is some 11 kJ (mol H 2 )21 more reactive. The effect of these increased negative enthalpies is qualitatively seen in the rates of hydrogenation of these phases when Y 5 Si 3 reacts with hydrogen and kinetic measurements are made. These kinetic data will form the content of another paper. At the heating stage, filling of sites in the b-phase is a statistical process depending on site energies and diffusion mechanisms for the hydrogen in the solid and hence time as well as the temperature and pressure of the system is important in determining the composition. This is probably the reason why there is a limited correlation between alloy composition and volume of hydrogen cycled as the temperature is varied across the phase change temperature. Calculations of site energies and their filling will be the subject of a future paper. Another unexpected feature of this research is the finding that hydrogen contents, y . 5.0, in Y 5 Si 3 C x H y , do not result in amorphous samples, such as those obtained when the hydrogen content exceeds y 5 5.0 in Y 5 Si 3 H y . The lines in the X-ray powder diffraction pattern are

extremely sharp and well defined as Fig. 1 shows. The change from crystalline to amorphous material in ternary systems, such as Y 5 Si 3 H x , is usually the result of decomposition reactions, for example, Y 5 Si 3 H x → YH 21 D 1 Y x Si y The amorphatisation in such cases results from ultrafine particle sizes and strain energy. It might be concluded therefore that carbohydrides, such as Y 5 Si 3 C 0.3 H 5.9 are resistant to such decomposition and remain well crystalline to the extent that the faint superlattice lines are easily visible on the X-ray films and the main lines remain sharp and clear. This strongly suggests that carbon stabilises the Y 5 Si 3 structure. However, calculations using a modified Miedema method [10] show that carbon acting as a wholly interstitial component in the octahedral sites makes the alloy structure less stable because the enthalpy of formation becomes more positive. Thus another explanation for the unusual X-ray results must be sought and for this we turn to some earlier neutron spectroscopy work. Some time ago we reported neutron spectroscopy results for a series of hydrides, amongst which were some

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Fig. 3. Van’t Hoff plot for Y 5 Si 3 C 0.5 1xH 2 ⇔Y 5 Si 3 C 0.5 H 2x .

carbohydrides [12]. The carbohydride spectra were completely different to all other hydrides examined in that they gave only low intensity peaks superimposed on a large, structured, background. The implication was made that the material had organised lattice vibrations with H atoms in a very unusual site. It was suggested that such a site could be H substituted into the partially covalent Si–Si sublattice, or H atoms in vacant Y metal sites along with H atoms in some of the alloy octahedral interstitial sites. Significant C–H interaction was indicated by the neutron spectra, a situation which would occur if some of the H was in the Y sublattice and so formed part of the octahedral site coordination environment. The need to get hydrogen in to such unusual sites might be the reason why a high temperature and raised pressure conditioning process is needed to make the carbides begin to react before reversibly absorbing and desorbing part of the hydrogen in the structure. Detailed single crystal X-ray analysis by AlShahery et al. [1] deduced that the carbon, when x.0.05, was present on two sites; one site was the expected 6gI, at the centre of a 6 Y atom octahedron. The other was a 4d site that caused metal atom vacancies in the yttrium sublattice. Thus the inclusion of carbon leads to significant structural disorder which might be the reason why these

carbides show such resistance to hydrogen amorphatisation. This view is based on the findings of Hughes [13] that the substitution of a small part of the nickel in LaNi 5 by tin introduces positional disorder, leading to a decrease of La–Ni interatomic distances and an increase in thermal degradation resistance during hydriding–dehydriding cycles. The carbide Y 5 Si 3 C 0.05 had the highest measured enthalpy of reaction with hydrogen and this too may reflect that after this composition positional disorder in the carbon sites appears. If the disorder introduces shorter Y–Si interatomic distances then greater alloy stability will result. The work referred to above, where a modified Miedema calculation predicted instability as carbon was added to the structure was only based on interstitial carbon atoms in the octahedral sites and could not make predictions for a situation of positional disorder. Now, if positional disorder leads to a more stable carbide, then applying the rule of reversed stability, the enthalpy of formation of the hydride will be less. This is seen in Table 4 where after x50.05 there is a sharp decrease in hydride enthalpy of formation for the suggested positional disordered carbide–hydride phases. Hence the suggestions in the early crystallographic work and the neutron spectroscopy are substantiated by the way the measured hydrogen reaction enthalpies depend on the carbon composition and the observed high temperature stability of the carbohydrides. We will report in another paper the effect that nitrogen has on the D8 8 structure where the positional disorder is not observed and amorphatisation is a general feature. This we offer as further evidence for the unusual behaviour of D8 8 carbides as reported in this paper.

5. Conclusions • Preparation of single phase carbides from a melt is dependent on the starting materials. Pre-formed Y 5 Si 3 re-melted with carbon is successful while melting the three elements is not. • Reversible absorption of hydrogen into Y 5 Si 3 C x requires a conditioning stage involving hydrogen gas pressures above 1 atm at high temperature. As the carbon content increases the pressure–temperature cycle is more severe. • From temperature–pressure equilibria a temperature of 4558C is determined as the a⇔b phase change that is responsible for the rapid sorption of 0.9 H (fu)21 . • When carbon exceeds 0.3 (fu)21 more than 1 H (fu)21 is rapidly exchanged in a process that may involve a b⇔b9 phase transition. Overall some of these carbides can cycle 2.5 H (fu)21 , in three stages, over a 5008C range. • Reversible hydrogen sorption into the D8 8 -structure containing as much hydrogen as the 7.33 H (fu)21 , as reported here, has never been observed before. Other

M. A. Hassen, I. J. McColm / Journal of Alloys and Compounds 313 (2000) 95 – 103

reports [6,9,11] make it clear that hydrogen reactions cease in non-carbide samples when H$5.5 H (fu)21 . The amount of recoverable hydrogen in all the carbides where x$0.3, is relatively constant at 2.5 H (fu)21 over a temperature interval of 5008C of which 0.9–1.2 H (fu)21 are cycled with fast kinetics over a limited temperature range around 4558C. Given this fact, the good stability of the materials and the low densities of the solids, these carbohydrides may have some potential as hydrogen storage materials. • Positional disorder induced by the carbon in the structure is responsible for a significant resistance to thermal degradation and hydrogen amorphatisation.

References [1] G.Y.M. Al-Shahery, D.W. Jones, I.J. McColm, R. Steadman, J. Less-Common Metals 85 (1982) 233.

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[2] G.Y.M. Al-Shahery, D.W. Jones, I.J. McColm, R. Steadman, J. Less-Common Metals 87 (1982) 99. [3] G.Y.M. Al-Shahery, R. Steadman, I.J. McColm, J. Less-Common Metals 92 (1983) 329. [4] T.W. Button, I.J. McColm, J. Less-Common Metals 97 (1984) 237. [5] I. Mayer, I. Shidlovsky, Inorg. Chem. 8 (1969) 1240. [6] T.W. Button, Ph.D. Thesis, University of Bradford, 1982. [7] G.Y.M. Al-Shahery, Ph.D. Thesis, University of Bradford, 1978. [8] T.W. Button, I.J. McColm, in: G.J. McCarthy, J.J. Rhyne, H.B. Silber (Eds.), Rare Earths in Modern Science and Technology, Vol. 1, 1980, pp. 415–422. [9] T.W. Button, I.J. McColm, J. Ward, J. Less-Common Metals 159 (1990) 205. [10] M.A. Hassen, Ph.D. Thesis, University of Bradford, 1997. [11] I.J. McColm, V. Kotroczo, J. Less-Common Metals 131 (1987) 191. [12] S.M. Bennington, I.J. McColm, D.K. Ross, J. Less-Common Metals 172 (1991) 307. [13] J.M. Hughes, Int. J. Hydrogen Energy 22 (1997) 347.