Interaction of energetic oxygen ions with lithium-containing amorphous hydrogenated carbon: an in-situ photoelectron spectroscopy study

ELSEVIER

Journal of Nuclear Materials 228 (1996) 290-301

materials

Interaction of energetic oxygen ions with lithium-containing amorphous hydrogenated carbon: an in-situ photoelectron spectroscopy study J.-U. Thiele *, P. Oelhafen Institut ~ r Physik, Universitlit Basel, Klingelbergstr. 82, Ch-4056 Basel, Switzerland

Received 8 September 1995; accepted 21 December 1995

Abstract

Pure and Li-containing amorphous hydrogenated carbon films were deposited using an ion beam deposition technique. Depending on Li content and deposition conditions, differences in chemical structure are detected. Samples were irradiated at room temperature with energetic oxygen ions of 100 eV to 1 keV. Transient time during which layers retain oxygen is significantly longer for Li-containing as compared to pure amorphous hydrogenated carbon films. Oxygen uptake is found to be mainly determined by Li content. For lower contents (CLi < 15 at.%) oxygen uptake is only about 12 at.%. For intermediate Li contents (15 at.% < CLi < 50 at.%) oxygen uptake increases and saturates above Li concentrations of about 30 at.%. In dependence on Li content, and chemical structure of the films, and irradiation conditions varying amounts of Li carbonate like, Li2CO3, and Li oxide like, Li20, compounds are detected in the irradiated samples. At room temperature about 35 at.% oxygen are retained at the surface of these films. For high initial Li contents (CLi > 50 at.%) carbon is almost completely removed upon irradiation and the resulting composition of the films is found to be close to the stochiometry of Li oxide.

1. Introduction

Wall conditioning of the plasma facing walls of controlled fusion test devices decisively influences many plasma parameters and can in particular drastically decrease the oxygen contamination level of the plasma. A well established and actually widely used technique in tokamaks is the 'boronization' of the walls, i.e., the deposition of pure boron or boron-carbon layers. Various techniques have been used, but the deposition of such coatings by means of plasma assisted chemical vapour deposition (PACVD) has proved its suitability as an easy to handle deposition technique. Detailed investigations on the oxidation behaviour and the stability of such coatings have been performed both in tokamaks and in simulation experiments [1-3]. A further improvement of these

* Corresponding author. Tel.: +41-61 267 3720; fax: +41-61 267 3784; e-mail: [email protected].

favourable properties can be expected by the use of carbon films containing the even lighter and more reactive element lithium instead of boron. First promising results have been obtained in model experiments [4,5], by lithium pellet injection on the tokamak device TFTR [6] and by lithium evaporation in the tokamak device JIPP T-IIU [4]. Significant reduction of carbon and oxygen impurities in the plasma was achieved. However, the underlying mechanisms of lithium based wall conditioning have not yet been fully understood. In Ref. [5] we reported on lithium-containing amorphous hydrogenated carbon (a-C : H / L i ) films prepared by radio frequency PACVD and their chemical behaviour upon exposure to an oxygen plasma. In the present work more detailed investigations on lithium-containing films are presented. In order to get better defined deposition parameters, films have been deposited here by Ion Beam Deposition (IBD) combined with lithium evaporation from dispensers. Particular attention shall be paid to the question how conditions of film preparation influence the chemical

0022-3115/96/$15.00 © 1996 Elsevier Science B.V. All rights reserved SSDI 0022-3115(96)00008-6

J.-U. Thiele, P. Oelhafen/Journal of Nuclear Materials 228 (1996) 290-301 structure and composition of the films and how this is correlated to the chemical behaviour and oxygen uptake upon irradiation with energetic oxygen ions. Therefore, interaction of these fihns with energetic oxygen ions has been studied by X-ray photoelectron spectroscopy (XPS) and additionally has been monitored by a quartz crystal microbalance (QCM).

2. Experimental Samples were deposited in the preparation chamber of the photoelectron spectrometer. A broad beam Kaufrnanntype DC ion source (Iontech Inc.) using pure methane as process gas and dispensers for the evaporation of Li (SAES getters) were used simultaneously. The process gas pressure was set at l.. × 10 -3 mbar and kept constant throughout deposition. The resulting composition of the films could be controlled by the ion current of the hydrocarbons and the heating current applied to the lithium dispensers, respectively. Films were deposited at the same time onto silicon substrates and a QCM (Leybold Inficon Inc.). In two sets of experiments the samples were mounted either on a heatable or a transferable sample holder. Films ranging from pure a-C:H to a-C:H containing up to 60 at.% lithium have thus been prepared. Oxygen contaminations in the as-prepared films could be kept below 0.2 at.% for pure and in the range of 1 to 5 at.% for lithium-containing films.

I

5.5

I

I

4.5

In the same deposition/preparation chamber, samples were irradiated with energetic oxygen ions, produced by a filament-free Penning-type ion source. Oxygen pressure was chosen to 1 × 10 -4 mbar, gas flow to 1 seem, ion energy was varied from 0.1 to 1 keV. Ion currents from both ion sources could be measured by a Faraday cup mounted close to the sample position on the tip of the transfer rod. Assuming the ion beam to contain only O~--ions, typical oxygen ion fluxes were of the order of 1014 O~- ions s -1 cm -2. Deposition and etching rates were controlled by means of a QCM film thickness monitor that could be moved close to the sample position. Typical deposition times were 10 min, whereby film thicknesses of about 5 nm have been achieved. Stage one graphite intercalational compound, C6Li, was chosen as a crystalline reference sample. The sample was prepared following the procedure described in detail by Pfluger et al. [7]. The sample's surface was cleaned by in situ cleaving until the wellknown bright golden surface of C6Li emerged [7,8]. In order to compare the results obtained here to those of our previous studies also some samples were prepared by PACVD following the procedure described in Ref. [5]. Photoelectron spectroscopy measurements were performed in situ on a combined XPS/UPS/EELS spectrometer (Leybold, EA 11/100). Only XPS core level spectra are presented here. Photoelectrons have been excited by Mg K,~ (hv = 1253.6 eV) X-ray radiation• Energy resolutions of about 1.0 eV have been obtained, core line shifts

I

a)

XPS (hv = 1253.6eV) C ls, a-C:H ~ Tsubs~,,,=25°c ] |

291

I

I

I

XPS (hv = 1253.6e V ) / ~ 2.5 Cls o /-',% T "mra'C:H/Li (30 at. )/o)/ / \ \

Eion= 1 0 0

"substrate= ~ ° ~

It

I

I

b)

~

m

& 3.5 1.5 .,.~ 2.5 1.5 0.5 0.5

.-. . . . . . . . . . . . . . . . . --L'-:J. . . . . ,--I

290

I

I

"" .....

I

288 286 284 binding energy [eV]

= ~= = - a

I

282

I

I

I

I

I

I

I

I

I

I

290

288 286 284 bindingenergy [eV]

282

Fig. 1. XPS C ls core level spectra of (a) pure and (b) lithium-containing(34 at.%) a-C: H deposited with an hydrocarbonion energyof 100 eV onto room temperature substrates. Background (dashed line) has been subtracted and spectra have been fit by one (a-C:H) and two (a-C : H/Li) Doniach Sunji6functions(dotted lines) accordingto descriptionin the text, respectively.

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are measured with an accuracy of 0.1 eV. Core electron binding energies are referred to the Fermi level of an Au test sample and are calibrated with respect to the Au 4f7/2 core line for which a binding energy of 84.0 eV is assumed. In the following the course of the experiments shall be shortly described. Samples were deposited in parallel onto the substrate mounted on the sample holder and the QCM in the preparation chamber of the electron spectrometer. After deposition, samples were transferred into the measuring chamber and photoemission spectra of the as-prepared state were taken. In the case of the samples mounted onto the transferable sample holder, following the initial investigation, samples were transferred via the deposition/preparation chamber into a third little storage chamber. While there, the layer deposited onto the QCM was irradiated with oxygen ions. From the course of frequency of the QCM the time dependence of the etching behaviour of the respective sample could be read. Afterwards the stored sample was irradiated in several steps for characteristic times as monitored by the QCM. Succesively they were each time reinvestigated by photoelectron spectroscopy. In the case of samples mounted onto the heater on the transfer rod and of the reference samples irradiation of samples and QCM had to be done in parallel. Samples were irradiated until the steady-state regime of the irradiation process was reached, as indicated by the QCM. Subsequently these samples were also investigated by photoelectron spectroscopy.

3. Results

3.1. As prepared samples Besides the mere elemental composition a detailed analysis of the XPS data supplies information on the

chemical structure of the a - C : H / L i . Main information about the chemical structure and the chemical reactions of the samples investigated here can be obtained from the C ls core level peak. In Fig. 1 the C ls peak of a pure a - C : H film and a lithium-containing film, both deposited onto a substrate at room temperature with a hydrocarbon ion energy of 100 eV, are shown. Obviously the peak of the lithium-containing Film is broadened and a shoulder can be seen on the low binding energy side. In order to quantify and understand these differences we shall try to describe the experimental data by a theoretical line shape. A slight asymmetry of the peak prevents the description of the core level line of pure a-C : H by a single Lorentzian profile. Good correspondence of numerical fitting with experimental data can be achieved either by deconvoluting the C ls peak into two or more Lorentzian functions or by additionally accounting for asymmetry. This can be done by convolving a Gaussian profile for the line itself with a Doniach Sunji6 (DS) function, i.e., a Lorentzian profile with an additional parameter, the singularity index a, accounting for asymmetry [9]. The background shown in Fig. 1 is obtained by convolving the DS line with the loss function and subtracting a linear background (for details see Ref. [10]). In the fitting procedure all parameters were free adjustable. Fitting parameters for the C ls core level lines shown in Fig. 1 are given among others in Table 1. In the model developed by Doniach and Sunji6 for simple metals, asymmetry of the peaks is correlated to a density of states at the Fermi level [9]. This model was successfully adapted for the line shape observed in the photoemission spectra of graphite [10]. Detailed understanding of atomic and electronic structure of amorphous hydrogen-free and hydrogenated carbon has been evolved over the last years (see e.g. Refs. [11,12]), but the correlation between structural properties - i.e., ratio of sp 2- to spa-coordinated carbon, hydrogen content or clustering of sp2-coordinated regions - and the features of the C ls core level peak as observed in photoemission is not yet really understood.

Table 1 Fitting parameters obtained by numerically fitting the XPS C ls core level lines of a pure a-C:H film, a carbide containing and a non-carbide-containing a-C : H/Li film to one or two Gaussian-broadened Doniach Sunji6 functions, respectively. All parameters were left free adjustable. All films were deposited onto substrates at room temperature. For the a-C: H/Li film deposited with an ion energy of 100 eVa sactisfactory description of the C ls core level line can only be achieved by decovolution into two Gaussian broadened Doniach Stlnji6 functions. See also Fig. 1 Sample

Peak

Binding energy (eV)

Gaussian width (eV)

Lorentzian width (eV)

Singularity index

Residual

a-C : H, Eion = 100 eV

a-C : H

284.5

1.26

0.43

0.013

1.31

a-C : H/Li, (32 at.%) Eion = 100 eV

a-C: H carbide

285.5 284.1

1.49 1.50

0.99 0.01

0.025 0.009

1.36

a-C : H/Li, (20 at.%) Eion = 200 eV

a-C: H

285.1

1.65

0.59

0.099

1.50

a-C : H/Li, (10 at.%) Eion = 200 eV

a-C : H

285.1

1.36

0.62

0.036

1.45

J.-U. Thiele, P. Oelhafen/Journal of Nuclear Materials 228 (1996) 290-301 F

'

I

'

I

'

I

'

I

CPS (hv = 1253.6 eV), C ls I a.-C:H(/Li), Eion = 100 eV

.

I

I

.

.

.

I

I

I

b) after 0 2 + irradiation, Ech+= 100 eV

a) as prepared

.

I

293

.

:5 25°C

~

50°C

A

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:

i

292

I

i

290

I

I

I

I

288

286

284

282

292

:

:

290

~

=

=

288

......

286

284

282

binding energy [eV]

binding energy [eV]

Fig. 2. (a) XPS C ls core level spectra of pure a-C: H deposited onto a room temperature substrate and a-C : H / L i films deposited onto substrates of different temperatures as indicated. Hydrocarbon ion energy for deposition was 100 eV for all films. (b) XPS C ls core level spectra of the samples she.wn in (a) after irradiation with 100 eV O~ ions for times sufficient to reach the steady-state as determined by the QCM. The components of the curves as obtained by fitting to two Doniach ,~unji6 functions are displayed as dashed curves. Where no dashed curves are displayed fitting to one component was sufficient. All curves are normalized arbitrarily to the same height. For some of the lithium-containing samples the C l s core level spectra cannot be described by a single DS line. As indicated by the dashed curves in Fig. l b these spectra can be decovoluted into two DS lines. Again all fitting parameters were left free adjustable. For the main peak at

However, in order to keep the picture simple we use the DS model function for our further analysis, but one should be well aware that in the case of amorphous carbon no physical meaning can be attributed to the fitting parameters. '

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Eion = 75 eV,

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|

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b) after 02 + irradiation, Eo2+= 100 eV

XPS (hv =1253.6 eV), C ls a-C:H/Li, Tsubarate= 25°C

a) as prepared

~"',.~

100 ov,

200 eV,

500 eV, 12 at.°/oLi I--

292

,

I

290

~ ,

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288

,

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286

,

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284

binding energy [eV]

,

I

282

,

292

I

I

I

I

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290

288

286

284

282

binding energy [eV]

Fig. 3. (a) XPS C ls core level spectra of a - C : H / L i films deposited onto room temperature substrates using different hydrocarbon ion energies as indicated. (b) XPS C ls core level spectra of the samples shown in (a) after irradiation with 100 eV Of ions (for explanation see Fig. 2).

294

J.-U. Thiele, P. Oelhafen/Journal of Nuclear Materials 228 (1996) 290-301

higher binding energy, broadening of both Gaussian and Lorentzian line width as compared to the values obtained for pure a-C :H is found, whereas asymmetry remains almost unchanged (see Table 1). In accordance with previous results on plasma deposited films [13] this peak shifts steadily towards higher binding energy in dependence of the lithium content of the films. Additionally, the line shape reveals a close similarity to the C ls line of pure a-C:H. Therefore, we assign this component of the peak to an a - C : H phase. For the lower binding energy peak about the same Gaussian width as for the main peak is obtained, but for Lorentzian width and asymmetry different values were obtained, as can be seen in Table 1. Following the interpretation of XPS data of a - C : H containing other carbide-forming metals as, e.g., reported for tantalum [14] or chromium [15], we assign this lower binding energy component to the formation of a carbidic lithium-carbon phase. In a - C : H / L i samples prepared by plasma CVD such a carbidic component was always found, whose contribution to the overall carbon concentration depended solely on the lithium content of the samples [13]. In contrast, in the present ion beam deposited samples the ratio of carbidic to a-C : H-related carbon is always lower. It can b e influenced by varying particular deposition parameters as well as by the lithium content. At a given ion energy, e.g., 100 eV like in Fig. 2a, the carbide content of the resulting films can steadily be changed from about 25% of the overall carbon at room temperature to almost 0% at 400°C substrate temperature. The lithium content of this particular film deposited at 400°C substrate temperature was still as high as 20 at.%, a concentration where at lower substrate temperatures and all the more in CVD deposited films a large component of carbidic carbon is observed [13]. Keeping on the other hand substrate temperature at room temperature and varying ion energy for deposition reveals a maximum carbide formation at about 75 eV ion energy, decreasing drastically at ion energies above 100 eV as can be seen in Fig. 3a. Apparently, depending on deposition parameters up to 20 at.% of lithium can be incorporated into the films without forming carbidic bonds to the carbon. However, also in these films a shift of the C ls core level line in dependence of the lithium content is observed (see Ref. [13]). Differences in chemical bonding are displayed only in the XPS C ls core level line, whereas the Li ls core level line does not change much at all. A binding energy of about 56.5 ___0.2 eV was observed for all films, irrespective of their lithium content or chemical structure. For comparison, binding energies of 54.9 eV for metallic lithium and 57.1 eV for lithium in C6Li are reported by Wertheim et al. [8], respectively. In summary, discussion of chemical bonding in a-C : H / L i in terms of absolute core level binding energies is difficult due to shifts in dependence of the lithium content. These shifts might be partly caused by charge transfer from lithium to the carbon matrix [13]. However,

information on the chemical structure can be derived from the analysis of the C Is core level line shape. Whereas for deposition with an ion energy in the range of 100 eV on a room temperature substrate the formation of a carbide phase is indicated by a second component in the C ls core level spectra, raising either substrate temperature or ion energy obviously prevents the formation of a carbide phase in a-C: H / L i deposited by IBD.

3.2. 0~- reactions on pure and lithium-containing carbon 3.2.1. Time dependence In Fig. 4a a typical dependence of QCM frequency upon irradiation with oxygen ions of 100 eV energy is shown for a pure and a lithium-containing a - C : H film, respectively. The same behaviour was qualitatively observed for all investigated films in the range of irradiation parameters used here. In particular, varying the oxygen ion energy between 100 eV and 1 keV does not result in different oxygen saturation concentrations at the surface of the samples. Only processes become faster with increasing ion energy, but since ion energy and flux of the ion source used here cannot be varied over a wide range independently these effects can hardly be distinguished. So for all investigations presented in the following sections we chose the ion energy as 100 eV and ion fluxes as approximately 1 X 1014 O f - i o n s S-1 cm -2. Whereas in the case of pure a - C : H after a transient time of only about 100 s the weight starts to decrease steadily, obviously in the case of the lithium-containing

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irradiationtime [s] Fig. 4. (a) Frequency change of the quartz crystal microbalance coated with a pure and a lithium-containing (34.8 at.%) a-C:H film upon irradiation with 100 eV O 3 ions, respectively. A decrease of frequency herein means a gain of mass. (b) Composition of the lithium-containing film from (a) as calculated from the

area ratios of the XPS peaks. Spectra were taken after times of irradiation with 100 eV Of ions as indicated.

J.-U. Thiele, P. Oelhafen/Journal of Nuclear Materials 228 (1996) 290-301 ~.

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60

Fig. 5. From the frequency-time curves similar to the one shown in Fig. 4 the (a) transient time until oxygen saturation is reached, the (b) erosion, and mass erosion rates (A = area of the quartz = 1 cm2) were determined. Solid lines are guide lines to the eye.

films the transient time until a steady state is reached is much longer. The duration of this transient regime depends on the lithium content of the films as shown in Fig. 5a. From the slope of the QCM frequency f in the steady-state regime erosion rates A f l A t were determined, and mass erosion rates A m ~ A t were ascertained according to the

I

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I

I

XPS (hv = 1253.6 eV) a-C:H/Li (34.8 at.%)

I

basic relationship between QCM frequency and mass given by Lu [16]. For lower oxygen ion energies, i.e., 100 to 500 eV, erosion rates of the lithium-containing films are found to be lower by a factor of about 5 as compared to pure a-C:H, as shown in Fig. 5b. At higher ion energies this difference in erosion rates is reduced to a factor of about 2. As shown in Fig. 4 for a lithium-containing sample, good coincidence between the time dependence of erosion as monitored by the QCM and the composition of the surface as observed by XPS is obtained. A series of XPS core level spectra of this particular sample taken at characteristic times is shown in Fig. 6a. Upon irradiation with oxygen the carbidic component of the C ls peak vanishes and at the same time a peak at 291 eV binding energy emerges. Upon further irradiation both components are shifted by about 0.5 eV towards lower binding energy. This shift might be due to changes in the work function upon oxygen irradiation. But while for the as-prepared samples the position of the main peak in the C ls core level line depended on the lithium content of the films (see above), for the oxygen irradiated samples the main peak was always found at 284.6 ___0.2 eV. Therefore the work function of all irradiated samples can be assumed to be constant within this limit, and binding energies of the core levels in the irradiated samples can contribute information on differences in the chemical environment. The peak at 290.5 eV is also seen in Fig. 2b and 3b in the spectra of the carbide containing films. By comparison to spectra reported by Contarini and Rabalais [17] we assign this peak to a carbonate-like configuration, i.e., Li2CO 3 or LiHCO 3. This change in chemical bonding is reflected also by changes in the binding energies of lithium

b;Lil's/

I

60

d

58 •

I

c)Ols

Os 290

288

'

after irradiation with O2+-ions

7

292

295

286

284

binding energy [eV]

282

536

56 54 binding energy [eV] '

~

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52

50

I

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!-

~

534

532

~-~-~"'~--

550

a

528

526

binding energy [eV]

Fig. 6. XPS core level spectra of the a-C: H/Li (34.8 at.%) film used in Fig. 4, taken after indicated times of irradiation with 100 eV O~ ions. (a) C ls core level peak, (b) Li Is core level peak and (c) O is core level peak. In (a) and (b) curves are normalized arbitrarily to the same height.

J.-U. Thiele, P. Oelhafen/Journal of Nuclear Materials 228 (1996) 290-301

296

and oxygen as shown in Figs. 6b and c. Especially the change in binding energy of the Li ls core level supports our conclusion that a carbonate-like configuration is formed whereas in the as-prepared film a binding energy of 56.4 eV is observed it changes upon irradiation to 55.2 eV (Fig. 7). For comparison, in Ref. [17] a binding energy of 55.1 eV for Li in Li2CO 3 is reported. In the case of pure a-C: H the transient regime is much shorter and accordingly also in the XPS spectra the steady state is reached after rather short irradiation. However, as can be seen in the uppermost curves in Figs. 2a and 2b, the C Is core level peak does not change much at all. Only a slight broadening due to an increase of the shoulder at the higher binding energy side of the peak can be observed, that might be caused by the formation of carbon oxides, CO and CO 2. But oxygen concentration saturates at a value of about 12 at.%. All pure a-C: H films investigated here revealed qualitatively the same behaviour and oxygen uptake, irrespective of the particular choice of deposition and irradiation parameters.

3.2.2. Influence of chemical composition and structure Following the arguments of the previous section it seems near at hand to connect the evolution of the carbonate-related component to the previous existence of the carbidic phase in the as-prepared samples. One might then suspect the oxygen uptake of the samples to be also connected to their carbide content. As discussed in Section 3.1, the formation of the carbidic phase could be influenced by deposition parameters. The C ls core level spectra of these samples are shown in Figs. 2b and 3b after exposure to oxygen ions of 100 eV energy for times long enough for the steady-state regime of erosion to be reached, respectively. At first sight these two figures seem to support the idea of the carbide content being the relevant parameter, but a more detailed look at the changes in elemental composition and chemical configuration upon irradiation reveals a somewhat more complicated picture. In dependence of the lithium content of the films, three types of films can roughly be distinguished.

:> 56.5

f

i

i

I

I

I

10

20

i

i

i

56.0 .~ 55.5 55.0

+ I

30 40 Cu~im [at.%]

I

I

50

60

Fig. 7. Binding energy of the Li Is core level peak of various a-C: H/Li films versus lithium content of the films after irradiation with 100 eV 0~ ions.

0 10 20 30 16[- i i i i

5J

Clithium [at. %] 40 50 i i

60 t o

!, {

0o

2 "

0 0.0

0.5

1.0

1.5

Clithium/CcarbonI initial Fig. 8. Ratio of lithium to carbon content of as-prepared and irradiated samples. The solid line is a guide line to the eye.

For films with low initial lithium (i.e., below 15 at.%) and carbide content, the ratio of Li to C remains almost constant upon irradiation, as can be seen in Fig. 8. The binding energy of the Li ls core level changes only slightly from 56.6 eV in the as prepared films to 56 eV after irradiation, indicating either bonding in a Li20-like configuration or chemical inactivity of at least part of the lithium. Almost no carbonate can be detected after irradiation. Oxygen uptake amounts to only about 12 at.%, the same value obtained upon irradiation of pure a-C : H. In contrast, for films with higher initial lithium and carbide contents, lithium is always enriched upon irradiation as compared to carbon. For intermediate initial lithium concentrations, i.e., about 20 to 40 at.%, oxygen uptake upon irradiation reaches values of typically 30-35 at.%. The binding energy of the Li ls core level shifts upon irradiation to 55.2 eV, close to 55.1 eV binding energy reported for Li in Li2CO 3 [17]. The C ls core level spectra reveal an enrichment of the relative carbonate content as compared to the initial relative carbide content. This might be due either to preferential etching of the a - C : H as compared to the carbide/carbonate component or to the formation of the carbonate being independent of the initial carbide concentration. In the latter case the formation of the carbonate would rather be dependent on a sufficient lithium concentration than on the previous existence of the carbide phase. We shall return to this question later. For lithium-rich films, i.e., films containing about 60 at.% Li, upon irradiation the carbon is almost completely removed. For these films as well as for the graphite intercalationai compound, C6Li, discussed in some detail in the following section, after irradiation a binding energy of about 56 eV is observed for the Li ls core level peak, indicating bonding in a Li20-like configuration. Like in the films of intermediate lithium content oxygen uptake reaches values of typically about 34 at.%.

3".-U. Thiele, P. Oelhafen/Journal of Nuclear Materials 228 (1996) 290-301

XPS(hv = 1253.6eV) C6Li,C ls

"*" " ,

j

as-prepared

t %

after irradiation with 02+-ions ~ ' = 100,:V

d ~t.

•.,

k.___

.~"'~ ~. :

.

::."-~.. .,:.~.¢. "~.

gard to this aspect the stage one graphite intercalational compound, C6Li , was chosen as a reference sample. In this configuration all the lithium is situated between the graphite planes and no lithium is found in a carbide conguration, i.e., covalently bonded to carbon. The XPS C Is core level spectra of the freshly cleaved samples, shown in Fig. 9, correspond well with those reported by Wertheim et al. [8]. After irradiating the sample for 2 h with oxygen ions of 100 eV energy only a few atomic percent of carbon were detected on the surface. Like in a-C: H / L i a carbonate related component of the C ls core level line can be observed. As compared to the spectra of a - C : H / L i this line is shifted by about 0.8 eV towards higher binding energy. Whether this is due to different work functions in a - C : H / L i and C6Li or to the presence of hydrogen in a-C: H / L i cannot be decided on the base of the data presented here. However, oxygen uptake at the surface of C6Li is about 32 at.%, yielding an overall composition close to lithium oxide, Li20, with only small contaminations of carbon. This oxygen uptake is rather high as compared to a - C : H / L i containing the same amount of lithium, indicating an additional influence of the structure of the carbon matrix.

/. j

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..~

;.

..-t.::~."~'"

.~,~,~ :..-....

- . : , ~ . % , .¢..¢, ~ u I

1

I

I

I

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294 292 290 288 286 284 282 bindingenergy[eV]

'280

Fig. 9. XPS C ls core level spectra of graphite intercalational compound, C6Li, before and after irradiation with 100 eV O~ ions. Curves are normalized arbitrarily to the same height. In summary, chemical reactions in the a - C : H / L i sampies show upon irradiation distinct differences in dependence of their initial lithium content, in particular, oxygen uptake is increased only by incorporation of lithium concentrations above 15 at.%. We shall discuss this in more detail in Section 4.

3.2.3. Influence of parameters of irradiation In a previous study a - C : H / L i films prepared by PACVD were exposed to an oxygen plasma of 300 V bias voltage. Films containing about 30 to 40 at.% lithium revealed a distinctly higher oxygen uptake than observed here. The resulting composition was close to that of lithium carbonate, Li2CO3, i.e., above 50 at.% oxygen were retained at the surface. In order to correlate the results of the

3.2.3. Crystalline reference C6Li Regarding the results of the previous sections one might suspect some features of the interaction of the oxygen ions and the a-C : H / L i surface to be related to the formation of the carbidic lithium-carbon phase. With re-

/

prepared by P A - C V D ~

Z. . . .o2 _t . O2 -ion irradiation

/

\

I

I

____./ --I

297

Ib) a-CI:H/Lii(20at/%) I A I preparedbyIBD / ~

'

1 /

t

',---4 oxygen

A

uptake

v/'N

~ + = 100 eV /~ oxygen atsampletemperaturesof/tltk uptake

t

......

X,,..ya,.%

I

I

I

I

294 292 290 288 286 284 282 280 binding energy [eV]

,

294 292 290 288 286 284 282 280 binding energy [eV]

Fig. 10. XPS C ls core level spectra of (a) an a-C : H/Li film prepared by PA CVD exposed succesively to 100 eV O~ ions and oxygen plasma with a substrate hias of 100 eV and (b) an a-C: H/Li film prepared by IBD exposed at different substrate temperatures to 100 eV O~ ions.

298

J.-U. Thiele, P. Oelhafen/Journal of Nuclear Materials 228 (1996) 290-301

present work to those of this previous study again a few samples were prepared by plasma CVD according to Ref. [5]. Irradiating CVD prepared samples of intermediate lithium contents (20 to 40 at.%) with oxygen ions of 100 eV energy from the Penning source results in an oxygen uptake of only about 32 at.%, in the same range as observed for the IBD prepared films investigated here. In contrast, subsequent exposure to an oxygen plasma again resulted in an oxygen uptake of about 48 at.%, comparably high as those reported in Ref. [5]. Additionally, a shift of the carbonate related component of the XPS C ls core level line can be seen in Fig. 10a. It is now near at hand to ascribe the higher oxygen concentrations found in the previous study to different conditions of irradiation in oxygen plasma and oxygen ion beam, respectively. As mentioned above, upon varying ion energy in the range available here no changes in oxygen uptake and chemical structure are detectable. So, main parameters are presumably ion flux and sample temperature, which both are somewhat higher in the case of the plasma. In order to distinguish effects due to changes in these two parameters IBD deposited sample were on the one hand irradiated at room temperature with varying ion fluxes and on the other hand with aequivalent fluxes at different substrate temperatures. The flux range available with the Penning source used for oxygen irradiation is rather small, i.e., between 0.1 and 1.50~- ions s -1 cm -2. However, no flux dependence of the chemical structure or composition of the irradiated samples could be observed in this range, but of course it cannot be excluded for distinctly higher fluxes. More significant changes were found upon heating the sample during irradiation to different temperatures up to 750°C. Still oxygen concentrations do not exceed 40 at.%, but for temperatures above 500°C the contribution of the carbonate component to the C ls core level peak is somewhat higher as compared to irradiation at room temperature. Furthermore, the position of the carbonate related component of the XPS C ls core level line shifts upon increasing the sample temperature during irradiation towards higher binding energy in a similar way as observed above for interaction of a - C : H / L i with oxygen plasma and C6Li with oxygen ions. In accordance with investigations of oxygen retention and reemission of various carbon and boron-carbon layers [18,19] it can be assumed that sample temperature is an important parameter determining the interaction of oxygen ions and a-C : H / L i films.

4. Discussion

4.1. As-prepared samples Comparing incorporation of different metals into a-C: H two general types can be distinguished. On the one side there are metals such as chromium or tantalum known from inorganic chemistry to form carbides [20]. When

brought together with hydrocarbons in a non-equilibrium deposition process - such as sputtering of metal targets in the presence of hydrocarbon ions [15] - they tend to form mainly carbidic clusters. Only excess carbon or metal is present in a non-carbidic form [14]. On the other side, non carbide-forming metals such as gold tend to form metallic dusters of varying sizes embedded in a carbon matrix [15]. Additionally, lithium can occupy an interstitial position in the carbon network, like, e.g., in intercalational lithium graphite or lithium implanted diamond [21]. As reported in Ref. [13] a - C : H / L i films prepared by CVD revealed properties very similar to a-C:H containing carbide-forming metals in so far that the carbide content depended strictly on the lithium content. Already for very low lithium concentrations a carbidic component was detected in the XPS C ls core level spectra. In contrast, in Figs. 2a and 3a it was shown that by changing particular deposition parameters of the IBD deposition films of similar elemental composition but different chemical structure can be deposited. Apparently up to 20 at.% of lithium can be incorporated into the films without forming carbidic bonds to the carbon detectable by XPS. The general resemblance of the C ls core level line in these non-carbide-containing films to the line observed in pure a-C:H gives evidence for lithium being incorporated into an otherwise scarcely changed a-C:H matrix. For the carbide-containing films the picture is somewhat more complicated. The shift of the a-C : H related component of the C ls core line on the one hand and the existence of a carbidic phase on the other hand suggest that lithium is incorporated in different carbon environments [13]. However, the picture of these processes is still incomplete. In order to come to a detailed understanding of the underlying processes further investigations using various experimental techniques will be necessary.

4.2.1. Time dependence When pure graphite or a-C:H are exposed to energetic oxygen ions the implanted oxygen is found to be partly trapped and partly reemitted in the form of CO and CO 2 [18]. In a simple model one can assume that oxygen is captured at its implantation depth until a saturation concentration is reached and excess oxygen starts being reemitted. To give a complete picture one has to account also for the surface sputtering and changes in the samples' chemical composition [22]~ Then a transient time can be calculated after which the material is saturated from the surface down to the implantation depth. This transient time in general depends on ion energy, flux and target temperature. Further irradiation results in a quasi steady-state, where the net rate of oxygen implantation and reemission or film erosion are equal, respectively. The form in which reemission takes place can either be the implanted species itself, as would certainly be the case for noble gases, or any chemical product of the implanted and the target material able to reach the samples surface. As mentioned above in

J.-U. Thiele, P. Oelhafen/Journal of Nuclear Materials 228 (1996) 290-301 the case of pure graphite and a-C:H solely CO and CO 2 are re-emitted. In the case of lithium-containinga-C:H it might be expected that - in analogy to the effects of energetic oxygen ions on boron-carbon materials [18] besides CO and CO 2 different forms of lithium oxides and perhaps carbonates will be reemitted from the films. This shall be investigated in more detail elsewhere [23]. Monitoring such a process in terms of film weight should consequently reveal an initial regime in which the weight remains constant or even increases slightly and a subsequent steady decrease. Such behaviour has been reported for pure a-C and a-C: H [24] and was also observed here for pure as well as for lithium-containinga-C : H. In order to explain the different time scales for pure and lithium-containingfilms different possibilities have to be taken into account. One possibility is the ability to retain oxygen, which is obviously larger in the case of lithium-containing films due to the formation of solid lithium oxide and carbonate in addition to volatile CO and CO 2. Another possibility might be deviation from the simple model that chemical reactions are restricted to the depth where oxygen is implanted into the material. This model is certainly a good approximation for pure a-C:H and a-C : H / L i containing moderate amounts of lithium. In this case only a rather thin layer at the surface is affected by the oxygen ions that can be removed again by sputtering [5]. In contrast, it is well known that even massive pieces of metallic lithium when brought into air oxidize completely. This should certainly result in a further prolongation of the transient regime additionally to prolongation due to enhanced retention. In order to understand the delayed emergence of the carbonate peak in Fig. 6a, i.e., only after 1700 s of irradiation, one has to take into account the ratio of implantation depth of the oxygen ions to escape depth of the photoelecrons. Using the TRIM code for a sample of a composition like the one shown in Fig. 6, assuming additionally a hydrogen content of 25 at.% in order to obtain a realistic density, one tinds an implantation depth of about 1.5 nm for oxygen ions of 100 eV ion energy. For pure a-C : H with an hydrogen content of 25 at.% an implantation depth of about 1.2 nm is obtained. For comparison, the escape depth of C Is core level photoelectrons excited wlth 1253.6 eV photons is about 1 nm in pure graphite and certainly lower in amorphous carbon. Therefore assuming a narrow Gaussian profile for the implantation depth distribution of the oxygen ions it can be concluded from the delayed emergence of the carbonate peak that formation of the carbonate does not take place directly at the samples surface but somewhere below, presumably at the depth of implantation of the oxygen ions. On the other hand this means, once the steady-state regime is reached the composition of the surface as determined from XPS peak areas can be assumed to be constant within the information depth of the method. As mentioned above this can be confirmed experimentally by varying the energy of the

299

impinging ions and thus their implantation depth. No changes in samples composition or chemical structure were observed upon increasing the ion energy from 100 eV to 1 keV. In particular oxygen saturation concentration remains constant. In Ref. [18] Refke et al. reported saturation fluences of 1 × 1016 cm -2 ions for graphite irradiated with oxygen ions of 1 keV energy. Typical ion flux in our experiments was about 1014 ions cm -2 s -1. Taking into account the different implantation depths for ions of different energy, for pure a-C: H irradiated with 100 eV oxygen ions a transient time of the order of 100 s can be estimated, well in the range of times observed. As shown in Fig. 5a the transient regime is significantly extended for a - C : H / L i like for the boron containing graphite in Ref. [18]. In conclusion, the QCM monitor gives a good measure for the relevant time scale for changes in the XPS spectra upon irradiation we shall discuss in the following sections. 4.2. Chemical composition It should be noted right beforehand that a simple understanding as in form of a single chemical reaction equation is not to be expected due to the complexity of the system and the variety of possible chemical reactions in processes far from equilibrium as the ones investigated here. However, some general trends and correlations can be derived. The main question herein is, which deposition parameters resulting in which initial composition are preferable in order to obtain both a maximum oxygen uptake at the surface and high stability of the reacted surface. Therefore we shall try to derive some clues to the question, by which processes the oxygen uptake at the surface is determined. In the previous section we had noticed that in some samples the formation of a carbide was observed and that it was transformed to a carbonate by irradiation with oxygen ions. However, as can be seen in Fig. 11a the oxygen concentration changes more or less regardless of the carbonate concentration. It can anyway be concluded that the formation of a carbonate-like phase is connected to the achievement of higher oxygen concentrations in a way that concomittant to a high oxygen uptake always the formation of some carbonate is observed. However, regarding Fig. l l b one has to assume, that the oxygen uptake of the films is decisively influenced by their lithium content. As already mentioned in Section 3.2.2 three regions can be distinguished in dependence of the lithium content. For films with low initial lithium content (i.e. below 15 at.%) the effects of irradiation with oxygen ions on the films do not differ significantly from those on pure a-C : H. In particular no formation of a carbonate-like phase is observed and the oxygen uptake is only about 12 at.%. Exceeding a lithium content of about 15 at.% the oxygen uptake rises steadily in dependence of the lithium content to about 33 at.% at 30 at.% initial lithium content.

300

J.-U. Thiele, P. Oelhafen /Journal of Nuclear Materials 228 (1996) 290-301 I

30

a)

25 20 15 10

I

I

I

1

2

3

I

I

I

4 5 Ccarbonat e [at.%]

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6

7

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0.0 I'

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1.5

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''-

-/---/7:~--

.

•

20 15

"~

10 ~ 0

10

20

•1 I 30 40 cli~um [at.%]

f 50

I 60

Fig. 11. Correlations determiningthe compositionof a-C:H/Li films prepared by IBD after irradiationwith 100 eV Of ions. (a) Oxygen content versus carbonate content• (b) Oxygen content vs lithiumcontent and ratio of lithiumto carbon content (set off by 5 at.% for clarity). Solid lines are guide lines to the eye.

Beyond this lithium to carbon ratio of about L i / C = 0.5 in the as-prepared films the oxygen uptake cannot further be incerased by increasing the lithium content. Only a nearsurface region approximately down to implantation depth is affected by irradiation with oxygen ions. In the XPS C ls core level spectra evidence for the formation of lithium oxide, Li20, as well as lithium carbonate, Li2CO3, is found. By heating samples of this composition to above 500°C during irradiation the oxygen uptake can be altered to about 40 at.% and a shift of the chemical reaction path towards enhanced formation of lithium carbonate is observed. In the case of lithium rich films, i.e., with an initial composition of 60 at.% Li, irradiation with oxygen ions results in an almost complete removal of carbon and a final composition of 64 at.% lithium and 34 at.% oxygen, close to the stochiometry of lithium oxide, Li20. Taking into account the prolongation above proportionality of the transient regime during which the films retain oxygen it has to be assumed that chemical reactivity of these films is not restricted to the depth of implantation of the oxygen ions. This is in reasonable correspondence with the results obtained under 'real tokamak conditions' in the tokamak device JIPP T-IIU, where the total amount of gettered

oxygen atoms was found to be about one half of the previously evaporated lithium atoms, thus suggesting the formation of Li20 [4]. Remarkable differences are found in the behaviour of a-C : H / L i films and intercalational lithium graphite, C6Li. While the initial composition of C6Li lies in the range of films of low or intermediate lithium content, it reveals upon irradiation with oxygen ions properties much closer to lithium rich films with a content of about 60 at.% lithium. In particular, carbon is almost completely removed from its surface, the final composition being close to the stoichiometry of lithium oxide, Li20. Apparently lithium content is the determining parameter for the interaction of IBD deposited a - C : H / L i films with oxygen ions, but the structure of the carbon matrix has also to be taken into consideration if it is as different as in the case of a-C : H and graphite. Returning to our key question, which initial composition yields the highest oxygen uptake upon oxygen irradiation we can therefore summarise as follows: In order to improve on the oxygen gettering behaviour of lithium-containing as compared to pure a-C:H fdms a minimum lithium content of about 15 at.% is necessary. In the range of ion energies, ion currents, and sample temperatures available for the present investigations the main reaction for lithium rich films (CLi > 40 at.%) seems to be the formation of lithium oxide• For films containing less lithium, i.e., 15 at.% < CLi < 40 at.%, to some extent also lithium carbonate is formed, but the oxygen uptake does not exceed 40 at.%. In the XPS spectra of lithium rich films the same maximum oxygen uptake is observed. However, the total oxygen uptake of lithium rich films is higher due to a larger implantation depth resulting in an extended transient regime before oxygen saturation of the films is reached. On the other hand, films of intermediate lithium content reveal a controlable stability in a sense that only a surface layer of the films approximately down to the implantation depth is affected by irradiation.

5. Conclusions The interaction of pure and lithium-containing a-C:H films as observed by QCM and XPS has been found to follow the same general trends derived for pure carbon and boron-carbon materials [18,19]. In a transient regime at the beginning of O f irradiation, oxygen is implanted into the films up to a saturation concentration. Depending on the available chemical reaction paths it is retained in form of solid compounds or released in form of volatile species such as CO or CO 2. In the case of lithium-containingfilms the transient time during which the film retains oxygen is significantly longer and the etching rate after oxygen saturation of the surface layer is lower as compared to pure a-C : H. Upon interaction with oxygen ions the solid corn-

J.-U. Thiele, P. Oelhafen /Journal of Nuclear Materials 228 (1996) 290-301 pounds lithium oxide, Li20, and lithium carbonate, Li2CO3, are formed. However, oxygen uptake is mainly determined by the lithium content of the film. In dependence on the lithium content three regions can be distinguished. For lower lithium contents (CLi < 15 at.%) the oxygen uptake is only about 12 at.%, the same as observed for pure a-C:H. For intermediate lithium contents (15 at.% < CLi < 40 at.%) 30-35 at.% oxygen are retained at the surface at room temperature. Relative oxygen uptake saturates for lithium concentrations above 30 at.%, but chemical behaviour and absolute oxygen uptake are different lithium rich films (CLi > 40 at.%). For these films irradiation results in an almost complete removal of carbon and a stochiometry close to that of lithium oxide, Li20. Whereas in films of intermediate lithium content only a surface layer approximately down to implantation depth is affected by irradiation, it can be concluded from the extended transient regime that in lithium rich films chemical reactivity is not restricted to the implantation depth of oxygen ions. This behaviour is in reasonable correspondence with results obtained with lithium evaporation in the tokamak device JIPP T-IIU [4]. Reasonable correspondence was also obtained of the results for films of intermediate lithium content with those of a previous study on interactions of plasma CVD deposited a - C : H / L i films with an oxygen plasma [5]. The higher oxygen uptake observed in the previous study was shown to be due to conditions of irradiation. In conclusion, whi,:h chemical reactions are dominating the interaction of a-C : H / L i layers and oxygen ions will depend on the one ihand on composition, in particular lithium content, and chemical structure of the films, on the other hand on conditions of interaction, in particular sample temperature.

Acknowledgements We would like to thank E. Vietzke and J. Winter from Forschungszentrum Jiilich and P. Favia, C. Hollenstein and R. Pitts from Ecole iPolytechnique F6d6rale de Lausanne for numerous valuable and stimulating discussions. Thanks to V. Thommen for providing the C6Li sample and to G. Francz for careful reading of the manuscript. Financial support by the Swiss Bundesamt fiir Energiewirtschaft and Bnndesamt fftr Bildung nnd Wissenschaft is gratefully acknowledged.

References [1] S. Vep~ek, S. Rambert, M. Heintz¢, F. Mattenberger, M. Jur6ik-Rajman, W. Po~mann, D. Ringer and U. Stiefel, J. Nucl. Mater. 162-164 (1989) 724.

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[2] R. Zehringer, H. Kiinzil and P. Oelhafen, J. Nuci. Mater. 176&177 (1990) 370. [3] J. Winter, H.G. Esser, L. KSnen, H. Reimer, J. yon Seggem, F. Waelbroeck, P. Wienhold, V. Philipps, E. Vietzke, T. Banno, D. Ringer and S. Veprek, J. Nucl. Mater. 162-164 (1989) 713. [4] H. Sugai, H. Toyoda, K. Nakamura, IC Furuta, M. Ohofi, K. Toi, S. Hirokura and K. Sato, J. Nucl. Mater. 220-222 (1995) 254. [5] J.-U. Thiele and P. Oelhafen, J. Nucl. Mater. 220-222 (1995) 1047. [6] J.A. Snipes, E.S. Marmar, J.L. Terry, M.G. Bell, R.V. Budny, K.W. Hill, D.U Jassby, D.K. Mansfield, D.M. Meade, H.K. Park, J.D. Strachan, B.C. Stratton, E.J. Synakowski, G. Taylor, D.N. Ruzic and M. Shaheen, J. Nucl. Mater. 196-198 (1992) 686. [7] P. Pfluger, P. Oelhafen, H.U. Kiinzi, R. Jeker, E. Hauser, K.P. Ackermann, M. Miiller and H.J. Giintherodt, Physica B 99 (1980) 395. [8] G.K. Wertheim, P.M.Th.M. van Attekum and S. Basu, Solid State Commun. 33 (1980) 1127. [9] S. Doniach and M. Sunji6, J. Phys. C 3 (1970) 285. [10] F. Sette, G.K. Wertheim, Y. Ma, G. Meigs, S. Modesti and C.T. Chert, Phys. Rev. B 41 (1990) 9766. [11] P. Oelhafen, D. Ugolini, S. Sehelz and J. Eitle, in: Diamond and Diamond-like Films and Coatings, eds. R.E. Clausing et al. (Plenum, New York, 1991) pp. 377-415. [12] U. Stephan, Th. Frauenheim, P. Blaudeck and G. Jungnickel, Phys. Rev. B 50 (1994) 1489; J. Robertson, Diamond Relat. Mater. 4 (1995) 297; D.A. Drabold, P.A. Fedders and P. Stumm, Phys. Rev. B 49 (1994) 16 415. [13] J.-U. Thiele, P. Oelhafen, Solid State Commun. 93 (1995) 29. [14] M. Grischke, A. Brauer, C. Bcnndorf, R. Memming and F. Thieme, in: P. Koidl and P. Oelhafen, eds., Amorphous Hydrogenated Carbon Films, EMRS Symp. Proc. 17, Editions de Physique (1987) pp. 491-497. [15] P. Oelhafen, P. Gantenbein and R. Gampp, Proc. 13th Syrup. on Optical Materials Technology for Energy Efficiency and Solar Energy Conversion, Freiburg, Germany, 1994. [16] C. Lu, J. Vac. Sci. Technol. 12 (1975) 578. [17] S. Contarini and J.W. Rabalais, J. Electron Spectrosc. Relat. Phenom. 35 (1985) 191. [18] A. Refke, V. Philipps, E. Vietzke, M. Erdweg and J. yon Seggern, J. Nucl. Mater. 212-215 (1994) 1255. [19] E. Vietzke, T. Tanabe, V. Philipps, M. Erdweg and K. Flaskamp, J. Nucl. Mater. 145-147 (1987) 425. [20] L.E. Toth, Transition Metal Carbides and Nitrides (Academic Press, New York, 1971). [21] M. Restle, K. Baruth-Ram, H. Quintel, C. Ronning, H. Hofs~iss, S.G. Jahn and I3. Wahl, Appl. Phys. Lett. 66 (1995) 2733. [22] D.K. Brice and W.R. Wampler, J. Nucl. Mater. 145-147 (1987) 368. [23] J.-U. Thiele, E. Vietzkc and P. Oeihafen, to be published. [24] E.B.D. Bourdon, R.H. Prince, W.D. Morison and R.C. Tennyson, Surf. Coat. Tecimol. 52 (1992) 51.