Chemical vapour deposition of tungsten carbides on tantalum and nickel substrates

Thin Solid Films 272 ( 1996)

I 16-I 23

Chemical vapour deposition of tungsten carbides on tantalum and nickel substrates H. Hijgberg, P. Tigtstrijm, J. Lu, U. Jansson Thin Film and Surface Chemistry Group, Department

of Inorganic Chemistry, University of Uppsala. Box 531, S-75121 Uppsalu. Sweden

Received 25 May 199.5; accepted 21 August 1995

Abstract Tungsten carbide films have been deposited by low-pressure chemical vapour deposition from a WF6/C3H,/H2 mixture on Ta and Ni substrates. Single-phase WC films could be deposited on Ta in a broad vapour composition range at 900 “C. A mixture of WC and W,C was deposited in the temperature range 700-850 “C, while an amorphous film was obtained at 650 “C. The temperature behaviour suggests that deposition of carbon is a limiting factor in the growth process. The deposition process on Ta could be separated into two parts: a fast substrate reduction step of WF, leading to the formation of metallic W followed by a slower formation and deposition of WC. The growth behaviour on Ta was also affected by tantalum carbides at the film-substrate interface. A different growth behaviour was observed on Ni. It was found that several q-carbides (i.e. N&W& and N&W&) were formed during a fast initial growth stage. Later on, the v-carbides reacted with carbon under the formation of WC and free Ni particles. It was also found that the carbon deposition rate on Ni substrates was higher than on Ta. This was explained by a catalytic process where Ni particles on the film surface favoured carbon deposition. Keywords; Nickel; Tungsten;

Tantalum

1. Introduction Tungsten carbides have an interesting combination of mechanical, electrical and chemical properties, suggesting a potential use of these compounds in many thin film applications. Chemical vapour deposition (CVD) of tungsten carbide films has been reported by several authors (see, for example, Refs. [ l-41 ) . The most commonly used tungsten precursors have been different metal halide compounds (WF, and WC&), while various hydrocarbons usually have been employed as carbon sources. A survey of the literature shows that most processes investigated so far have yielded tungstenrich films, containing several different phases such as WC, W2C and metallic W [ l-31. An exception to this statement can be found in a paper by Takahashi and Itoh [4]. They were able to deposit single-phase WC from a mixture of WC16/C3H8/H2 at temperatures above 1 200 “C. In a recent study, Tsgtstrijm et al. suggested that the problems in depositing single-phase WC are due to the high reactivity of WF, (or WCl,) compared with most hydrocarbons [ 51. This leads to a preferential deposition of tungsten unless the reaction probability of the halide can be reduced. It was demonstrated that this could be achieved by carrying out the CVD process at very low pressures (0.1 Torr) and with high linear gas flow velocities (7 m s-*) [S]. 0040-6090/96/$15.00 0 1996 Elsevier Science S.A. All rights reserved SSDIOO40-6090(95)08010-4

An important factor in CVD is different types of substratevapour and substrate-film interactions. Examples of such interactions in CVD of carbides are: substrate etching by precursor molecules, formation of substrate carbides, solid solutions between film and substrate, etc. The transition metals exhibit a large variation in chemical properties making them more or less suitable as substrate materials in WC CVD. In principle, it is possible to distinguish between transition metals belonging to groups 4-6 and the metals in groups 810 in the periodic table. This can be illustrated by comparing two model substrates, Ta (group 5) and Ni (group 10) in CVD of WC from a WF,/C,H,/H, mixture. Ta represents a group of metal substrates which, in contrast to Ni, is highly reactive towards WF, [ 61. Typically, these reactions involve a substrate etching by the halide which can strongly affect the growth behaviour during the initial stages of growth. The deposition process can also be influenced by the formation of interfacial substrate carbides and alloys. Also in this case, Ta and Ni exhibit a quite different behaviour. For example, as the other metals in group 46, Ta forms stable binary carbides, while Ni only forms a metastable carbide Ni3C. It should be noted that W and C can form complex ternary carbides (e.g. q-carbides) with several metals in groups 810 (i.e. Fe, Co and Ni) [ 7,8]. No such carbides are formed between W, C and metals in groups 4-6 [ 91. Finally, as an

H. Hiigberg et al. /Thin Solid Films 272 (1996) II&123

alloying element W also exhibits a quite different behaviour towards Ta and Ni. For example, W exhibits a complete miscibility with Ta, while the maximum solid solubility of W in Ni is only about 20 at.% [ 101. W and Ni also form several intermetallic compounds such as N&W and NiW

[Ill. The objective of this study was to investigate CVD of tungsten carbides in a low-pressure CVD system from a reaction gas mixture of WF6, C,Ha and HZ. The process has been studied as a function of vapour composition, temperature and deposition time. In addition the influences of substrateetching and carbide formation have been investigated in more detail on two substrate materials (Ta and Ni)

2. Experimental In a previous study it was shown that the deposition conditions for tungsten carbide CVD could be improved by using very low pressures and high linear gas flow velocities [ 51. For this reason, most experiments were carried out in a very low-pressure CVD reactor (hot-wall type), pumped by a diffusion pump with a capacity of 2000 I s- ‘. The base pressure and leak rate in the reactor were typically 1 X 1O-6 Torr and 5 X lo-’ Torr 1-l s-‘, respectively. This system made it possible to deposit tungsten carbide films under very clean conditions using deposition pressures below 0.5 Torr. A more detailed description of the system has been presented elsewhere [5]. The carbide films were deposited from a reaction gas mixture of WF, (claimed purity, 99.98%), C3HR (claimed purity, 99.95%) and H2 (claimed purity, 99.9999%). Preliminary experiments showed that high H2 concentrations favoured the deposition of metallic tungsten and suppressed the carbide growth. For this reason, all experiments were carried out with an Hz concentration of 50% in the vapour. The p(C,H,) / p(WF,) ratio was varied from 3 to 39 (i.e. a large excess of C3Hx was used in all experiments). Thin foils of Ni (99.999%) and Ta (99.9%) were used as substrate materials. Some deposition experiments were carried out on precarburized Ta substrates consisting of a mixture of TaC and Ta,C. The carburization procedure was performed in the CVD reactor by exposing Ta foils to 25 mTorr C3H, at 1 000 “C during 30 min. Prior to an experiment, the substrates were cleaned in trichloroethane followed by an methanol/ethanol rinse. The substrates were then loaded into the reactor using an Ar flow to minimise contamination. All substrates were placed vertical to the gas how. During heating and cooling, the system was flushed with an H, flow to reduce the oxide formation. The deposition process started by transporting a pre-mixed reaction gas into the reactor through a leak valve. Most experiments were performed at a total pressure of 0.1 Torr and a linear gas flow velocity of 7.5 m so ‘. The deposition temperatures ranged from 650 to 1 000 “C. Some experiments were also carried out in an ultra-high vacuum system (UHV) allowing for both thin film growth and surface anal-

I17

ysis. The UHV system has been described elsewhere [ 121. In this system, surface analysis was performed with X-ray photoelectron spectroscopy (XPS) using Al Ka radiation. Prior to the UHV experiments, the substrates were cleaned in situ by argon ion sputtering to remove surface oxides. The phase composition of the deposited films were determined by X-ray diffraction (XRD) using a conventional K20 geometry (Cu K~z radiation). X-ray fluorescence spectroscopy (XRFS) was used to measure the total amount 01 deposited W atoms per cm’. The tilm thickness can then be obtained by assuming ideal density. The film morphology and the presence of free carbon in the surface regions were investigated by atomic force microscopy ( AFM) and Raman spectroscopy, respectively. Finally, some films were studied with transmission electron microscopy (TEM) . Cross-sections of the samples were prepared by gluing several pieces of the sample with an epoxy glue into a sandwich structure. The sandwich was then encapsulated into a stainless steel tube, cut with a diamond cutter and polished to a thickness of 0.1 mm. Finally, the sample was dimpled to 20 ym and ion-milled at an angle of 4”.

3. Results and discussion Preliminary experiments showed that single-phase WC films as well as mixtures of WC and W,C could be deposited in a broad temperature and vapour composition range. It was also found that the deposition process was strongly affected by, for example, substrate material, total pressure, vapour composition and temperature. All experiments in this study have been carried out at experimental conditions favourable for the deposition of single phase WC in a broad vapour composition range. No attempts have yet been made to optimize the deposition process with respect to deposition rate. 3.1. CVD of tungsten carbides on Ta substrates 3. I. I. lnjhence of vapour composition and temperature Experiments carried out at 900 “C and at a total pressure of 0.1 Torr showed that tungsten-rich films containing a mixture of metallic W and W,C were obtained for low p(C,H,) / p(WF6) ratios in the vapour. However, single-phase WC films could be deposited on Ta substrates in a broad temperature range for p(C,Hx) Ip(WF,) ratios above 10. Fig. 1 shows a typical X-ray diffractogram of a WC film deposited on a pure Ta substrate (no precarburization of Ta) using a p( C,H,) lp( WF,) ratio of 10 and a temperature of 900 “C. As can be seen, the diffractogram only shows peaks from Ta, TaC, Ta,C and WC. No indication of W,C and/or metallic W can be observed. Raman spectroscopy also showed that no free carbon was present in the surface region of the film (see Fig. 2(a) ) . Further studies showed that single-phase WC films could be obtained at 900°C withp( C,H,) lp( WF,) ratios ranging from 10 to 39. Films deposited in this vapour composition range reached a thickness of 2.4 X 10” to

H. Htigberg et al. /Thin Solid Films 272 (1996) 116-123

118

25

35

45

55

65

20 Fig. 1. Diffractogram of WC film deposited on Ta at 900 “C and ap( C,H,) / p(WF,) ratio of 10. +, Ta; V, Ta*C; 0, TaC.

I

1

1

1000

1200

1400

1600

1800

Raman shift [cm-‘] Fig. 2. Raman spectra from (a) WC deposited on Ta, (b) WC/WJ deposited on carburized Ta, and (c) WC deposited on Ni. Deposition temperature 900 “C,p(C3Hs)/p(WF~) = 19. ml

--1

. .‘y+ l

l

10

20

40

PW,H~YP&;

Fig. 3. Amount of deposited W deposited during 1 h on (A) Ta and (0) precarburized Ta. Deposition temperature, 900 “C. Total pressure, 0.1 Torr. The lines in the figure are guidelines for the eye.

1 X 1018 W atoms cm- 2 after a deposition time of 60 min (Fig. 3). This corresponds to a deposition rate of about 0.050.2 pm h-‘. Furthermore, the results in Fig. 3 clearly show that the deposition rate decreased with increasing p( C,Hs) / p(WF,) ratios. The morphology of the WC films was also investigated by AFM. It was found that all films exhibited a nodular morphology which was slightly dependent on the vapour composition (cf. Fig. 4(a) and 4 (b) ) . Furthermore, cross-section TEM clearly showed that the WC films were

continuous and that a thin TaC/Ta,C layer was formed between WC and the Ta substrate (see Fig. 5). Electron diffraction confirmed that this layer mainly consisted of TaC. An interesting observation was that a large amount of pores was formed between Ta and the TaC/Ta,C layer. The high pore density in this region will certainly have a detrimental effect on the film adhesion. It is possible that the pore formation is due to the diffusion behaviour (i.e. Kirkendall diffusion) during carbide formation but further studies are required to confirm this assumption. The influence of temperature on the deposition process was also studied. It was found, however, that the film growth was affected by the presence of substrate carbides. During the initial stages of growth, tantalum carbides (TaC and Ta,C) are formed by diffusion of carbon atoms into the substrate. This process is highly temperature dependent and yields large variations in the surface composition at different deposition temperatures. The temperature experiments were therefore carried out on precarburized Ta substrates with a defined mixture of TaC and Ta,C (for a description of the carburization procedure, see Section 2). A series of films were then deposited on the precarburized substrates in a temperature range from 650 to 1 000 “C, using a p( C,H,) lp( WF6) ratio of 19. It was found that lower deposition temperature favoured the formation of more tungsten-rich films (see Fig. 6). On the precarburized substrate, single-phase WC was only deposited at 1 000 “C, while a mixture of WC and W,C was obtained in the temperature range 700-950 “C. X-ray diffractograms of films deposited at 650 “C showed only a few, very broad peaks indicating that the films deposited at this temperature are amorphous. Complementary XPS analyses showed that the amorphous carbide contained only tungsten and carbon. A rough quantitative analysis with WC as a standard suggest that the C/W ratio in these films was about 0.3-0.35 (i.e. a composition corresponding to the composition of about W,C). The results in Fig. 6 shows that the apparent activation energy for tungsten deposition is about 40 kJ mol- ‘. This value is considerably less than the 67-73 kJ mol- ’ observed for CVD of metallic W from a WF6/H, mixture [ 131. The activation energy for carbon deposition is more difficult to estimate since the carbon content in the film is temperature dependent. However, a rough estimation assuming that the carbon contents at 1 000 and 650 “C are about 50 at.% and 25 at.%, respectively, yields an activation energy for carbon deposition of about 80 kJ mol- ’ (the estimation is based on the assumption that no tungsten and/ or carbon are transported into the substrate during the deposition process). Consequently, the results above suggest that some step in the deposition of carbon limits the growth process.

3.1.2. Infruence ofTa etching on WC CVD In a previous study, Boman and Carlsson showed that Ta is a highly reactive substrate which should be able to reduce WF, according to the following reaction [ 61:

119

H. Hiigberg et al. /Thin Solid Films 272 (1996) 116-123

Fig. 4. AFM micrographs

of WC films deposited on Ta at 900°C with (a) p(C,Hs)Ip(WF,)

Fig. 5 I’hM cross-section

at a WC film deposited on Ta at 900 “C with ap(C?H,)

TW 42.0

,

1300 I

900 I

1100 1

1

67 "0 . ..

n

;? g s n

41.0

= 10 and (b) p(C,H,)/p(WF,)

-

? "a ? 2 3

= 15.7 Deposition time, 60 min

/p( WF,) ratio of 19. Deposition time, 60 min

exposing Ta to WF, at extremely low pressures ( 1 X 10 --’ Torr) in an UHV chamber. Fig. 7 shows an XPS spectrum of a Ta substrate prior to and after 10 s of WF, exposure ( no H2 or hydrocarbon present in the vapour). As can be seen, the Ta 4f peaks in the spectra almost completely disappear after only 10 s of WF6 exposure. The appearance of strong W 4f peaks and a subsequent XRFS analysis show that the Ta surface has been covered by a thin ( lo-20 A), continuous W film. Furthermore, a series of experiments in the low-

l 40.0

3 "'I'

) "I'*

s

Id/r

x n

’

”

1.1

0.9

0.7

”

(K-l)

Fig. 6 Influence of temperature on the deposition rate of W (atoms cm-> h) on precarburized Ta substrates. p(C,Hs)Ip(WF,) = 19. Phase composition of the films: l , single phase WC; v , W,C/WC; and 0, amorphous tungsten carbide. The line in the figure is a guideline for the eye.

SWF,(g)

+6Ta(s)

--SW(s)

+6TaF,(g)

The high reactivity of Ta suggests that the CVD of tungsten carbrde films can be separated into two different steps. A fast substrate-driven deposition of W (or tungsten-rich carbide films) can be expected during the initial stages of growth, when Ta is exposed to the vapour (assuming that C3H8 is more or less inert compared with WF,) . Later on, the deposition process can proceed without any major influence of the substrate material. This growth model was confirmed by

40

30

20

Binding energy (eV) Fig. 7. XPS spectra of the W 4f peaks from (a) a Ta substrate prior to WF, exposure and (b) after 10 s. WF, exposure at lO--5 Ton. Exposure temperature, 400 “C.

H. Higberg

120

et al. /Thin Solid Films 272 (1996) 116-123

deposition rate during the show that the total amount TaC/Ta,C layer after 60 1.2X 10’s C atoms cm-* TaC is formed).

35

25

45

55

65

28 Fig. 8. Diffractogram p(C,H,)Ip(WF,)

of W film deposited

after 3 min on Ta at 900 “C.

= 19. +, Ta; 0, Ta,C; 0, TaC.

1500

2

2 1000 % 2

E

0

zf

so0

0 0

20

40

60

Time (min) Fig. 9. Influence of deposition time on amount of deposited W atoms on (A) Ta and (0) Ni. Deposition temperature, 900 “C. p(C3Hx)l B(WF,) = 19.

pressure CVD reactor confirmed that metallic W is formed during the initial growth stage (see Fig. 8). As can be seen the film consists of mainly metallic W and small amounts of W&J and Ta,C after a deposition time of 3 min. The W film exhibits a strong (200) orientation similar to the texture of the cold-rolled substrate. Moreover, it should be noted that only small amounts of tantalum carbides have been formed. After 15 min deposition time the phase composition has changed. The W peaks in the diffractogram have almost disappeared and WC is now the dominating phase. The relative intensities of the peaks from the tantalum carbides have increased suggesting that the tantalum carbides are formed by diffusion of carbon into the substrate material. Finally, after a deposition time of 1 h, the film consists of only WC without any indication of metallic W and W,C (see Fig. 1) . The change in phase composition also leads to a variation in growth rate (Fig. 9). As can be see, the initial deposition rate of metallic W (measured as amount of deposited W cm - 2, is very high. This is followed by a much lower deposition rate of W atoms as the phase composition is changed to WC. It should be noted, however, that Fig. 9 only shows the amount of deposited W cm-*. The C deposition rate exhibits a quite different behaviour. The phase transformation from W to WC as well as the formation of tantalum carbides require that the deposition rate of C must be higher than the W

later growth stages. Fig. 5 of deposited C in WC and min can be estimated to (assuming that about 120

and 9 in the about nm of

3.1.3. Influence of tantalum carbide formation on WC CVD An interesting observation is that the presence of tantalum carbides could affect the deposition process. This was studied by carrying out deposition experiments on carburized substrates, containing a mixture of TaC and Ta,C on the surface (for details of the carburization procedure, see Section 2). It was found that no metallic W was formed on the precarburized substrates. Short deposition experiments (three minutes exposures) revealed that mainly W& was formed during the initial growth stages. The estimated W deposition rate for the W,C formation was about 30% lower than for the initial deposition of metallic W on the pure Ta substrates. Another observation was that the films deposited on precarburized substrates at low p(C,Ha) /p( WFn) ratios were considerably thicker than films deposited on pure, metallic Ta (Fig. 3). It is conceivable that this is due to a growth mechanism where significant amounts of carbon arc taken from the carburized substrate surface. Furthermore, phase analysis of the films showed that single-phase WC films could only be deposited on precarburized substrates withp( C,H,) / p( WF& ratios higher than 26. At lower ratios the films consisted of a mixture of W& and WC. No free carbon was detected in the films by X-ray diffraction and/or Raman spectroscopy (see Fig. 2(b) ). It is likely that the thickness variations as well as the change in phase composition are due to different nucleation and deposition conditions during the initial stages of growth. Further studies are therefore required to evaluate the correlation between experimental growth parameters and the growth behaviour. 3.2. CVD of tungsten carbides on nickel substrates Continuous WC films could also easily be deposited on Ni substrates at 900 “C (see the cross-section TEM micrograph in Fig. lO( a)). The WC films exhibited a rough morphology with “whisker-like” outgrowths. An interesting observation was that the initial deposition rate of W atoms were higher on Ni than on Ta substrates (see Fig. 9). However, in contrast to Ta, no metallic W was formed during this growth stage. The X-ray diffractograms showed that films deposited for 3 min consisted mainly of the cubic q-phases, N&W& (M,Ctype) and Ni,W,C (M&-type) together with some W,C (see diffractogram (a) in Fig. 11) . No diffraction lines from metallic W, W-Ni compounds, WC and/or the metastable carbide Ni,C could be observed at this point. After 15 min, the phase composition in the film was dramatically changed (diffractogram (b) in Fig. 11) . The relative intensities of the diffraction peaks from the n-phases have now decreased significantly. The diffractogram indicate that W,C is the majority phase and that small amounts of WC have been formed.

H. Hfigberg et al. /Thin Solid Films 272 (19961 116-123

Fig. 10 TEM micrographs of (a) continuous WC film deposited on Ni at 900 “C and ap(C?H,) grain covered with graphite (arrow) on the WC film surface.

Finally, after a deposition time of 60 min, the diffractograms showed peaks from WC and Ni without any traces of W,C and/or ternary carbides (diffractogram (c) in Fig. 11). No indication of any ternary carbides could be observed in TEM of WC films deposited for 60 min. The results above show that WC on Ni is formed by a series of phase transformations: Ni2W,C,Ni,W,C

+ W,C --+WC

A consequence of the phase-transformation sequence is that the chemical composition of the film must change as the deposition proceeds. This can be seen as a slight shift of the diffraction peaks of the N&W& phase indicating an increased Ni content (cf. diffractograms 11 (a) and 11 (b) ) [ 71. Moreover, it is also likely that the Ni will form free metal particles and/or dissolve into the WC phase as the q-carbides decom-

25

30

35

40

45

50

55

20 Fig. 1 I. Diffractograms of films deposited on Ni at 900 “C for (a) 3 min. (b) 15 min and (c) 60 min. Deposition temperature, 900 “C; p(C,Hx) / p( WF‘,) = 19.0, WC; A, Ni; n, W,C; qz4, Ni,W,C; qno. Ni,W,C; +, C.

lp( WF,)

ratio of 19. (b) Ni inclusions in WC grains. (c) Ni

pose. This was also confirmed by XPS analysis showing large amounts of Ni in the WC films deposited for 60 min. A rough quantitative analysis shows that the Ni/W ratio in the surface region was about 0.5 to 0.75. This means that significant amounts of metallic Ni must be present as free particles in the film since the solid solubility of Ni in WC is only about 2 at.% ]7]. The presence of Ni in the films was also confirmed by cross-section EM. It was found that continuous WC films contained metallic Ni grains with an average diameter of 50200 nm (Fig. 10(b) and IO(c)). The Ni grains were found in the bulk as well as on the surface of the lilm. Another observation was that the WC films deposited on Ni contained significant amounts of amorphous carbon and graphite in the lilm. Diffractogram (c) in Fig. 11 showed a peak at a d-value of about 3.34 A which can be attributed to graphite and/or amorphous carbon. XPS analysis also showed significant amounts of “graphitic” carbon on the WC surface. Furthermore, Raman spectroscopy showed two broad peaks centred at 1 360 and I 590 cm ’ (see Fig. 2). The two peaks can be assigned to amorphous carbon in the film [ 141. The presence of graphite in the fihns was also confirmed by cross-section TEM. An interesting observation was that the Ni grains on the hhn surface in general were covered by a thin layer of graphite (Fig. lO( c) ). Moreover, the Ni grains were also frequently found to be starting places for the whisker-like outgrowths of WC (see Fig. lO( c) ). The results above show that the Ni substrates seems to favour the formation of free carbon (amorphous carbon and graphite). In fact, the formation of free carbon shows that the total amount of carbon deposited on the Ni substrates are considerably higher than on Ta using identical experimental conditions. A possible explanation to this deposition behaviour is that no stable substrate carbides are formed on Ni and that the carbon atoms only are used to form WC and free

122

H. Hiigberg et al. /Thin Solid Films 272 (1996) 116-123

carbon on the Ni substrate. However, we would like to propose another explanation based on the presence of Ni grains on the films surface. It is well known that Ni is an active catalysts in many reactions involving hydrocarbons (see e.g. ref. [ 1.51). We therefore suggest that the Ni grains on the surface catalyse the decomposition of C3Hs thereby yielding a higher carbon deposition rate on Ni than on Ta substrates. In fact, it is also likely that the Ni grains can serve as C sinks since the solid solubility of C in Ni at 900 “C is about 1 at.% [ 101. This is supported by the formation of thin graphite layers on the Ni grains (see Fig. 10(c)). Such layers are known to form on Ni substrates exposed to hydrocarbons when the temperature is reduced and the solid solubility of C decreases (see, for example, Ref. [ 161). The presence of Ni grains on the WC film surface can also explain the whiskerlike outgrowths observed in this study. It is well known that Ni substrates yields whiskers in TiC CVD. It has been proposed that the whisker formation is due to a vapour-liquidsolid (VLS) or a vapour-solid (VS) mechanism (see, for example, Refs. [ 17,181). The absence of any low-melting eutectics in the W-Ni-C system suggest that whisker-like structures in our study are due to a VS mechanism.

4. Concluding

remarks

Tungsten carbide films have been deposited by low-pressure CVD from a WF,/C3H,/H, gas mixture using Ta and Ni as substrate materials. It was found that single-phase WC films could be deposited on Ta in the temperature range 9001 000 “C, while lower temperatures (700-850 “C) yielded a mixture of WC and W$Z. Amorphous W-C films with about 25 at.% C could be obtained at 650 “C. Typical WC deposition rates at 900 “C were about 2.5 X 10” to 1.0X 1018 W atoms cm-* h- ‘. This corresponds to a deposition rate of 0.05-0.2 km h-i. No attempts have been made to optimize the deposition process with respect to temperature and deposition rate. However, the results suggest that the deposition process is limited by the deposition of C from C3Hs. Higher deposition rates can therefore probably be obtained also at low temperatures by using alternative carbon sources (e.g. aromatic hydrocarbons). The results in this study also showed that different types of vapour-substrate and film-substrate interactions strongly affected the WC CVD process. This was investigated by comparing two transition metals (Ta and Ni), which exhibit different properties with respect to, for example, chemical reactivity and carbide formation. During the initial stages of growth, it was found that WF, can be reduced by chemically reactive substrates (e.g. Ta) under the formation of metallic W and volatile substrate fluorides. A WC film is then formed by carburization of the initially deposited W followed by a much slower simultaneous deposition of W and C. It is conceivable that the substrate reactivity can explain the low apparent activation energies observed for the deposition of tungsten and carbon in this study. The deposition process was

also affected by the presence of substrate carbides (TaC and Ta,C). It is likely that the higher deposition rates observed on the carburized Ta substrates are due to a growth mechanism where carbon in the WC is partly supplied by the substrate material. WC films could also be obtained on Ni. This substrate, however, exhibited a quite different behaviour compared with Ta. A high initial deposition rate was also observed on Ni, but the WC films were formed by a series of phase-transformations starting with q-carbides. The decomposition of the q-carbides resulted in the formation of metallic Ni grains in the bulk as well as on the surface of the final WC films. The WC films also contained significant amounts of free carbon amorphous carbon and graphite). A tentative model was proposed suggesting that the high carbon contents in the films was due to a catalytic effect of Ni grains on the surface. It was also suggested that the Ni grains caused whisker-like outgrowths of WC by a VS mechanism in analogy with similar observations in TiC CVD. In summary, the results above clearly show that the substrate material is an important parameter in CVD of WC films. It is likely that the growth behaviour observed on Ta and Ni can be found on many substrates of other transition metals and alloys thereof. The influence of the substrate material can certainly be reduced but not eliminated by a modified process working at lower deposition temperatures. This means that systematic investigations of substrate/vapour and substrate/ films interactions are of the highest importance in WC CVD.

Acknowledgements This work was supported by the Swedish Natural Science Research Council (NFR) and the Swedish Research Council for Engineering Studies (TFR). The authors are also most grateful to Lars Norin for his skilful assistance with the AFM study.

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