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Chemical vapour deposition of molybdenum carbides: aspects of nanocrystallinity
Thin Solid Films 396 Ž2001. 53᎐61
Chemical vapour deposition of molybdenum carbides: aspects of nanocrystallinity J. LuU , U. Jansson Department of Inorganic Chemistry, Uppsala Uni¨ ersity, P.O. Box 538, S-751 21 Uppsala, Sweden Received 23 September 2000; received in revised form 28 March 2001; accepted 15 May 2001
Abstract Nanocrystalline films of ␦-MoC 1yx have been deposited from a MoCl 5rH 2rC 2 H 4 mixture at 800⬚C and a total pressure of 1.7 torr. The average grain size, free carbon content and the deposition rate were affected by different experimental parameters such as vapour composition and substrate position in the reactor. One of the most important parameters was the C 2 H 4 concentration, which at higher concentrations reduces the grain size and the deposition rate. The results suggest that the nanocrystallinity was not caused by a high supersaturation. A more likely explanation is that the observed growth behaviour is caused by strongly adsorbed species from decomposing C 2 H 4 which reduce the surface diffusion lengths during film growth. 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Chemical vapour deposition ŽCVD.; Nanocrystalline structure; Transmission electron microscopy ŽTEM.; Carbide
1. Introduction Today the development of new, nanocrystalline materials is a rapidly growing field in materials science Žsee, e.g. w1᎐3x.. The research is driven by the fact that the nanocrystalline materials frequently exhibit different properties compared to conventional coarse-grained polycrystalline materials. This is a consequence of a large surfacerbulk ratio, which gives materials dominated by disordered interfacial regions such as grain boundaries and solidrvapour surfaces. In general, these regions have deviations of atomic densities and interatomic spacing from the bulk. Misfits in the interfaces can also induce strain fields from the disordered regions into the crystallites and remove atoms from their ideal lattice sites. The large fraction of interfacial regions and associated defects leads to a new type of material with unique chemical and physical properties.
U
Corresponding author. Tel.: q46-018-417-6837. E-mail address: [email protected] ŽJ. Lu..
An interesting group of nanocrystalline materials is the transition metal carbides ŽTiC, WC and NbC, etc... The conventional coarse-grained materials of the transition metal carbides exhibit high hardness, high melting points, and a chemical inertness making them useful as wear-resistant materials. Furthermore, several transition metal carbides ŽWC and MoC. also have a catalytic activity comparable to the platinum metals suggesting potential use in many catalytic processes w4x. It has recently been suggested that composites of nanocrystalline grains in an amorphous matrix could lead to new materials with interesting properties w5x. Voevodin et al. have also demonstrated that composites of 10-nm TiC crystallites in an amorphous carbon matrix have unique mechanical and tribological properties w6x. It is quite clear that the preparation of nanocrystalline carbides should make it possible to tune the mechanical as well as the chemical properties of this group of compounds. The synthesis of well-defined, nanocrystalline carbide structures, however, requires not only a control of the grain size distribution, but
0040-6090r01r$ - see front matter 䊚 2001 Elsevier Science B.V. All rights reserved. PII: S 0 0 4 0 - 6 0 9 0 Ž 0 1 . 0 1 1 5 6 - 7
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J. Lu, U. Jansson r Thin Solid Films 396 (2001) 53᎐61
also control of the amount of free carbon Že.g. graphite. in the grain boundaries and the carbidervapour interface. The nanocrystalline carbides, can be synthesised by many techniques Žsee, e.g. w6᎐11x.. For example, one of the most widely used techniques to produce bulk nanocrystalline carbide is the mechanical milling of a coarser material into a fine-grained powder Žsee, e.g. w7x.. An alternative method to make a large amount of powder is to carbonise a metal oxide in a carbon-containing atmosphere. The higher density of the carbide leads to the formation of an ultra-fine powder which can be sintered to a higher density at elevated temperatures Žsee, e.g. w8x.. Thin films of nanocrystalline transition metal carbides have also been deposited by physical vapour deposition ŽPVD. using for instance a sputtering process. The carbide films can either be sputtered directly as a nanocrystalline material or be formed indirectly by an initial deposition of an amorphous film which is allowed to crystallise into a finegrained material in a subsequent annealing process Žsee, e.g. w5,9x.. Fine-grained or nanocrystalline transition metal carbides can also be deposited by chemical vapour deposition ŽCVD. using suitable reactants Žsee, e.g. w10,11x.. CVD has, for example, been used to produce a very fine-grained carbide powder, which in a later stage can be sintered to a bulk carbide sample w12,13x. CVD can also be used to produce thin films with a wide range of microstructures. An advantage with the CVD technique is that films can be deposited on substrates with a more complex shape. Furthermore, the grain size can to a large extent be controlled by a fine-tuning of different experimental parameters such as vapour compositions, total pressure, temperature, etc. It is generally accepted that the grain size of CVD films can be controlled in at least two different ways: Ži. by changes in the supersaturation; and Žii. by reducing the surface diffusion lengths of adsorbed species w14x. An increased supersaturation will, according to the classical nucleation theory, decrease the critical size of nuclei and reduce the activation energy for nucleation w15x. This will also lead to a general increase in not only the nucleation rate, but also in the final deposition rate. Consequently, a grain-size variation caused by changes in the supersaturation will result in a general relationship between the grain size and the deposition rate as shown by curve a in Fig. 1. The second possibility to control the nucleation rate and thereby the final grain diameter is to reduce the surface diffusion lengths by different mechanisms such as blocking of surface sites by strongly adsorbed molecules or by the formation of a passivating surface layer. Shorter diffusion lengths make grain growth more difficult and will favour the nucleation of new grains. It will also decrease the overall deposition rate and give a
Fig. 1. Relationship between growth rate and grain size due to Ža. change in supersaturation and Žb. reduced surface diffusion lengths.
relationship between growth rate and grain size as shown by curve b in Fig. 1. We have in previous studies observed that CVD of molybdenum carbides from MoCl 5rC 2 H 4rH 2 and MoCl 5rCH 2 I 2rH 2 gas mixtures yields nanocrystalline films of the cubic ␦-MoC 1y x phase w10,16x. It was also found that free carbon could be deposited together with the carbide. The grain size distribution as well as the formation of free carbon influence the physical and chemical properties of these materials. It is therefore important to learn how the microstructure of this type of CVD carbide film can be controlled. The main objective of the present study was to investigate how the experimental parameters in the CVD of MoC affect the grain size as well as the free carbon formation in the films and to find a possible explanation for the nanocrystalline growth behaviour. 2. Experimental The molybdenum carbide thin films were deposited in a horizontal hot-wall CVD quartz reactor from a gas mixture of MoCl 5 Ž99%., H 2 Ž99.9999%. and C 2 H 4 Ž99.95%.. The CVD system has been described previously by Lu et al. w16x. C 2 H 4 and H 2 were mixed before the reactor and transported into the heated region in a separate tube. The MoCl 5 powder was heated in a sublimator and then transported into the reactor through a separate tube using Ar Žpurity 99.9999%. as a carrier gas. The MoCl 5 partial pressure in the reactor was controlled by the sublimation temperature and the flow rate of carrier gas. The flow rates of all the gases were controlled by mass flowmeters. Mixing of all the gases was carried out approximately 30 cm into the reactor. The system was pumped by a mechanical pump yielding a base pressure of 10y3 torr. A liquid N2 trap was placed between the reactor and the pump to minimise the oil back diffusion. The system leaking rate was less than 10y5 torrrs. All the experiments in this study were carried out at 800⬚C using a total pressure of 1.7 torr. In most experiments, the total gas flow was 480 sccm, while the partial
J. Lu, U. Jansson r Thin Solid Films 396 (2001) 53᎐61
pressure of H 2 was kept constant at 0.4 torr Ž25% H 2 . and the partial pressure of MoCl 5 in the vapour was kept constant at approximately 2% by setting the sublimation temperature to 70⬚C and maintaining a constant carrier gas flow rate in all experiments. Small pieces of single-crystal ␣-Al 2 O 3 wafers Žsapphire. with a Ž012. orientation were used as substrates. Prior to a deposition experiment, the substrates were cleaned by acetone in ultrasonic vibration and then placed parallel to the gas flow in the reactor. The chemical composition of the carbide films was investigated with X-ray photoelectron spectroscopy ŽXPS. in a PHI 5500 multitechnique instrument using monochromated AlK ␣ radiation. The binding energy scale was calibrated by setting the C1s peak for adsorbed hydrocarbons to 284.6 eV. The free carbon content in the carbide films was measured by calculating the area ratio of the C᎐C peak at approximately 284.3 eV to the C᎐Mo peak at approximately 283.2 eV. The phase composition of the films were determined by X-ray diffraction ŽXRD. using a SIEMENS D5000 diffractometer with CuK ␣ radiation in a ᎐2 geometry. The grain sizes of the films were measured from XRD using the Scherrer formula w17x: t s arBcos Ž . where t is the grain size, a is an instrumental constant, B is the FWHM of the diffraction peak and is the wavelength of the X-rays. The FWHM of the peaks was calculated by subtracting the K ␣ 2 contribution and assuming a Gaussian peak shape. The grain size was found to vary for different diffraction peaks. All grain sizes given in the paper are average grain sizes calculated from the Ž111., Ž200. and Ž220. peaks. The grain size of the films was also studied by transmission electron microscopy ŽTEM. using a JEOL 2000 FXII instrument with a 200-kV working voltage. The film thicknesses were measured both by TEM and X-ray fluorescence spectroscopy ŽXRFS.. The latter technique gives the total amount of deposited Mo atomsrcm2 , which can be converted into film thickness by assuming ideal film density and uniformity. Finally, TEM was also used to study the nanocrystalline microstructures of the molybdenum carbide films. Crosssections of two samples were prepared by gluing two pieces of the films together face-to-face, then cut, polished, dimpled and finally ion milled.
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puter code EKVICALC w18x., where vapour composition, temperature and pressure are required as input data. The program calculates the composition of the system Žamounts of gaseous species and condensed phases. at equilibrium by minimising the Gibbs free energy, G. It also gives out different thermodynamic parameters such as ⌬G, ⌬ H and ⌬ S for the process. The thermodynamic data necessary for the calculations are taken from a database of approximately 3000 species. It is important to point out that the thermodynamic calculations give the equilibrium state, which is rarely obtained in CVD. However, the results could be used to predict general trends in a CVD process with changes in experimental parameters. In the following calculations, the temperature and total pressure were set at 800⬚C and 1.7 torr, respectively. The H 2 concentration was kept constant at 25%. The C 2 H 4rMoCl 5 ratio was varied from 0.1 to 0.5 keeping the total amount of MoCl 5 constant at 2.5%. The results, which are presented in Fig. 2, show that the driving force ⌬G is more or less linearly proportional to the ethene content but with a little change in slope at a C 2 H 4rMoCl 5 ratio of 0.25. At this ratio all MoCl 5 has reacted and formed carbide. Hence, the curve above this ratio represents the change in ⌬G when graphite starts to form. From a thermodynamic point of view, increasing C 2 H 4 concentrations in the vapour should increase not only the nucleation rate but also the deposition rate Žnote that kinetic restrictions are not considered at this point.. This means that if the grain size is affected only by the supersaturation, with increasing C 2 H 4 concentration we should expect a relationship between grain size and higher deposition rate similar to curve a in Fig. 1. 4. Experimental results 4.1. General obser¨ ations In a previous study we have shown that nanocrystalline films of ␦-MoC 1y x can easily be deposited from
3. Thermodynamic calculations As stated above, the grain size of a CVD film can be influenced by the supersaturation, which is related to the change in Gibbs free energy, ⌬G, for the process. In this study we have carried out thermodynamic calculations using a free energy minimising technique Žcom-
Fig. 2. ⌬G for carbide and graphite formation as a function of the C 2 H 4rMoCl 5 ratio.
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J. Lu, U. Jansson r Thin Solid Films 396 (2001) 53᎐61
a MoCl 5rC 2 H 4rH 2 mixture at 800⬚C w16x. Two other carbide phases ŽMo 2 C and ␥⬘-MoC 1yx . could be deposited at certain experimental conditions but they were never found to exhibit a nanocrystalline microstructure. In the present study, we carried out a series of preliminary experiments at different deposition conditions to gain a more detailed understanding of the nanocrystalline growth behaviour. The total pressure, temperature and gas flow velocity were kept constant in these experiments. We found that the C 2 H 4 concentration and the substrate position in the reactor were two very important parameters and they were therefore studied in more detail Žsee below.. The MoCl 5 concentration in the vapour also affected the grain size. An increase of the sublimator temperature Ži.e. higher MoCl 5 pressures. yielded films with larger grain size. Furthermore, from an earlier study we know that hydrogen plays an important role for the nanocrystalline growth w16x. It was shown that CVD from a MoCl 5rC 2 H 4 mixture Ži.e. no H 2 added to the vapour. yielded an initial deposition of the ␦-MoC 1y x phase ˚ with an average grain size of approximately 250᎐300 A. After a while, however, the phase composition changed ˚. as ␥⬘- MoC 1y x with a fairly large grain size Ž) 1500 A was formed on top of the ␦-MoC 1y x film. An addition of approximately 10% H 2 , however, changed the phase composition to the cubic ␦-MoC 1y x phase which always exhibited a nanocrystalline microstructure Žgrain ˚ .. Consequently, it can be concluded that sizes - 150 A hydrogen itself is required for inducing and maintaining the deposition of the nanocrystalline ␦-MoC 1y x phase. In the present study we have investigated the influence of different H 2 concentrations and found a minor influence on the grain size Žproviding that the content is larger than approx. 10%.. 4.2. Influence of substrate position Analysis of films deposited at different positions in the reactor showed variations in the phase composition and in the total carbon content. In general, molybdenum-rich films of metallic Mo and Mo 2 C were deposited the first 3 cm after mixing of the reactants Žthe gases were mixed approx. 30 cm into the reactor tube.. These films always exhibited a large grain size. The nanocrystalline ␦-MoC 1y x phase was only deposited further into the reactor. In a previous study, this effect was explained by differences in reactivity between MoCl 5 and C 2 H 4 w16x. MoCl 5 is a very reactive compound and reacts easily with H 2 above 500⬚C. In contrast, C 2 H 4 needs to be ‘activated’ by thermal excitation andror homogeneous reactions prior to decomposition w16x. This will lead to a depletion of MoCl 5 in the vapour and increase the relative amount of hydrocarbons further downstream in the reactor. ␦-MoC 1y x films deposited at different positions in
Fig. 3. C1s XPS spectra of ␦-MoC 1y x films deposited at different positions in the reactor. Ža. Substrate position 5 cm and Žb. 9 cm after gas mixing. C 2 H 4 content: 12.5%. H 2 content: 25%.
the reactor changed from a metallic to an almost black nuance with increasing distance from the reactor entrance suggesting an increase in the carbon content downstream in the reactor. This was confirmed by XPS showing a peak at approximately 284.3 eV, which can be attributed to C᎐C andror C᎐H bonds. In the following this carbon is denoted ‘free carbon’. In addition, a C᎐Mo peak was seen at approximately 283.2 eV in all the films Žsee Fig. 3.. The intensity of the free-carbon peak decreased with increasing sputtering time and finally reached a constant value for the ratio of free carbon to carbidic carbon approximately 5᎐30% Žwhich corresponds to approx. 2᎐11 at% free carbon concentration. after 10 min sputtering. The results show that the films contain more free carbon in the surface region than in the bulk. Furthermore, the carbon content in the surface region and in the bulk increased with increasing distance from the reactor entrance ŽFig. 4.. At present, we are unable to identify whether the free carbon consists of graphite or some other form of carbon. However, no graphite has been observed by XRD. The detection limit for graphite is estimated to be approximately 10%. A rough estimation shows that the amount of ‘free carbon’ observed in the XPS analy˚ thick ses corresponds to a graphite layer less than 1 A ˚ if the grain size is assumed to be 100 A. As can be seen in Fig. 4 most of the films consist of carbidic carbon. XRD also showed that the cubic ␦MoC 1y x phase was the only carbide formed at a substrate position of more than 4 cm from the mixing of the gases. An interesting observation was that ␦MoC 1y x films deposited at different positions in the reactor have an almost identical cell parameter of 4.26 ˚ , which corresponds to a composition of approxiA mately ␦-MoC 0.67 using the data for cell parameter vs. carbon content given by Rudy et al. w19x. This result is in good agreement with the XPS analysis showing a composition of approximately MoC 0.68 ᎐ 0.70 using elemental standards and only including the contribution from the carbidic carbon.
J. Lu, U. Jansson r Thin Solid Films 396 (2001) 53᎐61
Fig. 4. Variation of free carbon content in the films as a function of sample position after gas mixing. Ža. 1 min sputtering; Žb. 3 min sputtering; and Žc. 10 min sputtering.
In general, the X-ray diffractograms of the ␦-MoC 1y x films exhibit very broad diffraction peaks Žsee Fig. 5.. The line-broadening can be assigned to several factors such as a small grain size and stresses in the film. A plot of Bcos vs. sin shows no indication of stresses. Most of the line-broadening can therefore be attributed to a small grain size. The average grain diameter was estimated to vary from approximately 50 to 170 ˚ using the Scherrer formula. The grain size was also A found to be dependent on the sample position ŽFig. 6.. In general, the grain size decreased further downstream in the reactor. However, the change in grain size was also dependent on the C 2 H 4 content. The largest variation was found for films deposited at low C 2 H 4 concentrations while those deposited at higher concentration only showed a minor change in grain size at different substrate positions ŽFig. 6.. 4.3. Influence of C2 H4 content The influence of C 2 H 4 on the deposition process was investigated in a series of experiments with C 2 H 4 concentrations ranging from 0.625% to 12.5% Žthe MoCl 5 and H 2 flows were kept constant.. XRD showed that all the films in this series consisted of ␦-MoC 1y x . The cell parameter of the ␦-MoC 1y x films was de-
Fig. 5. X-Ray diffractograms of ␦-MoC 1y x films deposited at Ža. 0.625% C 2 H 4 and Žb. 12.5% C 2 H 4 , respectively.
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Fig. 6. Influence of sample position on grain size at different C 2 H 4 concentrations. Ža. 0.625% C 2 H 4 , Žb. 1.875% C 2 H 4 and Žc. 12.5% C2 H4.
˚ which corresponds to a compositermined to be 4.26 A tion of approximately ␦-MoC 0.67 for all films Ži.e. the carbide composition was independent of the C 2 H 4 concentrations .. Furthermore, the XPS results showed an unexpected relationship between free carbon content and vapour composition. It was found that the highest free carbon contents in the bulk Ž11 at.%. were observed in the films deposited at low ethene content Ž0.625%., while the film deposited with 12.5% ethane contained only 2 at.% free carbon. An interesting observation was that films deposited at ethene concentrations below 2% exhibited a Ž110. texture. In contrast, films deposited at ethene concentration above 3% were more or less texture-free Žcf. Fig. 5a,b.. Furthermore, it was also found the linebroadening of the XRD peaks was dependent on the ethene content in the vapour. A general observation was that the line-width of the diffraction peaks increased at higher ethene concentrations. From the Scherrer formula, the average grain sizes were esti˚ by decreasing the mated to vary from 50 to 170 A ethene content from 12.5 to 0.625% Žsee Fig. 7.. The deposition rate of the films was estimated from XRFS analyses. It was found that the deposition rate
Fig. 7. Influence of C 2 H 4 concentration on the grain size of a ␦-MoC 1y x film deposited at a sample position of 5 cm from the gas mixing.
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Fig. 8. Deposition rate of the ␦-MoC 1y x phase as a function of grain size.
decreased with increasing ethene content from approximately 1.4= 10 19 Morcm2 h at 0.625% C 2 H 4 to approximately 6 = 10 18 Morcm2 h at 12.5% C 2 H 4 . A comparison between deposition rates and grain sizes shows that higher deposition rates yielded a larger grain size Žsee Fig. 8.. The influence of the ethene content on the film microstructure was also studied by TEM. Two samples deposited at 0.625% and 12.5% C 2 H 4 were chosen for the TEM observations. The TEM image Žsee Fig. 9. clearly shows that the film deposited with 12.5% C 2 H 4 consisted of almost texture-free nanocrystalline grains ˚ which are with diameters of approximately 50᎐100 A, in good agreement with the value obtained by XRD. In contrast, the microstructure of the films obtained with 0.625% C 2 H 4 is quite different Žsee Fig. 10.. At this ethene composition, the carbide exhibits a dendrite-like structure with Ž110. texture Žsee the SAD inset in Fig. 10.. It was found that the dendrite-like microstructure exhibited a larger distribution of grain sizes. Moreover, the dendrite-like film also contained a lot of pores. The porosity was estimated to be approximately 50% by
Fig. 10. TEM micrograph of nanocrystalline ␦-MoC 1y x film with a dendrite-like microstructure deposited at 0.625% C 2 H 4 .
comparing the film thickness shown by cross-section TEM image with the thickness calculated from XRFS analysis. In contrast, the porosity of the texture-free film in Fig. 9 was found to be negligible. 5. Discussion The results above show that nanocrystalline ␦MoC 1y x films can be deposited from a MoCl 5r C 2 H 4rH 2 mixture. It is also clear that we, within certain limits, can control the average grain size of the films. To explain the mechanism behind the nanocrystalline growth behaviour we need to determine how the carbide phase is formed. One mechanism for carbide formation in CVD is that the carbide is formed directly by a simultaneous addition of metal and carbon by surface reactions involving adsorbed hydrocarbon and metal halide species. The final grain size will then be defined by the nucleation kinetics of the carbide grains. Several authors, however, have presented an alternative growth models for carbide CVD using halides as metal sources, which are based on the assumption that small grains of the metal initially are formed by a hydrogen-reduction step followed by a carbonisation of the metal by, e.g. a hydrocarbon: MeX n Ž g. q nr2 H 2 Ž g. ª Me Ž s . q nHXŽ g. Me Ž s . q C x H y Ž g. ª MeC Ž s . q H 2 Ž g.
Fig. 9. TEM micrograph of nanocrystalline ␦-MoC 1y x film deposited with 12.5% C 2 H 4 .
In principle, the metal grains can be formed either by a heterogeneous process on the carbide surface or by a homogeneous reaction in the vapour. In the latter case, the metalrcarbide grain will be deposited on the surface of the growing carbide film yielding a powdery deposit. In both cases, however, the size of the final
J. Lu, U. Jansson r Thin Solid Films 396 (2001) 53᎐61
carbide grains will be defined by the size distribution of the initially formed metal grains. This type of growth behaviour has been reported in processes where the main purpose has been to produce a carbide powder for further processing. Such processes are usually carried out at high pressures where the supersaturation and the nucleation rate of the metal is high w20,21x. A similar mechanism has also been reported by Hogberg ¨ et al. for WC CVD on a reactive substrate material Ži.e. Ta. using WF6 as a metal source w22x. In this case, the substrate reduced WF6 under the formation of a metallic W film, which was carbonised in a later stage. Our results, however, clearly show that this simple growth mechanism can be excluded. For example, deposition experiments without C 2 H 4 Ži.e. with a MoCl 5rH 2 gas mixture. yielded only films of metallic Mo with a rather large grain size Žapprox. 1 m.. Furthermore, rather coarse-grained films of metallic Mo and Mo 2 C were generally deposited the first 3 cm after mixing of the reactants, while more carbon-rich, nanocrystalline ␦MoC 1y x films were deposited after a certain distance into the reactor where the relative concentration of C 2 H 4 is higher due to the depletion of MoCl 5 . Thus it can be concluded that the nanocrystallinity is correlated to the presence of C 2 H 4 in the vapour. This is also clearly seen in Fig. 7, which shows that the grain size is strongly influenced by the C 2 H 4 concentration. As discussed above the nanocrystalline growth can be due to two main factors: Ži. high supersaturation; and Žii. reduced surface diffusion lengths. In many CVD processes, nanocrystalline films are caused by a high supersaturation resulting in a very fast homogeneous nucleation step in the gas phase. The film is then formed by particles ‘snowing’ towards the substrate surface yielding a very powdery and porous film microstructure. This growth behaviour, however, is rather unlikely in our process since we have used a rather low pressure Ž1.7 torr. and a diluted reaction gas mixture. We have also not visually observed any particle formation in the gas phase or seen any indications of particle transport after the deposition zone. Furthermore, most of our films Žwith exception of those deposited at lower C 2 H 4 contents. exhibit a dense microstructure without pores or agglomerated particles. These results suggest that the nanocrystalline ␦-MoC 1y x films be formed by an extensive heterogeneous nucleation on the substrate andror the growing film surface. We believe, however, that a high supersaturation alone cannot be responsible for the high nucleation rates required for the nanocrystalline growth. This conclusion is supported by the thermodynamic calculations which showed that ⌬G Žthe driving force. for the process will increase with increasing C 2 H 4rMoCl 5 ratios in the vapour Žsee Fig. 2.. If the nanocrystalline growth is caused by only a high supersaturation, an increase in the C 2 H 4 content should lead not only to smaller grain sizes but also to a
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higher deposition rate as shown by curve a in Fig. 1. The results in Fig. 8, however, show an opposite behaviour suggesting that some other mechanism than a high supersaturation must be responsible for the finegrained microstructure. The second possibility is that the small grain size in the ␦-MoC 1y x films is caused by reduced surface diffusion lengths. Our results are consistent with such a growth mechanism. This conclusion is supported by, for example, Fig. 8, showing a strong correlation between smaller grain sizes and reduced deposition rates. This relationship is in good agreement with a growth process controlled by a reduced diffusion lengths as shown by curve b in Fig. 1. The results above also suggest that the presence of C 2 H 4 is critical for this behaviour Žsee, e.g. Fig. 7.. This conclusion is also supported by the XPS analyses which show that all nanocrystalline films contain free carbon Žor more correct carbon with mainly C᎐C andror C᎐H bonds. at the film surface and in the grain boundaries. We therefore suggest that the nanocrystalline growth behaviour of the ␦-MoC 1y x films is due to a ‘poisoning’ of the carbide surface by strongly adsorbed reaction products from the decomposed C 2 H 4 . A likely mechanism can then be as follows: small carbide nuclei are formed on the surface of the substrate or on the growing film. The surfaces of the nuclei are quickly covered by strongly adsorbed species from the decomposing C 2 H 4 , which reduces the surface diffusion rates of the reactants. As a consequence the growth rate of the nuclei will be low. As further reactants are adsorbed, the supersaturation on the surface will increase leading to the nucleation of new carbide nuclei. The final film will exhibit a microstructure with nanometre-sized ␦-MoC 1y x grains with some free carbon at the grain boundaries. At present, we have no clear evidence for the exact nature of these carbon-containing reaction products. A possibility is that a graphitic-like overlayer is formed on the carbide surface. As mentioned above the free carbon ˚ thick if corresponds to a graphite layer less than 1 A ˚ grain size is assumed to be 100 A. This thickness is less than a monolayer of a graphite and show that no continuous graphite phase can exist in the boundary regions of the nanocrystalline grains. Thus it is clear that further studies with, e.g. other in situ surface analysis techniques andror high-resolution TEM is required to determine the nature of the free carbon and its importance for the observed growth behaviour. The model for nanocrystalline growth proposed above is in agreement with observations in the literature. For example as mentioned in Section 1, it is well known that molybdenum carbide surfaces in many respects exhibit large similarities with many platinum metals w4x. It is also known that so-called ‘carbonaceous overlayers’ are formed on platinum metal surfaces in the presence of hydrocarbons w23x. These overlayers are
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catalytically active at lower temperatures. At higher temperatures, however, they loose hydrogen under the formation of ‘graphite-like’ carbon on the surface. The presence of such ‘graphitic’ overlayer inhibits many surface reactions and makes the surface catalytically inactive. In addition, there is some experimental evidence that similar effects can be also found on molybdenum carbide surfaces. Carburisation experiments of Mo with hydrocarbons have shown that a carbon overlayer with graphitic-like characters Žas shown with Auger electron spectroscopy. is formed at elevated temperatures and high hydrocarbon exposures w24x. The presence of such an overlayer was found to drastically reduce the rate of the carbonisation process. Evidence of carbon-containing overlayers has also been observed in the CVD of boron carbides on transition metal substrates w25x. It was found that the nucleation and initial deposition rates of the boron carbide were strongly reduced by some hydrocarbons Že.g. C 2 H 4 .. The effect, which was observed on Mo substrates but not on other metals such as Ti, was attributed to the formation of a carbon-containing overlayer on the surface w25x. A growth mechanism controlled by a low surface mobility and reduced surface diffusion lengths can also explain other experimental observations. The deposition profile described in Section 4.2, for example, shows that molybdenum-rich phases Ži.e. Mo and Mo 2 C. with a large grain-size are deposited at the entrance of the reactor and that nanocrystalline ␦-MoC 1y x films only are formed at some distance into the reactor. This observation together with the fact that the grain size is reduced with increasing distance from the reactor entrance suggest that the surface blocking is more effective further downstream in the reactor Žsee Fig. 6.. This can attributed to the fact that many hydrocarbons have to be ‘activated’ by, e.g. thermal excitation andror homogeneous reactions prior to decomposition at the fairly low temperature used in our experiments Ž800⬚C.. The activation requires a certain time in the heated zone of the furnace and leads to strong variations in the growth conditions at different positions in the deposition reactor. This also suggests that a chemically more active carbon source should give a quite different deposition profile compared to C 2 H 4 . An example of such a precursor is CH 2 I 2 , which rapidly decomposes into I and CH 2 radicals upon heating. In a previous study, we have also shown that carbide CVD from a MoCl 5rCH 2 I 2rH 2 mixture yields nanocrystalline ␦MoC 1y x film in all parts of the reactor without any formation of large-grained Mo and Mo 2 C at the reactor entrance and with minor variations in grain sizes and free carbon content along the reactor w10x. This can be explained by the fact that the CH 2 radicals are more effective in forming the carbon-containing overlayer which reduces surface mobility.
Our results also show that hydrogen plays an important role for the nanocrystalline growth. In an earlier study, it was shown that CVD from a MoCl 5rC 2 H 4 mixture Ži.e. no H 2 added to the vapour. yielded an initial deposition of the ␦-MoC 1y x phase with an aver˚ After a age grain size of approximately 250᎐300 A. while, however, the phase composition changed as ␥⬘MoC 1y x with a fairly large grain size was formed on top of the ␦-MoC 1y x film w16x. An addition of approximately 10% hydrogen, however, changed the phase composition to the cubic ␦-MoC 1y x phase which always exhibited a nanocrystalline microstructure. The observations can be explained by the fact that hydrogen is likely to be involved in different homogeneous andror heterogeneous reaction steps in the decomposition of C 2 H 4 . It is therefore possible that the presence of H 2 in the vapour is required for the formation of a carbon-containing overlayer which can effectively reduce the surface mobility and diffusion lengths during growth. Finally, it is conceivable to assume that the formation of a carbon-containing overlayer which reduces the surface mobility of adsorbed reactants is dependent on the crystallographic orientation of the carbide surface. Some surface orientations are probably more easy to ‘poison’ than other. This should give rise to a film texture where some grains grow faster than other grains with unfavourable orientations. In fact, the XRD and TEM studies showed that films with a Ž110.-texture and a dendrite-like growth behaviour were deposited at low C 2 H 4 contents in the vapour Žsee Figs. 5 and 10.. These films also exhibited with a high degree of porosity and a large size distribution of the carbide grains. In contrast, films deposited at higher C 2 H 4 concentrations were almost texture-free, rather dense and exhibited a more narrow grain-size distribution. The growth behaviour at high C 2 H 4 concentrations can be explained by the fact that all carbide surface irrespective of crystallographic orientation is covered by the carbon-containing overlayer. Films deposited at these conditions will therefore be texture-free. At low C 2 H 4 concentrations, however, it is possible that some crystallographic orientations are less affected by this overlayer thereby giving rise to a textured film. The porosity of films deposited at 0.625% C 2 H 4 also explains the high contents of free carbon in these films. The large total surface area in the porous films will contain large amounts of free carbon Žsee Figs. 3 and 4. which will give an apparent high ‘bulk concentration’ also after sputtering. 6. Conclusions CVD of molybdenum carbides from a MoCl 5r H 2rC 2 H 4 mixture at 800⬚C leads to the formation of nanocrystalline ␦-MoC 1y x films. The average grain size,
J. Lu, U. Jansson r Thin Solid Films 396 (2001) 53᎐61
free carbon content and the deposition rate are affected by the experimental parameters. The experiments clearly show that high C 2 H 4 concentrations favour the formation of nanocrystalline grains. Our results suggest that a high supersaturation is an unlikely explanation for the observed growth behaviour. We have therefore proposed a model where the nanocrystalline growth is caused by a reduced surface mobility induced by the presence of some strongly adsorbed reaction products from C 2 H 4 . Hitherto, this explanation is based on general growth trends Že.g. Figs. 7 and 8.. Further studies using, for example, in situ analyses with different surface science techniques under well-defined conditions Ži.e. UHV. are required to elucidate the exact nature of such carbon-containing reaction products and their influence on the growth process. It is clear, however, that the proposed growth model is in agreement with many observations in the literature. Acknowledgements The Swedish Research Council for Engineering Sciences ŽTFR., the Swedish Natural Science Research ˚ Council ŽNFR., and the Angstrom ¨ Consortium are acknowledged for financial support. References w1x w2x w3x w4x
H. Gleiter, Prog. Mater. Sci. 33 Ž1989. 223. K. Lu, Mater. Sci. Eng. R 16 Ž1996. 161. R. Birringer, Mater. Sci. Eng. A 117 Ž1989. 33. S.T. Oyama, Catalysis Today 15 Ž1992. 179.
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w5x S. Veprek, S. Reiprich, Thin Solid Films 268 Ž1995. 64. w6x A.A. Voevodin, S.V. Prasad, J.S. Zabinski, J. Appl. Phys. 82 Ž1997. 855. w7x M.S. El-Eskandarany, T.J. Konno, K. Sumiyama, K. Suzuki, Mater. Sci. Eng. A 217r218 Ž1996. 265. w8x J.S. Lee, L. Volpe, F.H. Ribeiro, M. Boudart, J. Catalysis 112 Ž1988. 44. w9x F. Okuyama, Y. Fujimoto, S. Kato, T. Kondo, Appl. Phys. A 38 Ž1985. 275. w10x H. Siriwardane, O. Pringle, J.W. Newkirk, W.J. James, Thin Solid Films 287 Ž1996. 8. w11x J. Lu, U. Jansson, Proceedings of the Fourteenth International Conference and EUROCVD-11. Electrochem. Soc. Pennington, NJ, USA; 1997, xxii q 1652, pp. 333᎐340. w12x C-W. Won, B-S. Chun, H.Y. Sohn, J. Mater. Res. 8 Ž1993. 2702. w13x J. Hojo, M. Tajika, A. Kato, J. Less-Common Met. 75 Ž1980. 11. w14x J.M. Blocher, Jr., J. Vac. Sci. Technol. 11 Ž1974. 680. w15x R.E. Reed-Hill, R. Abbaschian, Physical Metallurgy Principles, 3rd ed., PWS-Kent Publishing Company, Boston, 1992, p. 479. w16x J. Lu, H. Hugosson, O. Eriksson, L. Nordstrom, ¨ U. Jansson, Thin Solid Films 370 Ž2000. 203. w17x B.D. Cullity, Elements of X-ray Diffraction, 2nd ed., AddisonWesley Publishing Company, London, 1978, p. 102,’. w18x B. Nolang, ¨ Thesis, Acta Univ. Ups ŽFac. Sci.. 691 Ž1983.. w19x E. Rudy, St. Windisch, A.J. Stosick, J.R. Hoffman, Trans. Metallurg. Soc. AIME 239 Ž1967. 1247. w20x X. Tang, R. Haubner, B. Lux, B. Kieffer, J. Phys. II C5 Ž1995. 1013. w21x J. Hojo, T. Oku, A. Kato, J. Less-Common Metals 59 Ž1978. 85. w22x H. Hogberg, P. Tagtstrom, ¨ ¨ ¨ J. Lu, U. Jansson, Thin Solid Films 272 Ž1996. 116. w23x G.A. Somorjai, Introduction to Surface Chemistry and Catalysis, Wiley-Interscience, New York, 1994, p. 505. w24x E.V. Rutkov, A. Ya. Tontegode, M. Usufov, N.R. Gall, Appl. Surf. Sci. 78 Ž1994. 179. w25x U. Jansson, J-O. Carlsson, B. Stridh, J. Vac. Sci. Technol. A 5 Ž1987. 2823.
Chemical vapour deposition of molybdenum carbides: aspects of nanocrystallinity J. LuU , U. Jansson Department of Inorganic Chemistry, Uppsala Uni¨ ersity, P.O. Box 538, S-751 21 Uppsala, Sweden Received 23 September 2000; received in revised form 28 March 2001; accepted 15 May 2001
Abstract Nanocrystalline films of ␦-MoC 1yx have been deposited from a MoCl 5rH 2rC 2 H 4 mixture at 800⬚C and a total pressure of 1.7 torr. The average grain size, free carbon content and the deposition rate were affected by different experimental parameters such as vapour composition and substrate position in the reactor. One of the most important parameters was the C 2 H 4 concentration, which at higher concentrations reduces the grain size and the deposition rate. The results suggest that the nanocrystallinity was not caused by a high supersaturation. A more likely explanation is that the observed growth behaviour is caused by strongly adsorbed species from decomposing C 2 H 4 which reduce the surface diffusion lengths during film growth. 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Chemical vapour deposition ŽCVD.; Nanocrystalline structure; Transmission electron microscopy ŽTEM.; Carbide
1. Introduction Today the development of new, nanocrystalline materials is a rapidly growing field in materials science Žsee, e.g. w1᎐3x.. The research is driven by the fact that the nanocrystalline materials frequently exhibit different properties compared to conventional coarse-grained polycrystalline materials. This is a consequence of a large surfacerbulk ratio, which gives materials dominated by disordered interfacial regions such as grain boundaries and solidrvapour surfaces. In general, these regions have deviations of atomic densities and interatomic spacing from the bulk. Misfits in the interfaces can also induce strain fields from the disordered regions into the crystallites and remove atoms from their ideal lattice sites. The large fraction of interfacial regions and associated defects leads to a new type of material with unique chemical and physical properties.
U
Corresponding author. Tel.: q46-018-417-6837. E-mail address: [email protected] ŽJ. Lu..
An interesting group of nanocrystalline materials is the transition metal carbides ŽTiC, WC and NbC, etc... The conventional coarse-grained materials of the transition metal carbides exhibit high hardness, high melting points, and a chemical inertness making them useful as wear-resistant materials. Furthermore, several transition metal carbides ŽWC and MoC. also have a catalytic activity comparable to the platinum metals suggesting potential use in many catalytic processes w4x. It has recently been suggested that composites of nanocrystalline grains in an amorphous matrix could lead to new materials with interesting properties w5x. Voevodin et al. have also demonstrated that composites of 10-nm TiC crystallites in an amorphous carbon matrix have unique mechanical and tribological properties w6x. It is quite clear that the preparation of nanocrystalline carbides should make it possible to tune the mechanical as well as the chemical properties of this group of compounds. The synthesis of well-defined, nanocrystalline carbide structures, however, requires not only a control of the grain size distribution, but
0040-6090r01r$ - see front matter 䊚 2001 Elsevier Science B.V. All rights reserved. PII: S 0 0 4 0 - 6 0 9 0 Ž 0 1 . 0 1 1 5 6 - 7
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also control of the amount of free carbon Že.g. graphite. in the grain boundaries and the carbidervapour interface. The nanocrystalline carbides, can be synthesised by many techniques Žsee, e.g. w6᎐11x.. For example, one of the most widely used techniques to produce bulk nanocrystalline carbide is the mechanical milling of a coarser material into a fine-grained powder Žsee, e.g. w7x.. An alternative method to make a large amount of powder is to carbonise a metal oxide in a carbon-containing atmosphere. The higher density of the carbide leads to the formation of an ultra-fine powder which can be sintered to a higher density at elevated temperatures Žsee, e.g. w8x.. Thin films of nanocrystalline transition metal carbides have also been deposited by physical vapour deposition ŽPVD. using for instance a sputtering process. The carbide films can either be sputtered directly as a nanocrystalline material or be formed indirectly by an initial deposition of an amorphous film which is allowed to crystallise into a finegrained material in a subsequent annealing process Žsee, e.g. w5,9x.. Fine-grained or nanocrystalline transition metal carbides can also be deposited by chemical vapour deposition ŽCVD. using suitable reactants Žsee, e.g. w10,11x.. CVD has, for example, been used to produce a very fine-grained carbide powder, which in a later stage can be sintered to a bulk carbide sample w12,13x. CVD can also be used to produce thin films with a wide range of microstructures. An advantage with the CVD technique is that films can be deposited on substrates with a more complex shape. Furthermore, the grain size can to a large extent be controlled by a fine-tuning of different experimental parameters such as vapour compositions, total pressure, temperature, etc. It is generally accepted that the grain size of CVD films can be controlled in at least two different ways: Ži. by changes in the supersaturation; and Žii. by reducing the surface diffusion lengths of adsorbed species w14x. An increased supersaturation will, according to the classical nucleation theory, decrease the critical size of nuclei and reduce the activation energy for nucleation w15x. This will also lead to a general increase in not only the nucleation rate, but also in the final deposition rate. Consequently, a grain-size variation caused by changes in the supersaturation will result in a general relationship between the grain size and the deposition rate as shown by curve a in Fig. 1. The second possibility to control the nucleation rate and thereby the final grain diameter is to reduce the surface diffusion lengths by different mechanisms such as blocking of surface sites by strongly adsorbed molecules or by the formation of a passivating surface layer. Shorter diffusion lengths make grain growth more difficult and will favour the nucleation of new grains. It will also decrease the overall deposition rate and give a
Fig. 1. Relationship between growth rate and grain size due to Ža. change in supersaturation and Žb. reduced surface diffusion lengths.
relationship between growth rate and grain size as shown by curve b in Fig. 1. We have in previous studies observed that CVD of molybdenum carbides from MoCl 5rC 2 H 4rH 2 and MoCl 5rCH 2 I 2rH 2 gas mixtures yields nanocrystalline films of the cubic ␦-MoC 1y x phase w10,16x. It was also found that free carbon could be deposited together with the carbide. The grain size distribution as well as the formation of free carbon influence the physical and chemical properties of these materials. It is therefore important to learn how the microstructure of this type of CVD carbide film can be controlled. The main objective of the present study was to investigate how the experimental parameters in the CVD of MoC affect the grain size as well as the free carbon formation in the films and to find a possible explanation for the nanocrystalline growth behaviour. 2. Experimental The molybdenum carbide thin films were deposited in a horizontal hot-wall CVD quartz reactor from a gas mixture of MoCl 5 Ž99%., H 2 Ž99.9999%. and C 2 H 4 Ž99.95%.. The CVD system has been described previously by Lu et al. w16x. C 2 H 4 and H 2 were mixed before the reactor and transported into the heated region in a separate tube. The MoCl 5 powder was heated in a sublimator and then transported into the reactor through a separate tube using Ar Žpurity 99.9999%. as a carrier gas. The MoCl 5 partial pressure in the reactor was controlled by the sublimation temperature and the flow rate of carrier gas. The flow rates of all the gases were controlled by mass flowmeters. Mixing of all the gases was carried out approximately 30 cm into the reactor. The system was pumped by a mechanical pump yielding a base pressure of 10y3 torr. A liquid N2 trap was placed between the reactor and the pump to minimise the oil back diffusion. The system leaking rate was less than 10y5 torrrs. All the experiments in this study were carried out at 800⬚C using a total pressure of 1.7 torr. In most experiments, the total gas flow was 480 sccm, while the partial
J. Lu, U. Jansson r Thin Solid Films 396 (2001) 53᎐61
pressure of H 2 was kept constant at 0.4 torr Ž25% H 2 . and the partial pressure of MoCl 5 in the vapour was kept constant at approximately 2% by setting the sublimation temperature to 70⬚C and maintaining a constant carrier gas flow rate in all experiments. Small pieces of single-crystal ␣-Al 2 O 3 wafers Žsapphire. with a Ž012. orientation were used as substrates. Prior to a deposition experiment, the substrates were cleaned by acetone in ultrasonic vibration and then placed parallel to the gas flow in the reactor. The chemical composition of the carbide films was investigated with X-ray photoelectron spectroscopy ŽXPS. in a PHI 5500 multitechnique instrument using monochromated AlK ␣ radiation. The binding energy scale was calibrated by setting the C1s peak for adsorbed hydrocarbons to 284.6 eV. The free carbon content in the carbide films was measured by calculating the area ratio of the C᎐C peak at approximately 284.3 eV to the C᎐Mo peak at approximately 283.2 eV. The phase composition of the films were determined by X-ray diffraction ŽXRD. using a SIEMENS D5000 diffractometer with CuK ␣ radiation in a ᎐2 geometry. The grain sizes of the films were measured from XRD using the Scherrer formula w17x: t s arBcos Ž . where t is the grain size, a is an instrumental constant, B is the FWHM of the diffraction peak and is the wavelength of the X-rays. The FWHM of the peaks was calculated by subtracting the K ␣ 2 contribution and assuming a Gaussian peak shape. The grain size was found to vary for different diffraction peaks. All grain sizes given in the paper are average grain sizes calculated from the Ž111., Ž200. and Ž220. peaks. The grain size of the films was also studied by transmission electron microscopy ŽTEM. using a JEOL 2000 FXII instrument with a 200-kV working voltage. The film thicknesses were measured both by TEM and X-ray fluorescence spectroscopy ŽXRFS.. The latter technique gives the total amount of deposited Mo atomsrcm2 , which can be converted into film thickness by assuming ideal film density and uniformity. Finally, TEM was also used to study the nanocrystalline microstructures of the molybdenum carbide films. Crosssections of two samples were prepared by gluing two pieces of the films together face-to-face, then cut, polished, dimpled and finally ion milled.
55
puter code EKVICALC w18x., where vapour composition, temperature and pressure are required as input data. The program calculates the composition of the system Žamounts of gaseous species and condensed phases. at equilibrium by minimising the Gibbs free energy, G. It also gives out different thermodynamic parameters such as ⌬G, ⌬ H and ⌬ S for the process. The thermodynamic data necessary for the calculations are taken from a database of approximately 3000 species. It is important to point out that the thermodynamic calculations give the equilibrium state, which is rarely obtained in CVD. However, the results could be used to predict general trends in a CVD process with changes in experimental parameters. In the following calculations, the temperature and total pressure were set at 800⬚C and 1.7 torr, respectively. The H 2 concentration was kept constant at 25%. The C 2 H 4rMoCl 5 ratio was varied from 0.1 to 0.5 keeping the total amount of MoCl 5 constant at 2.5%. The results, which are presented in Fig. 2, show that the driving force ⌬G is more or less linearly proportional to the ethene content but with a little change in slope at a C 2 H 4rMoCl 5 ratio of 0.25. At this ratio all MoCl 5 has reacted and formed carbide. Hence, the curve above this ratio represents the change in ⌬G when graphite starts to form. From a thermodynamic point of view, increasing C 2 H 4 concentrations in the vapour should increase not only the nucleation rate but also the deposition rate Žnote that kinetic restrictions are not considered at this point.. This means that if the grain size is affected only by the supersaturation, with increasing C 2 H 4 concentration we should expect a relationship between grain size and higher deposition rate similar to curve a in Fig. 1. 4. Experimental results 4.1. General obser¨ ations In a previous study we have shown that nanocrystalline films of ␦-MoC 1y x can easily be deposited from
3. Thermodynamic calculations As stated above, the grain size of a CVD film can be influenced by the supersaturation, which is related to the change in Gibbs free energy, ⌬G, for the process. In this study we have carried out thermodynamic calculations using a free energy minimising technique Žcom-
Fig. 2. ⌬G for carbide and graphite formation as a function of the C 2 H 4rMoCl 5 ratio.
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J. Lu, U. Jansson r Thin Solid Films 396 (2001) 53᎐61
a MoCl 5rC 2 H 4rH 2 mixture at 800⬚C w16x. Two other carbide phases ŽMo 2 C and ␥⬘-MoC 1yx . could be deposited at certain experimental conditions but they were never found to exhibit a nanocrystalline microstructure. In the present study, we carried out a series of preliminary experiments at different deposition conditions to gain a more detailed understanding of the nanocrystalline growth behaviour. The total pressure, temperature and gas flow velocity were kept constant in these experiments. We found that the C 2 H 4 concentration and the substrate position in the reactor were two very important parameters and they were therefore studied in more detail Žsee below.. The MoCl 5 concentration in the vapour also affected the grain size. An increase of the sublimator temperature Ži.e. higher MoCl 5 pressures. yielded films with larger grain size. Furthermore, from an earlier study we know that hydrogen plays an important role for the nanocrystalline growth w16x. It was shown that CVD from a MoCl 5rC 2 H 4 mixture Ži.e. no H 2 added to the vapour. yielded an initial deposition of the ␦-MoC 1y x phase ˚ with an average grain size of approximately 250᎐300 A. After a while, however, the phase composition changed ˚. as ␥⬘- MoC 1y x with a fairly large grain size Ž) 1500 A was formed on top of the ␦-MoC 1y x film. An addition of approximately 10% H 2 , however, changed the phase composition to the cubic ␦-MoC 1y x phase which always exhibited a nanocrystalline microstructure Žgrain ˚ .. Consequently, it can be concluded that sizes - 150 A hydrogen itself is required for inducing and maintaining the deposition of the nanocrystalline ␦-MoC 1y x phase. In the present study we have investigated the influence of different H 2 concentrations and found a minor influence on the grain size Žproviding that the content is larger than approx. 10%.. 4.2. Influence of substrate position Analysis of films deposited at different positions in the reactor showed variations in the phase composition and in the total carbon content. In general, molybdenum-rich films of metallic Mo and Mo 2 C were deposited the first 3 cm after mixing of the reactants Žthe gases were mixed approx. 30 cm into the reactor tube.. These films always exhibited a large grain size. The nanocrystalline ␦-MoC 1y x phase was only deposited further into the reactor. In a previous study, this effect was explained by differences in reactivity between MoCl 5 and C 2 H 4 w16x. MoCl 5 is a very reactive compound and reacts easily with H 2 above 500⬚C. In contrast, C 2 H 4 needs to be ‘activated’ by thermal excitation andror homogeneous reactions prior to decomposition w16x. This will lead to a depletion of MoCl 5 in the vapour and increase the relative amount of hydrocarbons further downstream in the reactor. ␦-MoC 1y x films deposited at different positions in
Fig. 3. C1s XPS spectra of ␦-MoC 1y x films deposited at different positions in the reactor. Ža. Substrate position 5 cm and Žb. 9 cm after gas mixing. C 2 H 4 content: 12.5%. H 2 content: 25%.
the reactor changed from a metallic to an almost black nuance with increasing distance from the reactor entrance suggesting an increase in the carbon content downstream in the reactor. This was confirmed by XPS showing a peak at approximately 284.3 eV, which can be attributed to C᎐C andror C᎐H bonds. In the following this carbon is denoted ‘free carbon’. In addition, a C᎐Mo peak was seen at approximately 283.2 eV in all the films Žsee Fig. 3.. The intensity of the free-carbon peak decreased with increasing sputtering time and finally reached a constant value for the ratio of free carbon to carbidic carbon approximately 5᎐30% Žwhich corresponds to approx. 2᎐11 at% free carbon concentration. after 10 min sputtering. The results show that the films contain more free carbon in the surface region than in the bulk. Furthermore, the carbon content in the surface region and in the bulk increased with increasing distance from the reactor entrance ŽFig. 4.. At present, we are unable to identify whether the free carbon consists of graphite or some other form of carbon. However, no graphite has been observed by XRD. The detection limit for graphite is estimated to be approximately 10%. A rough estimation shows that the amount of ‘free carbon’ observed in the XPS analy˚ thick ses corresponds to a graphite layer less than 1 A ˚ if the grain size is assumed to be 100 A. As can be seen in Fig. 4 most of the films consist of carbidic carbon. XRD also showed that the cubic ␦MoC 1y x phase was the only carbide formed at a substrate position of more than 4 cm from the mixing of the gases. An interesting observation was that ␦MoC 1y x films deposited at different positions in the reactor have an almost identical cell parameter of 4.26 ˚ , which corresponds to a composition of approxiA mately ␦-MoC 0.67 using the data for cell parameter vs. carbon content given by Rudy et al. w19x. This result is in good agreement with the XPS analysis showing a composition of approximately MoC 0.68 ᎐ 0.70 using elemental standards and only including the contribution from the carbidic carbon.
J. Lu, U. Jansson r Thin Solid Films 396 (2001) 53᎐61
Fig. 4. Variation of free carbon content in the films as a function of sample position after gas mixing. Ža. 1 min sputtering; Žb. 3 min sputtering; and Žc. 10 min sputtering.
In general, the X-ray diffractograms of the ␦-MoC 1y x films exhibit very broad diffraction peaks Žsee Fig. 5.. The line-broadening can be assigned to several factors such as a small grain size and stresses in the film. A plot of Bcos vs. sin shows no indication of stresses. Most of the line-broadening can therefore be attributed to a small grain size. The average grain diameter was estimated to vary from approximately 50 to 170 ˚ using the Scherrer formula. The grain size was also A found to be dependent on the sample position ŽFig. 6.. In general, the grain size decreased further downstream in the reactor. However, the change in grain size was also dependent on the C 2 H 4 content. The largest variation was found for films deposited at low C 2 H 4 concentrations while those deposited at higher concentration only showed a minor change in grain size at different substrate positions ŽFig. 6.. 4.3. Influence of C2 H4 content The influence of C 2 H 4 on the deposition process was investigated in a series of experiments with C 2 H 4 concentrations ranging from 0.625% to 12.5% Žthe MoCl 5 and H 2 flows were kept constant.. XRD showed that all the films in this series consisted of ␦-MoC 1y x . The cell parameter of the ␦-MoC 1y x films was de-
Fig. 5. X-Ray diffractograms of ␦-MoC 1y x films deposited at Ža. 0.625% C 2 H 4 and Žb. 12.5% C 2 H 4 , respectively.
57
Fig. 6. Influence of sample position on grain size at different C 2 H 4 concentrations. Ža. 0.625% C 2 H 4 , Žb. 1.875% C 2 H 4 and Žc. 12.5% C2 H4.
˚ which corresponds to a compositermined to be 4.26 A tion of approximately ␦-MoC 0.67 for all films Ži.e. the carbide composition was independent of the C 2 H 4 concentrations .. Furthermore, the XPS results showed an unexpected relationship between free carbon content and vapour composition. It was found that the highest free carbon contents in the bulk Ž11 at.%. were observed in the films deposited at low ethene content Ž0.625%., while the film deposited with 12.5% ethane contained only 2 at.% free carbon. An interesting observation was that films deposited at ethene concentrations below 2% exhibited a Ž110. texture. In contrast, films deposited at ethene concentration above 3% were more or less texture-free Žcf. Fig. 5a,b.. Furthermore, it was also found the linebroadening of the XRD peaks was dependent on the ethene content in the vapour. A general observation was that the line-width of the diffraction peaks increased at higher ethene concentrations. From the Scherrer formula, the average grain sizes were esti˚ by decreasing the mated to vary from 50 to 170 A ethene content from 12.5 to 0.625% Žsee Fig. 7.. The deposition rate of the films was estimated from XRFS analyses. It was found that the deposition rate
Fig. 7. Influence of C 2 H 4 concentration on the grain size of a ␦-MoC 1y x film deposited at a sample position of 5 cm from the gas mixing.
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J. Lu, U. Jansson r Thin Solid Films 396 (2001) 53᎐61
Fig. 8. Deposition rate of the ␦-MoC 1y x phase as a function of grain size.
decreased with increasing ethene content from approximately 1.4= 10 19 Morcm2 h at 0.625% C 2 H 4 to approximately 6 = 10 18 Morcm2 h at 12.5% C 2 H 4 . A comparison between deposition rates and grain sizes shows that higher deposition rates yielded a larger grain size Žsee Fig. 8.. The influence of the ethene content on the film microstructure was also studied by TEM. Two samples deposited at 0.625% and 12.5% C 2 H 4 were chosen for the TEM observations. The TEM image Žsee Fig. 9. clearly shows that the film deposited with 12.5% C 2 H 4 consisted of almost texture-free nanocrystalline grains ˚ which are with diameters of approximately 50᎐100 A, in good agreement with the value obtained by XRD. In contrast, the microstructure of the films obtained with 0.625% C 2 H 4 is quite different Žsee Fig. 10.. At this ethene composition, the carbide exhibits a dendrite-like structure with Ž110. texture Žsee the SAD inset in Fig. 10.. It was found that the dendrite-like microstructure exhibited a larger distribution of grain sizes. Moreover, the dendrite-like film also contained a lot of pores. The porosity was estimated to be approximately 50% by
Fig. 10. TEM micrograph of nanocrystalline ␦-MoC 1y x film with a dendrite-like microstructure deposited at 0.625% C 2 H 4 .
comparing the film thickness shown by cross-section TEM image with the thickness calculated from XRFS analysis. In contrast, the porosity of the texture-free film in Fig. 9 was found to be negligible. 5. Discussion The results above show that nanocrystalline ␦MoC 1y x films can be deposited from a MoCl 5r C 2 H 4rH 2 mixture. It is also clear that we, within certain limits, can control the average grain size of the films. To explain the mechanism behind the nanocrystalline growth behaviour we need to determine how the carbide phase is formed. One mechanism for carbide formation in CVD is that the carbide is formed directly by a simultaneous addition of metal and carbon by surface reactions involving adsorbed hydrocarbon and metal halide species. The final grain size will then be defined by the nucleation kinetics of the carbide grains. Several authors, however, have presented an alternative growth models for carbide CVD using halides as metal sources, which are based on the assumption that small grains of the metal initially are formed by a hydrogen-reduction step followed by a carbonisation of the metal by, e.g. a hydrocarbon: MeX n Ž g. q nr2 H 2 Ž g. ª Me Ž s . q nHXŽ g. Me Ž s . q C x H y Ž g. ª MeC Ž s . q H 2 Ž g.
Fig. 9. TEM micrograph of nanocrystalline ␦-MoC 1y x film deposited with 12.5% C 2 H 4 .
In principle, the metal grains can be formed either by a heterogeneous process on the carbide surface or by a homogeneous reaction in the vapour. In the latter case, the metalrcarbide grain will be deposited on the surface of the growing carbide film yielding a powdery deposit. In both cases, however, the size of the final
J. Lu, U. Jansson r Thin Solid Films 396 (2001) 53᎐61
carbide grains will be defined by the size distribution of the initially formed metal grains. This type of growth behaviour has been reported in processes where the main purpose has been to produce a carbide powder for further processing. Such processes are usually carried out at high pressures where the supersaturation and the nucleation rate of the metal is high w20,21x. A similar mechanism has also been reported by Hogberg ¨ et al. for WC CVD on a reactive substrate material Ži.e. Ta. using WF6 as a metal source w22x. In this case, the substrate reduced WF6 under the formation of a metallic W film, which was carbonised in a later stage. Our results, however, clearly show that this simple growth mechanism can be excluded. For example, deposition experiments without C 2 H 4 Ži.e. with a MoCl 5rH 2 gas mixture. yielded only films of metallic Mo with a rather large grain size Žapprox. 1 m.. Furthermore, rather coarse-grained films of metallic Mo and Mo 2 C were generally deposited the first 3 cm after mixing of the reactants, while more carbon-rich, nanocrystalline ␦MoC 1y x films were deposited after a certain distance into the reactor where the relative concentration of C 2 H 4 is higher due to the depletion of MoCl 5 . Thus it can be concluded that the nanocrystallinity is correlated to the presence of C 2 H 4 in the vapour. This is also clearly seen in Fig. 7, which shows that the grain size is strongly influenced by the C 2 H 4 concentration. As discussed above the nanocrystalline growth can be due to two main factors: Ži. high supersaturation; and Žii. reduced surface diffusion lengths. In many CVD processes, nanocrystalline films are caused by a high supersaturation resulting in a very fast homogeneous nucleation step in the gas phase. The film is then formed by particles ‘snowing’ towards the substrate surface yielding a very powdery and porous film microstructure. This growth behaviour, however, is rather unlikely in our process since we have used a rather low pressure Ž1.7 torr. and a diluted reaction gas mixture. We have also not visually observed any particle formation in the gas phase or seen any indications of particle transport after the deposition zone. Furthermore, most of our films Žwith exception of those deposited at lower C 2 H 4 contents. exhibit a dense microstructure without pores or agglomerated particles. These results suggest that the nanocrystalline ␦-MoC 1y x films be formed by an extensive heterogeneous nucleation on the substrate andror the growing film surface. We believe, however, that a high supersaturation alone cannot be responsible for the high nucleation rates required for the nanocrystalline growth. This conclusion is supported by the thermodynamic calculations which showed that ⌬G Žthe driving force. for the process will increase with increasing C 2 H 4rMoCl 5 ratios in the vapour Žsee Fig. 2.. If the nanocrystalline growth is caused by only a high supersaturation, an increase in the C 2 H 4 content should lead not only to smaller grain sizes but also to a
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higher deposition rate as shown by curve a in Fig. 1. The results in Fig. 8, however, show an opposite behaviour suggesting that some other mechanism than a high supersaturation must be responsible for the finegrained microstructure. The second possibility is that the small grain size in the ␦-MoC 1y x films is caused by reduced surface diffusion lengths. Our results are consistent with such a growth mechanism. This conclusion is supported by, for example, Fig. 8, showing a strong correlation between smaller grain sizes and reduced deposition rates. This relationship is in good agreement with a growth process controlled by a reduced diffusion lengths as shown by curve b in Fig. 1. The results above also suggest that the presence of C 2 H 4 is critical for this behaviour Žsee, e.g. Fig. 7.. This conclusion is also supported by the XPS analyses which show that all nanocrystalline films contain free carbon Žor more correct carbon with mainly C᎐C andror C᎐H bonds. at the film surface and in the grain boundaries. We therefore suggest that the nanocrystalline growth behaviour of the ␦-MoC 1y x films is due to a ‘poisoning’ of the carbide surface by strongly adsorbed reaction products from the decomposed C 2 H 4 . A likely mechanism can then be as follows: small carbide nuclei are formed on the surface of the substrate or on the growing film. The surfaces of the nuclei are quickly covered by strongly adsorbed species from the decomposing C 2 H 4 , which reduces the surface diffusion rates of the reactants. As a consequence the growth rate of the nuclei will be low. As further reactants are adsorbed, the supersaturation on the surface will increase leading to the nucleation of new carbide nuclei. The final film will exhibit a microstructure with nanometre-sized ␦-MoC 1y x grains with some free carbon at the grain boundaries. At present, we have no clear evidence for the exact nature of these carbon-containing reaction products. A possibility is that a graphitic-like overlayer is formed on the carbide surface. As mentioned above the free carbon ˚ thick if corresponds to a graphite layer less than 1 A ˚ grain size is assumed to be 100 A. This thickness is less than a monolayer of a graphite and show that no continuous graphite phase can exist in the boundary regions of the nanocrystalline grains. Thus it is clear that further studies with, e.g. other in situ surface analysis techniques andror high-resolution TEM is required to determine the nature of the free carbon and its importance for the observed growth behaviour. The model for nanocrystalline growth proposed above is in agreement with observations in the literature. For example as mentioned in Section 1, it is well known that molybdenum carbide surfaces in many respects exhibit large similarities with many platinum metals w4x. It is also known that so-called ‘carbonaceous overlayers’ are formed on platinum metal surfaces in the presence of hydrocarbons w23x. These overlayers are
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catalytically active at lower temperatures. At higher temperatures, however, they loose hydrogen under the formation of ‘graphite-like’ carbon on the surface. The presence of such ‘graphitic’ overlayer inhibits many surface reactions and makes the surface catalytically inactive. In addition, there is some experimental evidence that similar effects can be also found on molybdenum carbide surfaces. Carburisation experiments of Mo with hydrocarbons have shown that a carbon overlayer with graphitic-like characters Žas shown with Auger electron spectroscopy. is formed at elevated temperatures and high hydrocarbon exposures w24x. The presence of such an overlayer was found to drastically reduce the rate of the carbonisation process. Evidence of carbon-containing overlayers has also been observed in the CVD of boron carbides on transition metal substrates w25x. It was found that the nucleation and initial deposition rates of the boron carbide were strongly reduced by some hydrocarbons Že.g. C 2 H 4 .. The effect, which was observed on Mo substrates but not on other metals such as Ti, was attributed to the formation of a carbon-containing overlayer on the surface w25x. A growth mechanism controlled by a low surface mobility and reduced surface diffusion lengths can also explain other experimental observations. The deposition profile described in Section 4.2, for example, shows that molybdenum-rich phases Ži.e. Mo and Mo 2 C. with a large grain-size are deposited at the entrance of the reactor and that nanocrystalline ␦-MoC 1y x films only are formed at some distance into the reactor. This observation together with the fact that the grain size is reduced with increasing distance from the reactor entrance suggest that the surface blocking is more effective further downstream in the reactor Žsee Fig. 6.. This can attributed to the fact that many hydrocarbons have to be ‘activated’ by, e.g. thermal excitation andror homogeneous reactions prior to decomposition at the fairly low temperature used in our experiments Ž800⬚C.. The activation requires a certain time in the heated zone of the furnace and leads to strong variations in the growth conditions at different positions in the deposition reactor. This also suggests that a chemically more active carbon source should give a quite different deposition profile compared to C 2 H 4 . An example of such a precursor is CH 2 I 2 , which rapidly decomposes into I and CH 2 radicals upon heating. In a previous study, we have also shown that carbide CVD from a MoCl 5rCH 2 I 2rH 2 mixture yields nanocrystalline ␦MoC 1y x film in all parts of the reactor without any formation of large-grained Mo and Mo 2 C at the reactor entrance and with minor variations in grain sizes and free carbon content along the reactor w10x. This can be explained by the fact that the CH 2 radicals are more effective in forming the carbon-containing overlayer which reduces surface mobility.
Our results also show that hydrogen plays an important role for the nanocrystalline growth. In an earlier study, it was shown that CVD from a MoCl 5rC 2 H 4 mixture Ži.e. no H 2 added to the vapour. yielded an initial deposition of the ␦-MoC 1y x phase with an aver˚ After a age grain size of approximately 250᎐300 A. while, however, the phase composition changed as ␥⬘MoC 1y x with a fairly large grain size was formed on top of the ␦-MoC 1y x film w16x. An addition of approximately 10% hydrogen, however, changed the phase composition to the cubic ␦-MoC 1y x phase which always exhibited a nanocrystalline microstructure. The observations can be explained by the fact that hydrogen is likely to be involved in different homogeneous andror heterogeneous reaction steps in the decomposition of C 2 H 4 . It is therefore possible that the presence of H 2 in the vapour is required for the formation of a carbon-containing overlayer which can effectively reduce the surface mobility and diffusion lengths during growth. Finally, it is conceivable to assume that the formation of a carbon-containing overlayer which reduces the surface mobility of adsorbed reactants is dependent on the crystallographic orientation of the carbide surface. Some surface orientations are probably more easy to ‘poison’ than other. This should give rise to a film texture where some grains grow faster than other grains with unfavourable orientations. In fact, the XRD and TEM studies showed that films with a Ž110.-texture and a dendrite-like growth behaviour were deposited at low C 2 H 4 contents in the vapour Žsee Figs. 5 and 10.. These films also exhibited with a high degree of porosity and a large size distribution of the carbide grains. In contrast, films deposited at higher C 2 H 4 concentrations were almost texture-free, rather dense and exhibited a more narrow grain-size distribution. The growth behaviour at high C 2 H 4 concentrations can be explained by the fact that all carbide surface irrespective of crystallographic orientation is covered by the carbon-containing overlayer. Films deposited at these conditions will therefore be texture-free. At low C 2 H 4 concentrations, however, it is possible that some crystallographic orientations are less affected by this overlayer thereby giving rise to a textured film. The porosity of films deposited at 0.625% C 2 H 4 also explains the high contents of free carbon in these films. The large total surface area in the porous films will contain large amounts of free carbon Žsee Figs. 3 and 4. which will give an apparent high ‘bulk concentration’ also after sputtering. 6. Conclusions CVD of molybdenum carbides from a MoCl 5r H 2rC 2 H 4 mixture at 800⬚C leads to the formation of nanocrystalline ␦-MoC 1y x films. The average grain size,
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free carbon content and the deposition rate are affected by the experimental parameters. The experiments clearly show that high C 2 H 4 concentrations favour the formation of nanocrystalline grains. Our results suggest that a high supersaturation is an unlikely explanation for the observed growth behaviour. We have therefore proposed a model where the nanocrystalline growth is caused by a reduced surface mobility induced by the presence of some strongly adsorbed reaction products from C 2 H 4 . Hitherto, this explanation is based on general growth trends Že.g. Figs. 7 and 8.. Further studies using, for example, in situ analyses with different surface science techniques under well-defined conditions Ži.e. UHV. are required to elucidate the exact nature of such carbon-containing reaction products and their influence on the growth process. It is clear, however, that the proposed growth model is in agreement with many observations in the literature. Acknowledgements The Swedish Research Council for Engineering Sciences ŽTFR., the Swedish Natural Science Research ˚ Council ŽNFR., and the Angstrom ¨ Consortium are acknowledged for financial support. References w1x w2x w3x w4x
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