Low temperature epitaxial growth of metal carbides using fullerenes

Surface and Coatings Technology 142᎐144 Ž2001. 817᎐822

Low temperature epitaxial growth of metal carbides using fullerenes a , J.-P. Palmqvist a , L. Norina , J.O. Malmb, U. Jansson a,U , H. Hogberg ¨ L. Hultman c , J. Birch c

˚ Department of Inorganic Chemistry, Angstrom ¨ Laboratory, Uppsala Uni¨ ersity, P.O. Box 538, SE-751 21, Uppsala, Sweden Department of Inorganic Chemistry2, National Center of HREM, Lund Uni¨ ersity, P.O. Box 124, SE-22100, Lund, Sweden c Department of Physics, Thin Film Physics Di¨ ision, Linkoping Uni¨ ersity, SE-581 83, Linkoping, Sweden ¨ ¨

a b

Abstract Epitaxial transition metal carbides can be deposited at low temperatures by simultaneous evaporation of C 60 and either metal e-beam evaporation or metal d.c. magnetron sputtering. Hitherto, epitaxial films of TiC, VC, NbC, MoC, W2 C and WC have been deposited on MgOŽ100., MgOŽ111. and in some cases 6H- and 4H-SiCŽ0001.. Epitaxial TiC films with a good quality have been deposited at temperatures as low as 100⬚C with metal sputtering, while somewhat higher temperatures Ž) 200⬚C. are required for the other metals. In general, the plasma-assisted process allows lower deposition temperatures than the co-evaporation process. Most carbides can be deposited in a wide range of compositions within their homogeneity ranges by a fine-tuning of the MerC 60 flux. However, the results suggest that the formation of free surface carbon can be a limiting factor. The processes have also been used to deposit superlattices of TiCrNbC and TiCrVC at 400᎐500⬚C as well as epitaxial ternary Ti xV1yx C y films. Furthermore, epitaxial films of ternary carbides with well-controlled metal concentration profiles can be deposited at temperatures below 500⬚C. 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Metal carbides; Epitaxy; C 60 ; Sputtering; Evaporation

1. Introduction The transition metal carbides have many interesting properties. The most well-known application is as a wear-resistant material in, e.g. thin film coatings on cemented carbide inserts. The mechanical properties are strongly dependent on the microstructure and can be controlled by a careful tuning of chemical composition, phase composition, grain size, etc. For example, it has recently been demonstrated that nanocrystalline carbide films with small amounts of amorphous carbon may exhibit a combination of high hardness, low friction and a high toughness w1x. Also epitaxial carbide films may exhibit interesting properties. For example,

U

Corresponding author. Fax: q46-18-513548. E-mail address: [email protected] ŽU. Jansson..

layered epitaxial structures Žsuperlattices . of transition metal nitrides Že.g. TiNrVN and TiNrNbN. are known to have unique mechanical properties including a substantial increase in strength and hardness w2x. It is possible that superlattices of monocarbides ŽNaCl structure-type . of group 4 and 5 metals exhibit similar or even superior properties to nitride superlattices. Less well-known is the fact that the transition metal carbides are also very good electrical conductors with resistivity values of 20᎐70 ␮⍀ cm. Recently, it has also been demonstrated that metal carbides may be potential candidates as contact materials in SiC-based microelectronics w3x. The high melting points of the metal carbides makes it difficult to deposit epitaxial films at low temperatures. Most chemical vapour deposition ŽCVD. and physical vapour deposition ŽPVD. processes yield fine-grained andror amorphous films at temperatures

0257-8972r01r$ - see front matter 䊚 2001 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 7 - 8 9 7 2 Ž 0 1 . 0 1 1 1 1 - 2

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U. Jansson et al. r Surface and Coatings Technology 142᎐144 (2001) 817᎐822

below 800⬚C. However, we have recently demonstrated that C 60 is an excellent carbon source in PVD of carbides Žsee, e.g. w4᎐8x.. The films can be deposited either by co-evaporation or by metal sputtering combined with evaporation of C 60 . The process allows epitaxial growth at temperatures as low as 100⬚C on suitable substrates. This paper presents a review of the most important results from this project including the growth of single-phase binary and ternary films superlattices and gradient structures.

Cross-section samples for TEM and HREM studies were prepared by attaching two carbide coatings together with epoxy and sectioning through the glue joint. The sections were first dimple-polished to ; 10 ␮m and finally thinned by ion milling to transparency using PIPS ŽGatan. at a 4⬚ angle.

3. Results and discussion 3.1. Carbide formation

2. Experimental The carbide films were deposited using two different ultra high vacuum ŽUHV. processes: Ži. co-evaporation of C 60 and transition metals; and Žii. evaporation of C 60 and magnetron sputtering of the metal using an Ar pressure of approximately 0.1 Pa. In both processes, the C 60 ŽMER 99.95%. was introduced into the chamber by a Knudsen effusion cell. The vapor pressure of C 60 in the cell was approximately 0.04᎐0.25 Pa. In the co-evaporation process, the metals were evaporated from standard e-beam evaporators using a metal rod while conventional d.c. magnetrons were used in the plasma-assisted process. MgOŽ100., MgŽ111., SiŽ100. and 4H- and 6H-SiCŽ0001. have been used as substrate materials. Details on the experimental parameters, the quality of the source materials and substrate cleaning w7,8x. can be found in Hogberg ¨ The carbide films were characterized by X-ray diffraction ŽXRD.. Epitaxial relationships to the substrate material were determined by X-ray azimuthal ␾-scans, using a four circle Siemens D 5000 equipped with point focus. Furthermore, the films were investigated by X-ray photoelectron spectroscopy ŽXPS. performed on a PHI 5500 multitechnique instrument using monochromatic AlK ␣ radiation. The microstructure was investigated with high resolution electron microscopy ŽHREM. ŽJEM4000EX. operated at 400 kV.

The interaction between co-evaporated C 60 and transition metal atoms on a surface is strongly dependent on the type of metal and the MerC 60 flux ratio. This can be seen in Fig. 1a which shows the C1s XPS peak for a typical carbide-forming transition metal Žvanadium. at different film compositions. C1s spectra from films with a low metal content show only one peak at approximately 285 eV which can be attributed to C᎐C bonds. Further studies with Raman and X-ray spectroscopies show that the C 60 cages in these films are intact. The films, which are amorphous and very oxidation-sensitive, can be described as a new type of transition metal fulleride compounds w9,10x. When the metal content in the film is gradually increased Žby changing the MerC 60 flux ratio., a second additional peak can be seen at lower binding energies Žapprox. 282.5 eV in Fig. 1a.. This peak can be attributed to a metal carbide. The carbide formation can be explained by a charge transfer from the metal to antibonding Ž ␲U . orbitals in the C 60 molecule. This results in a weakening of the intramolecular bonds in the fullerene. At a critical metal concentration the cage starts to decompose and a carbide is formed. Hitherto, we have investigated a number of different carbide-forming transition metals ŽTi, V, Nb, Mo and W. and found that they all exhibit similar behavior towards C 60 . This can be seen in Fig. 1b which shows that the three carbideforming metals Ti, V and Nb begin to form carbides at

Fig. 1. Ža. C1s XPS spectra from films deposited at 100⬚C on SiŽ100. with vanadium as metal source; and Žb. carbidic carbon content vs. metal content Žoverall concentration in at.%. for Ti, V, Nb and Co. Results obtained from C1s spectra. The lines in the figure are guidelines for the eye.

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Table 1 List of epitaxial metal carbides deposited with C 60 as the carbon source a Carbide

TiC

Substrate

MgOŽ100., MgOŽ111., 4HSiCŽ0001., 6HSiCŽ0001. MgOŽ100.

VC NbC MoCŽcubic. WCŽcubic. W2 CŽhex.

MgOŽ100. MgOŽ100., MgOŽ111. MgOŽ100., MgOŽ111. MgOŽ111.

TiCrVCU

MgOŽ100.

TiCrNbCU Ti1y xVx C Mo1y x Nbx C

MgOŽ100. MgOŽ100. MgOŽ111.

aU

Technique

Min. temp. for

Max.

epitaxy Ž⬚C.

˚. Thickness ŽA

Co-evaporation Sputtering

250 Žco-evap.. 100 Žsputtering.

) 5000

w4᎐8x

Co-evaporation Sputtering Co-evaporation Co-evaporation

400 Žco-evap.. 200 Žsputtering. 400 N.D

1500

w5᎐8x

200 N.D

w7x w11x

Sputtering

N.D

N.D

w11x

Sputtering

N.D 400 Žco-evap.. ) 200 Žsputter. 500 400 N.D

N.D

˚ - 7000 A

w11x w6᎐8x

N.D N.D N.D

w7x w11x w11x

Co-evaporation, Sputtering Co-evaporation Co-evaporation Co-evaporation

Ref.

Denotes superlattice structures, N.D, Not determined.

approximately 5᎐10 at.% metal and the amount of formed carbide as a function of metal content is similar for all three metals. In contrast, cobalt shows a quite different behaviour. This metal forms only metastable carbides and the co-evaporation of C 60 and Co yields no carbide for Co contents below approximately 50 at.%. We believe that other transition metals which cannot form stable carbides, such as Fe and Ni, would also exhibit a similar behaviour. 3.2. Epitaxial film growth in the co-e¨ aporation process Films deposited at 100⬚C with the co-evaporation process have a cubic structure ŽNaCl-type. and are always polycrystalline. The XRD peaks, however, are very broad, which suggests an extremely fine-grained microstructure. The grain sizes have been estimated to ˚ using the Scherrer formula w8x. XPS analysis 60᎐100 A show that the bulk contains small amounts of free carbon together with the carbide and that their microstructure therefore can be described as nanocrystalline carbide grains in an amorphous carbon matrix. With increasing deposition temperature, the microstructure rapidly changes and epitaxial carbide films are formed at a certain temperature providing that the film is grown on a suitable substrate material. Table 1 shows a summary of epitaxial film structures deposited in this project. As can be seen the critical temperature for epitaxy and the maximum attainable thickness are dependent on the metal. In general, TiC has the highest quality followed by VC and NbC. Epitaxial TiC can be deposited at 250⬚C with thicknesses greater than ˚ The deposition rate of TiC is typically 0.2᎐0.3 5000 A. ␮mrh but no attempts have been made to optimize the deposition rate. Epitaxial growth of VC and NbC re-

quires at least 400⬚C and a polycrystalline film starts to ˚ respectively. form after approximately 1500 and 200 A, The maximum thickness and overall quality of the films can be improved by a temperature increase. For example, the maximum thickness of epitaxial NbC can be ˚ at a deposition temincreased to approximately 400 A perature of 500⬚C. The epitaxial films have been characterized with XRD and HREM. Typically, a ␪᎐2␪ scan of epitaxial carbide films deposited on MgOŽ100. only show  n004 reflections and ␾-scans clearly show that the films are epitaxial with the orientation MeCŽ001.rrMgOŽ001. and with an in-plane orientation MeCw100xrrMgOw100x Žfor a more detailed description of the XRD analyses, w7x.. HREM studies of the filmrsubstrate see Hogberg ¨ interfaces show a very good film registry to the substrate Žsee Fig. 2.. No grain boundaries have been

Fig. 2. HREM image of the interface Žmarked with arrows. between ˚ thick VC film and MgOŽ100.. a ; 900-A

820

U. Jansson et al. r Surface and Coatings Technology 142᎐144 (2001) 817᎐822

observed in any of the studied films. A more detailed study of the interfaces, however, show that misfit dislocations have been formed Ži.e. the films have started to relax.. The density of the misfit dislocations is directly related to the lattice mismatch. NbC with an approximate misfit of approximately 5% to MgOŽ100. has an average distance between the dislocations of approxi˚ In contrast, TiC and VC with a misfit of mately 40 A. approximately 2 and ᎐1.5%, respectively, shows very few dislocations on each HREM sample Žthe dislocation density is too low to be calculated. w7x. Epitaxial TiCŽ111. films have also been deposited on 6H- and 4H-SiCŽ0001. at temperatures as low 250⬚C w5,12x. HREM of these films, however, show that the carbide grows with columnar grains in the w111x direction. The grains seem to originate from the surface steps on the SiC substrate. Furthermore, the grains exhibit twin relations which can be attributed to two different stacking sequences w5x. The epitaxial TiC films exhibit excellent properties as Ohmic contact materials in SiC microelectronics w3x. Monocarbides with the NaCl structure usually exhibit large homogeneity ranges caused by carbon vacancy formation. A good example is TiC which is stable from ; TiC 0.5 to ; TiC 0.97 . A general observation in all our studies is that the epitaxial carbide films are substoichiometric and that the composition range for epitaxial growth is more narrow than predicted from the phase diagram. For example, at 400⬚C epitaxial titanium carbide films can only be deposited in the composition range of TiC 0.5 ᎐TiC 0.85 . All attempts to deposit epitaxial films with higher carbon content failed and yielded only polycrystalline films. A similar behavior was observed for NbC and VC where the composition ranges for epitaxy were determined to NbC 0.7y0.8 and VC 0.8y0.9 , respectively. XPS analyses suggest that the narrow composition range for epitaxy can be due to the formation of free carbon on the surface. This can be seen in Fig. 3, which shows a C1s spectra from an epitaxial NbC 0.7 film with only one carbide peak at 282.7 eV. When the C 60 rNb flux ratio is increased in an attempt to deposit a more carbon-rich film, the amount of surface free carbon starts to increase to approximately 30 at.% of the total carbon content Žtop spectrum in Fig. 3.. The composition of the films is now approximately NbC 0.8 and an XRD analysis shows that the film is polycrystalline. A similar behavior has also been observed for the other carbides.

Fig. 3. C1s XPS spectra from an epitaxial NbC 0.7 film Žbottom. and a polycrystalline NbC 0.8 film Žtop..

film quality than for the corresponding co-evaporation process. For example, the minimum temperature for TiC and VC epitaxy on MgOŽ001. can be reduced to 100 and 200⬚C, respectively, when the metal is sputtered w8x. Furthermore, nanocrystalline films are also formed in the plasma-assisted process below the minimum temperature for epitaxy. However, the grain size in these films Žas estimated by applying the Scherrer formula on the XRD peak widths. is approximately twice as large as in co-evaporated films w8x. The mechanisms for the apparent improvement of film quality for sputtering vs. evaporation of the metal species are currently not known. Obvious differences are in the higher deposition rate control and ion-assisted deposition conditions of the sputtering source. Recently, the metal sputtering process has been used to deposit WC films w11x. An interesting observation is that epitaxial films of cubic WC ŽNaCl structure-type . can be deposited on MgOŽ100.. This is a high-temperature phase in the W᎐C system and is only thermodynamically stable above 2500⬚C. It is likely that this phase is stabilized by the low filmrsubstrate mismatch Žy1.5᎐0.5%.. Hitherto, all attempts to deposit epitaxial films of hexagonal WC have failed. It has been possible, however, to deposit epitaxial films of hexagonal W2 C on both MgOŽ100. and MgOŽ111. w11x. 3.4. Carbide superlattices

3.3. Epitaxial carbide film growth by C6 0 e¨ aporation and metal sputtering Epitaxial carbide films have also been deposited by d.c. magnetron sputtering of the metal together with C 60 evaporation from a Knudsen cell Žsee Table 1.. Typically, the sputtering process yields films with higher

Both the co-evaporation process and the sputtering process have been used to deposit carbide superlattices by using individually controlled shutters in front of the e-beam evaporators and magnetrons. As can be seen in Table 1, both TiCrVC and TiCrNbC superlattices have been deposited on MgOŽ100. by the co-evapora-

U. Jansson et al. r Surface and Coatings Technology 142᎐144 (2001) 817᎐822

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Fig. 4. Cross-sectional low-resolution TEM image Žlower part of picture. showing the contrast between individual layers of TiC and NbC. The ˚ TiCq ; 50 A ˚ NbC. superlattice. The upper left image show the interface between the HREM images show two regions of the 5 = Ž; 60 A MgOŽ001. substrate and the first epitaxial TiC layer Žmarked with arrows.. The upper right image shows the fifth TiCrNbC interface Žmarked with arrows.. In the region misfit dislocations are present Žedge dislocation. marked with T in the picture.

tion process. The TiCrNbC superlattice requires a higher deposition temperature Ž500⬚C. than TiCrVC ˚ TiCq 50 Ž400⬚C.. A HREM cross-section of a 5 = Ž60 A ˚ . A NbC superlattice is shown in Fig. 4. The upper right image shows that the interface between the fifth TiC and NbC layers. As can be seen this interface is semicoherent since misfit dislocations are present in this region. The sputtering process yields superlattices with a higher quality than the co-evaporation process w8x. TiCrVC superlattices can be deposited already at 200⬚C with the plasma-assisted process but the quality is rather low. At 400⬚C, the quality of the sputtered TiCrVC superlattice Žas deduced from the superlattice satellites in the diffractograms. is considerably higher than for corresponding superlattice deposited by co˚ thick TiCrVC evaporation. Hitherto, more than 4000-A superlattices with modulation wavelengths of 72 and ˚ have been deposited. 208 A 3.5. Epitaxial ternary carbide films Most monocarbides exhibit a broad homogeneity range for solid solution of other metals. For example, the TiC and VC systems exhibit a complete miscibility which makes it possible to deposit ternary, Ti 1y xVx C films. Preliminary studies on this system show that epitaxial films easily can be obtained on, e.g. MgOŽ100. w11x. The cell parameter of this ternary compound can

be fine-tuned to exactly match the substrate. Moreover, the relative deposition rate of V and Ti can be changed during the deposition process making it possible to obtain gradient Ti 1y xVx C films with, e.g. xs 0 close to the substrate and xs 1 at the top of the film. 3.6. Mechanisms for epitaxy The mechanism behind the low temperature for epitaxial growth Ž; 10% of the melting temperature. is yet unknown. One possibility is that C 60 and fragments of this molecule have such high mobility on carbide surfaces that epitaxy can be maintained also at low temperatures. A more likely explanation, however, may be the formation of metal᎐carbon clusters. Recent studies have shown that C 60 reacts with group 4᎐5 metal atoms in the vapour under the formation of small metal᎐carbon clusters such as Ti 8 C 12 and V8 C 12 Žmetcars. Žsee, e.g. Tast et al. w12x.. Many of these metcars exhibit a polyhedral framework without cubic symmetry but cubic cluster structures such as Ti 14 C 13 and Nb4 C 4 have also been observed. It is likely that such metal᎐carbon clusters are formed also on the surface of a growing carbide film. Cubic clusters can certainly be perfect building blocks and the assembly of such clusters may be the main reason for the epitaxial growth at low temperatures. However, further studies, with, e.g. HREM are required to determine the mechanism for epitaxy.

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4. Concluding remarks Epitaxial carbide films and superlattices can be deposited by both a co-evaporation process and by a sputter process. It is clear that the plasma-assisted process yields films with a better quality and allows epitaxial growth at lower temperatures. The deposition processes have yet not been optimized and it is likely that a fine-tuning of the experimental parameters can push the minimum temperature for epitaxy even lower and also make it possible to deposit stoichiometric films. The films can have a potential use in, e.g. SiCbased microelectronics.

Acknowledgements The Swedish Research Council for Engineering Sciences ŽTFR., the Swedish Natural Science Research Council ŽNFR. and the Swedish Foundation for Strategic Research ŽSSF. are acknowledged for financial support. Jorg ¨ Neidhardt and Ian Brunell are also acknowledged for experimental assistance and discussions.

References w1x A.A. Voevodin, J.P. O’Neill, S.V. Prasad, J.S. Zabinski, J. Vac. Sci. Technol. A 17 Ž1999. 986. w2x S.A. Barnett, A. Madan, Physics World 11 Ž1998. 45. ¨ w3x S.-K. Lee, C.-M. Zetterling, M. Ostling, J.-P. Palmquist, H. Hogberg, U. Jansson, Solid-State Electron. 44 Ž2000. ¨ 1179᎐1186. w4x L. Norin, S. McGinnis, U. Jansson, J.O. Carlsson, J. Vac. Sci. Technol. A 15 Ž1997. 3082᎐3085. w5x L. Norin, J. Lu, J.O. Malm, U. Jansson, J. Mater. Res. 14 Ž1999. 1589᎐1596. w6x L. Norin, H. Hogberg, J. Lu, J.O. Malm, U. Jansson, Appl. ¨ Phys. Lett. 73 Ž1998. 2754᎐2756. w7x H. Hogberg, J.O. Malm, A. Talyzin, L. Norin, J. Lu, U. Jansson, ¨ J. Electrochem. Soc. 147 Ž2000. 3361. w8x H. Hogberg, J. Birch, M.P. Johansson, L. Hultman, U. Jansson, ¨ J. Mater. Res. 16 Ž2001. 633. w9x L. Norin, U. Jansson, C. Dyer, P. Jacobsson, S. McGinnis, Chem. Mater. 10 Ž1998. 1184᎐1190. w10x L. Qian, L. Norin, J.H. Guo et al., Phys. Rev. B 59 Ž1999. 12667᎐12671. w11x J.P. Palmqvist, H. Hogberg, J.O. Malm et al. unpublished ¨ results. w12x F. Tast, N. Malinowski, S. Frank, M. Heinebrodt, I.M.L. Billas, T.P. Martin, Z. Phys. D 40 Ž1997. 351.