Reactive magnetron sputtering of molybdenum sulfide ...

Reactive magnetron sputtering of molybdenum sulfide thin films: In situ synchrotron x-ray diffraction and transmission electron microscopy study V. Weiss, W. Bohne, J. Röhrich, E. Strub, U. Bloeck et al. Citation: J. Appl. Phys. 95, 7665 (2004); doi: 10.1063/1.1736323 View online: http://dx.doi.org/10.1063/1.1736323 View Table of Contents: http://jap.aip.org/resource/1/JAPIAU/v95/i12 Published by the AIP Publishing LLC.

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JOURNAL OF APPLIED PHYSICS

VOLUME 95, NUMBER 12

15 JUNE 2004

Reactive magnetron sputtering of molybdenum sulfide thin films: In situ synchrotron x-ray diffraction and transmission electron microscopy study V. Weiss, W. Bohne, J. Ro¨hrich, E. Strub, U. Bloeck, I. Sieber, and K. Ellmera) Hahn-Meitner-Institut, Solarenergieforschung and Ionenstrahllabor, Glienicker Straße 100, D-14109 Berlin, Germany

R. Mientus Opto-Transmitter-Umweltschutz-Technologie e.V., Ko¨penicker Straße 325b, D-12555 Berlin, Germany

F. Porsch University of Bonn, Institute of Mineralogy and Petrology, Poppelsdorfer Schloß, D-53115 Bonn, Germany

共Received 2 September 2003; accepted 10 March 2004兲 The nucleation and growth of magnetron sputtered MoSx films has been investigated by in situ energy dispersive x-ray diffraction, electron microscopy, and elastic recoil detection analysis. The MoSx films (0.5⭐x⭐2) were prepared by reactive magnetron sputtering from a molybdenum target in an argon–hydrogen sulfide mixture at substrate temperatures up to 700 °C. Using time-resolved in situ x-ray diffraction it was found that the films start to grow with 共001兲 orientation where the van der Waals planes are parallel to the substrate surface. Depending on the deposition conditions a crossover of texture to the 共100兲 orientation occurs, which leads to very rough surfaces. This texture crossover occurs earlier at low substrate temperatures and/or high deposition rates and/or high energetic particle bombardment of the growing films. The MoSx films exhibit significant lattice strain 共up to 4%兲 in the c direction, i.e., perpendicular to the van der Waals planes, which decreases at high substrate temperatures and/or low deposition rates. This lattice expansion is not caused by film stress. Instead, it seems to be connected with disturbed or turbostratic growth due to crystallographic defects induced by energetic bombardment of the films. Also, intercalation of hydrogen could be responsible for lattice expansion, since significant amounts of hydrogen were detected by elastic recoil detection analysis. The sulfur deficiency 共up to 20%兲 found in films that were deposited at temperatures higher than 200 °C could have been initiated by the reducing effect of the atomic hydrogen in our Ar/H2 S plasma. © 2004 American Institute of Physics. 关DOI: 10.1063/1.1736323兴

perpendicular to the substrate 关 c⬜ or 共001兲 texture; previously referred to type II as opposed to type I, which means 共100兲 texture兴. Most of the studies on films deposited by methods like aqueous chemical synthesis,6,8 annealing of multilayers diode sputtering,11 magnetron Mo/S/Mo/S,...,9,10 12–14 and metalorganic chemical vapor deposition sputtering, 共MOCVD兲15,16 have been performed on a very empirical basis. In most cases only a small number of parameters was varied or a limited parameter range was explored. The aim of our work was to better understand the growth mechanisms of these van der Waals type semiconductors. For van der Waals type semiconductors MoS2 and WS2 to be used as thin film solar cell absorber materials a film thickness of only about 100 nm is sufficient because of their high absorption coefficients. It is expected that even the initial film growth influences the electrical properties of MoS2 films. For these reasons a detailed in situ study of the growth was performed in the present investigation. Nucleation, grain growth, and texture formation have been analyzed by in situ energy dispersive x-ray diffraction 共EDXRD兲 during reactive magnetron sputtering. EDXRD is a well established method for in situ investigations of bulk reactions:17 investigation of high pres-

I. INTRODUCTION

Investigations of molybdenum disulfide (MoS2 ) thin films have been widely published because the material is a good lubricant,1–3 exhibits catalytic properties,4 and is possible for use as an absorber material in thin film solar cells.5 The last application has been proposed since MoS2 , as well as the structurally comparable selenide and corresponding tungsten compounds (WS2 ,WSe2 ), shows a high absorption coefficient and a suitable band gap (E G ⫽1.71 eV) in relation to the intensity distribution of the solar spectrum.6 The crystallographic structure of these layered materials is characterized mostly by van der Waals bonds between the sulfur layers, leading to a saturated surface of sulfur atoms without dangling bonds.7 Therefore, the charge carrier recombination of such surfaces or interfaces should be reduced significantly compared to that of covalently bonded semiconductors like silicon or gallium arsenide. This offers new opportunities for the preparation of thin film solar cells if the films can be prepared with the van der Waals bonded layers parallel, i.e., with the c axis a兲

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sure phase compressibility, phase growth, or change in phase,18 –20 syntheses under hydrothermal conditions,21–23 high-temperature hydration of cement,24 or studies of the charge–discharge cycles of lithium intercalation batteries.25 Recently, this method has been used for in situ investigations of thin film growth26 and sulfurization of metal precursor thin films.27 This article is organized as follows: Sec. II summarizes the experimental details; the EDXRD results and analysis are given in Sec. III together with electron microscopy data, film density, and stoichiometry measurements; Sec. IV summarizes our results. Preliminary data on textured MoSx growth has been already published.28 II. EXPERIMENTAL PROCEDURE

The MoSx thin films were deposited onto oxidized silicon 共100兲 substrates by reactive magnetron sputtering from a molybdenum target 共99.95%兲 in mixtures of argon 共99.999%兲 and H2 S 共99.5%兲. A balanced magnetron with a target diameter of 51 mm and a target-to-substrate distance of 60 mm was used in radio frequency 共rf兲 mode at 13.56 MHz or under dc excitation conditions. The process chamber had a base pressure of about 5⫻10⫺5 Pa. The total sputtering pressure was varied between 0.18 and 9 Pa and the H2 S content of the sputtering atmosphere varied from 9% to 78%. A boron nitride coated graphite heater allowed substrate temperatures up to 750 °C. The setup for the in situ x-ray diffraction analysis which was used at the synchrotron radiation source HASYLAB in Hamburg, Germany, has been described recently.26,29 In situ EDXRD spectra were recorded under diffraction angles ␪ of about 5°. The width of the primary x-ray beam was at least 100 ␮m, which is much larger than the film thickness. The peak area of the Mo K ␣ fluorescence line was used as a measure of the number of the Mo atoms 共atoms cm⫺2兲 in the film after calibration by an independent method, i.e., elastic recoil detection analysis 共see below兲. The chemical composition of the films was studied by Rutherford backscattering spectrometry 共RBS兲 and elastic recoil detection analysis 共ERDA兲. RBS was performed with 1.8 MeV He ions. The film composition was simulated using the RUMP program.30 ERDA measurements were carried out with 86Kr, 129Xe, or 197Au ions with energies of 1.8 MeV/amu.31 Scanning electron microscopy 共SEM兲 was performed with a Hitachi field emission microscope 共S-4100兲 using an electron energy of 25 keV. Cross-sectional transmission electron microscopy 共TEM兲 samples were analyzed in a Philips CM12 microscope at 120 kV with a LaB6 cathode. Sample preparation was done according to standard mechanical and ion beam polishing techniques. On high resolution TEM images fast Fourier transforms 共FFT兲 were performed. Grazing incidence x-ray reflectometry 共GIXR兲 was carried out with Cu K ␣ radiation by a Siemens D5000 diffractometer. The density, thickness, and surface roughness of the films were calculated with the GIXR simulation REFSIM1.1 program.32 Generally, the film thickness was measured with a profilometer 共DekTak 3030, Veeco兲 at a step formed by scratching.

III. RESULTS AND DISCUSSION A. EDXRD

The EDXRD spectra were continuously measured during depositions in time intervals of 20–30 s. Depending on the deposition conditions film thicknesses of 60–5000 nm were obtained. Each spectrum was analyzed by fitting Gaussians to the diffraction and fluorescence peaks. The time series of the peak intensities, positions, and widths were used to calculate the film thickness and the structural parameters 共i兲 coherently diffracting domain size, 共ii兲 lattice parameter, and 共iii兲 grain size, and 共iv兲 microstrain during deposition.26 Deposition 关the EDXRD spectra series of which as shown in Fig. 1共a兲兴 was performed at a relatively low deposition rate and displays only 共002兲 and combined 共100兲/共101兲 Bragg reflections of hexagonal 2H-MoS2 phase 共Ref. 33兲 and Mo K ␣ , ␤ fluorescence lines. It can be seen that 共001兲 texture at the beginning of the deposition changes into 共100兲/ 共101兲 texture at the end. The development of the peak areas of the 共002兲 and combined 共100兲/共101兲 reflections with increasing film thickness or deposition time is shown in Fig. 3共c兲. In order to quantify the texture, measured by EDXRD, a 共001兲-texture parameter c tex normalized to the JCPDS data was introduced: c tex⫽A 002• 共 I 002共 JCPDS兲 ⫹I 100共 JCPDS兲兲 / 关共 A 002⫹A 100兲 •I 002共 JCPDS兲兴 ,

共1兲

where A hkl is the peak area and I hkl(JCPDS) the reflection intensity of 2H-MoS2 according to the JCPDS data base.33 A 共001兲-textured MoSx film, which shows only 共002兲 and no 共100兲 reflex, has a texture parameter value of 1.22, while for a perfectly 共100兲-textured film c tex is equal to zero. In Fig. 1共b兲 the texture parameter c tex and the lattice strain ⑀ of the c axis 关 ⑀ ⫽(c film – c powder)/c powder兴 are depicted with respect to their dependence on the number of Mo atoms deposited for the deposition in Fig. 1共a兲 performed at 3 Pa and for one performed at 9 Pa, i.e., at a lower deposition rate. Due to texture crossover the texture parameter steeply decreases with an increase in deposition time, i.e., for a linearly increasing atomic areal density of Mo, while the lattice strain in the direction of the c axis is reduced from 3.5% to about 2%. The lattice strain versus the Mo atomic areal density shows a behavior typical of thin film growth: a steep increase from negative strain values 共i.e., compression of the c axis兲 to a positive maximum value 共i.e., expansion of the c axis兲 and subsequent relaxation. Maximum expansion of the c axis is significantly smaller for the lower deposition rate, i.e., for deposition at sputtering pressure of 9 Pa. Calculation of the grain size26 in the direction of the c axis from the peak widths of the 共002兲 reflections resulted in values between 7 and 17 nm, significantly smaller than those for the strongly 共001兲 textured films deposited at 9 Pa where grain size up to about 50 nm was observed.28 1. Thickness dependence of the (002) diffraction peak

The 共002兲 diffraction peak areas versus the film thickness are shown in Fig. 2共a兲 for depositions performed at different sputtering pressures. The signals measured at sputtering pressures higher than 3 Pa in particular show a char-



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FIG. 1. 共a兲 Series of EDXRD spectra during magnetron sputtering of a MoSx film on an oxidized 共100兲 silicon substrate. The 共001兲 texture which is observed at the beginning of deposition turns into 共100兲 texture at larger film thickness. The vertical bars with triangles indicate the position and intensity of the powder diffraction pattern of 2H-MoS2 共JCPDS File No. 37-1492兲. Escape signals 共esc兲 are due to the detection system and occur at energies 9.9 keV (Ge K ␣␤ fluorescence energy兲 smaller than their ‘‘mother peaks.’’ Deposition parameters: Total deposition time 19 min, sputtering power 50 Wdc, gas flow ratio F H2 S /(F Ar⫹F H2 S)⫽0.75, sputtering pressure 3 Pa, substrate temperature 450 °C, and substrate bias ⫺8 V 共floating potential兲. 共Diffraction angle ␪⫽5.036°, stoichiometry MoS1.47 , total film thickness 1.6 ␮m兲. 共b兲 Texture parameter c tex and 共c兲 lattice strain, depending on the atomic areal density of Mo for this deposition 共closed symbols兲 and a film deposited at lower deposition rate 共open symbols; sputtering pressure 9 Pa兲 for comparison.

FIG. 2. 共a兲 Development of the peak area of 共002兲 diffraction signals 共normalized to constant primary x-ray beam intensity兲 vs the atomic areal density of Mo as a measure of the film thickness during MoSx deposition at different sputtering pressures. Deposition parameters: Sputtering power 50 Wdc, gas flow ratio F H2 S /(F Ar⫹F H2 S)⫽0.75, substrate temperature 450 °C, and floating substrate potential. 共b兲, 共c兲 Normalized saturation values depending on the deposition parameters sputtering pressure 共䊊兲, substrate temperature 共䊏兲, deposition parameters: Sputtering power 50 Wdc, gas flow ratio F H2 S /(F Ar⫹F H2 S)⫽0.75, sputtering pressure 1 Pa, and floating substrate potential, and substrate bias voltage 共䉭兲, deposition parameters: sputtering power 50 Wdc, gas flow ratio F H2 S /(F Ar⫹F H2 S)⫽0.75, sputtering pressure 4 Pa, substrate temperature 450 °C, and floating potential ⫺18 V兴. In 共c兲 the texture parameter c tex is also given with regard to its dependence on the substrate bias voltage 共䊉兲.

acteristic steep increase at the beginning of film growth. Independent of the sputtering pressure saturation is observed, which means that growth of the 共001兲-oriented grains stops when a certain thickness is reached. This can be inferred also from the fact that a new diffraction signal from 共100兲 and 共101兲 lattice planes occurs when the 共002兲 signal reaches saturation 关see Figs. 3共b兲 and 3共c兲兴. The dependence of the saturation values on the sputtering pressure, the substrate temperature, and the substrate bias voltage is shown in Figs. 2共b兲 and 2共c兲. At high sputtering pressures and high substrate

temperatures, saturation occurs at higher peak areas compared to depositions performed at low pressures or low temperatures. This means that the thickness of that part of the film that has 共001兲 orientation, i.e., the crystal orientation required, increases rapidly when the sputtering pressure is increased to more than 3 Pa and when the substrate temperature is increased to more than 530 °C. The advantageous effect of the substrate temperature is obvious, since higher temperatures increase the adatom mobility and, hence, the thickness of the coherently diffracting film regions. In order


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FIG. 3. Dependence of the peak areas of the Mo K ␣ fluorescence 共a兲, 共002兲, and 共100兲/共101兲 Bragg signals 共b兲–共d兲 vs the film thickness and deposition time of three MoSx thin films deposited at different sputtering pressures. Diagrams 共c兲 and 共d兲 clearly show texture crossover from 共001兲 to 共100兲 orientation. The deposition parameters are the same as in Fig. 1.

to decide which effect is responsible in the case of variation of pressure, additional experiments were performed: we varied the power in the dc and the rf plasma excitation modes and the substrate bias. The latter is shown in Fig. 2共c兲. By varying the power, it was found that the 共002兲-diffraction line intensity and the texture parameter increase with a decrease in discharge power, i.e., with a decrease in deposition rate. This also means that in the case of pressure variation the reduced deposition rate is mainly responsible for the better 共001兲-growth of the MoSx films. With an increase in negative substrate bias voltage the saturation value of the 共002兲 signal is reduced. This decrease is connected with a decrease in texture parameter c tex in this series 关closed circles in Fig. 2共c兲兴, which means that high bias voltages do not only reduce the coherently scattering volume of the basal 共001兲-oriented film but at the same time increase the volume of 共100兲/共101兲 crystallites. This can be explained by the detrimental effect of bombardment of the growing film by particles with higher energies 共mainly argon reflected from the target兲. Energetic bombardment leads to the formation of defects in the film surface. The defects, in our case, form nuclei for the growth of 共100兲/共101兲-oriented crystallites. This influence of energetic particle impact on the morphology and phase formation was also observed during magnetron sputtering of boron nitride where the cubic phase grows on top of the hexagonal or turbostratic phase of BN. The latter phase also crystallizes in a van der Waals type structure like MoS2 . 34 Figure 2共a兲 also reveals that 共002兲 reflection could even be observed for very thin films with atomic areal densities of molybdenum less than 2⫻1016 atoms cm⫺2 , corresponding to a film thickness of about 20 S–Mo–S slabs. Therefore, nucleation of the MoS2 grains must occur with 共001兲 orientation. This was the case for all deposition conditions with respect to substrate temperature, pressure, and H2 S content of the atmosphere. 2. Texture crossover

Figure 3 shows the development of the Mo K ␣ fluorescence line and the 共002兲 as well as the 共100兲/共110兲 peak areas

in three experiments performed at different sputtering pressures. The increasing slopes of the Mo K ␣ intensities with a decrease in sputtering pressure in Fig. 3共a兲 reflect the rise in deposition rate, i.e., number of Mo atoms deposited per time. The linear development of the fluorescence intensities is a clear indication of a constant deposition rate during our experiments. In contrast, the slope of the 共100兲 area increase is smaller at 9 and 3 Pa than at 0.5 Pa 关Figs. 3共b兲–3共d兲兴. At the same time, when the 共002兲 diffraction signal starts to saturate, combined 共100兲/共101兲 reflection begins to rise. The transition from initial grain growth with 共001兲 planes parallel to the substrate surface to grain growth with 共001兲 planes perpendicular to the substrate is known as texture crossover.35 It has to be noted that this crossover does not coincide perfectly with a change of quantitative texture parameter (c tex), which is normalized to the powder diffraction intensities and has a value of c tex⫽1 at texture crossover. Besides, the energy dependence of the incident x-ray beam intensity and the Ge-detector efficiency were neglected in our calculations of the texture parameter. Similar texture crossover was observed by Schell et al.35 during sputtering of TiN thin films. Therefore, this effect seems to be a common feature of sputtered films. The reason for texture crossover for TiN was explained by those authors as competition between the surface and stress energy. Up to a critical film thickness surface energy is dominant, leading to 共001兲 texture, whereas, with increasing film thickness mechanical stress in TiN favors 共111兲 orientation. The situation for growth of layered semiconductors is quite different. Due to van der Waals bonding between sulfur–Mo–sulfur stacks no significant mechanical stress can develop in the films. This has been proven by measuring the bending of thin Si cantilevers before and after MoSx deposition. Therefore, mechanical stress cannot be the driving force for texture crossover of this material. On the other hand, Moser and Le´vy36 drew attention to the match of the atomic sulfur–sulfur distance d S–S 共along the a axis in the van der Waals plane兲 and the Mo–Mo distance 共also along the a axis兲 and explained texture crossover of MoS2 thin films on the basis of this coincidence. However, this reasoning does not explain why



FIG. 4. SEM cross-sectional micrographs of MoSx films deposited at different substrate potentials. Deposition parameters: 50 Wdc; F H2 S /(F Ar⫹F H2 S) ⫽0.75; p⫽4 Pa; T sub⫽450 °C. 共a兲 Floating potential 共⫺10 V兲, film thickness 450 nm; 共b兲 ⫺300 V, film thickness 50 nm. The angle of view is 30° relative to the substrate surface.

this texture crossover depends on the deposition rate. Their argument means that texture crossover is an inevitable result of the crystal structure in thin film growth of these van der Waals materials and should also occur in single-crystal growth. Since we observe a distinct dependence of the texture crossover on the deposition conditions this argument cannot explain the phenomenon. Furthermore, it is evident that dangling bonds primarily occur at lattice defects where non-van der Waals planes are present at the film surface. It is plausible that at these defect sites crystallites with their c axes parallel to the substrate 关 c 储 or 共100兲 orientation兴 start to grow, since the dangling bonds are directed out of the 共001兲 lattice planes, favoring 共100兲 growth of the crystallites 共see also Sec. III B兲. B. Scanning and transmission electron microscopy

In order to confirm the EDXRD results, selected MoSx films were analyzed by SEM and cross-sectional TEM. The SEM micrographs 共Fig. 4兲 reveal that an increase in the negative substrate bias voltage changes the morphology from a dendritic structure with high surface roughness 关Fig. 4共a兲兴 to a denser structure with a relatively smooth surface 关Fig. 4共b兲兴. The EDXRD spectra of these films show that this change in morphology is connected with a decrease in tex-

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ture parameter c tex , i.e., c⬜ growth is disturbed. This behavior can be explained by an increase in the number of defects, like point defects, edge dislocations, stacking faults, and grain boundaries by the high energy of the positive ion species which hit the growing film. Further SEM investigations confirmed the strong influence of the deposition rate and the substrate temperature on the structure and morphology which have been published recently.28 From Fig. 4共a兲 it can be also understood why, once 共100兲-oriented crystallites arise, the texture changes irreversibly from 共001兲 to 共100兲. A porous rough film with 共100兲 crystallites does not allow further growth of the basal 共001兲oriented film because it becomes quickly covered with c 储 -oriented grains that have high growth velocity in the direction of the substrate normal. These fast-growing c 储 -oriented grains lead to a shadowing effect,37 which explains the open, porous morphology of these films, which appear velvet like to the naked eye. Cross-sectional transmission electron microscopy 共XTEM兲 investigations were performed to get detailed information about the region where orientation crossover occurs. In high-resolution TEM micrographs 共Figs. 5 and 6兲 the S–Mo–S slabs are well resolved as black lines with distance of approximately 0.62 nm. The bright spaces represent van der Waals gaps.36 The micrographs show that all c 储 -oriented crystals grow on 共001兲-oriented basal layers with thickness of at least three to four S–Mo–S slabs, as has also been observed by others 共e.g., in Ref. 37兲. Grazing incidence x-ray reflectometry on 共001兲 textured films with film thickness ⭐100 nm confirmed that the basal layers had densities between 4.7 and 5.1 g cm⫺3, which is close to the MoS2 bulk density of 5.06 g cm⫺3 共Ref. 33兲共see Sec. III C兲. The TEM micrographs confirmed the reduced density of the 共100兲oriented upper part of the film in comparison to the basal layer density 关see Fig. 6共b兲兴. Edge dislocations where one S–Mo–S slab branches into two can be identified in Fig. 5共b兲. This type of defect has already been observed in TEM investigations of MoS2 and WS2 thin films and nanotubes.38 – 40 Whitby et al.40 stated that layer branching is caused by metal vacancies in the hexagonal layered lattice and that wave-like structures occur where ‘‘the mean value of separation across the layers is maintained.’’ They demonstrated that vacancies are also essential for the formation of the onion-like structures of WSx nanotubes. Displacement defects in the 共110兲 plane which lead to wave-like distortion of the layered material were described by Jose´-Yacama´n et al.41 The TEM cross section in Fig. 6共b兲 demonstrates that such wave-like structures can also be found in our films and result in ring segments in the 共002兲 feature of the two-dimensional FFT in the TEM micrograph. Nonetheless, lattice spacing 共inverse distance from the center of the FFT兲 is constant in the crystallites. Due to the limited resolution of the microscope available the 2H 共with two S–Mo–S stacks in a hexagonal unit cell兲 and 3R 共with three S–Mo–S stacks in a rhombohedral unit cell兲 polytypes of the MoS2 structure could not be distinguished and possible stacking faults were not identified. Other groups have been able to resolve triangular S–Mo–S groups in one slab and relative orientation of these layers of


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FIG. 5. XTEM micrographs of MoSx film on oxidized silicon showing c 储 -oriented grains arising from the basal c⬜ layer. 共a兲 Along the white line the thickness of the c 储 crystallite and its lattice spacing were measured as 25 and 0.637 nm, respectively. Inset: Fast Fourier transform of the XTEM micrograph. 共b兲 The basal c⬜ layer has a minimum thickness of 13 S–Mo–S slabs. The open circles encompass crystallographic defects. Deposition parameters: 50 Wdc; F H2 S /(F Ar⫹F H2 S)⫽0.75; p⫽4 Pa; T sub⫽450 °C; U sub⫽⫺18 V 共floating potential兲; stoichiometry MoS1.87 .

triangles in succeeding slabs.42– 44 It is probable that stacking faults also occur in our MoSx films because there is a strong correlation between edge dislocations and grain boundaries, on the one hand, and stacking faults, on the other hand. A polycrystalline material can be regarded as a network of edge dislocations. If a dislocation moves through a crystal towards the edges, the resulting stacking order can differ from regular stacking in polytypes 2H or 3R of MoS2 throughout the entire crystal. A high dislocation density can, therefore, lead to distorted order of the S–Mo–S-slabs in the direction of the c axis. This so-called turbostratic growth has been observed in different materials that crystallize in a layer-type structure: graphite, hexagonal boron nitride, and transition metal sulfides.34,45,46 The d 002 distance, i.e., that of the S–Mo–S slabs, in the XTEM micrographs was analyzed by two-dimensional fast

FIG. 6. High resolution XTEM micrographs showing details of MoSx film on oxidized silicon. 共a兲 The van der Waals planes parallel to the substrate (c⬜ ) bend continually until they are perpendicular to the substrate (c 储 ). 共b兲 The MoSx film lifted from the substrate. It shows wave-like growth of the van der Waals planes. Bottom: Fast Fourier transforms of the TEM micrograph in 共b兲 and, for comparison, an XTEM micrograph of a MoS2 single crystal 共micrograph not shown兲. Deposition parameters: 共a兲 The same as in Fig. 5. 共b兲 t tot⫽13 min; 50 Wdc; F H2 S /(F Ar⫹F H2 S)⫽0.75; p⫽4 Pa; T sub ⫽450 °C; U sub⫽⫺10 V 共floating potential兲; stoichiometry MoS1.92 .

Fourier transformation of selected areas, which yielded lattice parameters d 002⫽c/2 between 0.62 and 0.65 nm. This confirms lattice expansion in the c direction of several percent measured by the XRD experiments 共powder value d 002 ⫽0.6149 nm).



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FIG. 7. Density 共䊐兲 and surface roughness 共䉭兲 depending on the film thickness of reactively sputtered MoSx thin films. Deposition parameters: Sputtering power 100 Wrf 共13.56 MHz兲, gas flow ratio F H2 S /(F Ar⫹F H2 S) ⫽0.75, sputtering pressure 1.6 Pa, substrate temperature 450 °C, grounded substrate, and substrate–target distance 150 mm. The dashed line indicates the density of a MoS2 single crystal 共5.06 g cm⫺3兲.

On the basis of the high resolution TEM micrographs a dislocation density of about 3⫻1012 cm⫺2 was estimated by counting the edge dislocations in a Fourier filtered TEM spectrum of the c⬜ -oriented basal layer. For comparison a TEM cross-sectional sample of MoS2 single crystal was prepared whose FFT is shown in Fig. 6共b兲共bottom兲. Using ERDA stoichiometry of a MoS1.94 and about 3 at. % oxygen contamination were measured for this single crystal. C. Film density

In order to evaluate possible reasons for the observed lattice expansion, the density of the MoSx films was measured for its dependence on the film thickness by grazing incidence x-ray reflectometry.47 Figure 7 shows the result for four films which had a mirror-like appearance, i.e., the films were highly reflective and had thicknesses below the point of texture crossover. The film density decreases slightly with an increase in film thickness. At the same time the surface roughness of the films increases. ERDA measurements on thicker samples revealed densities down to about 1 g cm⫺3. These low densities are due to the very porous dendritic morphology of the 共100兲-oriented crystallites. D. Stoichiometry

It is well known that molybdenum sulfide films often exhibit stoichiometries the MoSx with x⭐2.7,48,49 Such deviation from stoichiometric sulfur-to-molybdenum ratio x ⫽2 leads to unacceptable electronic film quality with respect to photovoltaic application. Films that are sulfur deficient mostly exhibit low resistivity and almost no photoactivity. Therefore, to optimize the deposition parameters for MoS2 films, the chemical composition of the MoSx films was analyzed. The dependence of the sulfur-to-molybdenum ratio, determined by RBS and ERDA, is displayed in Figs. 8共a兲 and 8共b兲 as a function of the gas flow ratio F H2 S /(F Ar⫹F H2 S) for two sputtering powers 共100 and 200 Wdc) at two substrate temperatures 共410 and 150 °C兲.

FIG. 8. Stoichiometry 关S兴/关Mo兴 and concentration of contaminants H, C, N, O and Ar of reactively sputtered MoSx thin films on glassy carbon depending on the 共a兲 and 共b兲 H2 S/Ar gas flow ratio and 共c兲 substrate temperature. The stoichiometry was measured by ERDA 共⫻兲 and by RBS 共䊊兲; contamination data were obtained from ERDA. Deposition conditions: Sputtering power 共a兲, 共c兲 100 and 共b兲 200 Wdc; substrate temperature 共a兲 410 °C and 共b兲 150 °C; total sputtering pressure 共a兲, 共b兲 0.18 and 共c兲 0.3 Pa, and gas flow ratio F H2 S /(F Ar⫹F H2 S)⫽0.78 and floating substrate potential.

For low values of F H2 S /(F Ar⫹F H2 S) the S-to-Mo ratio increases and reaches a saturation of about 1.6 –1.7, independent of the substrate temperature and sputtering power, i.e., the deposition rate. The ratio F H2 S /(F Ar⫹F H2 S), where the saturation is reached, is shifted toward higher values for the 200 W series 关Fig. 8共b兲兴, which is due to the higher deposition rate 共about a factor of 2兲. This is typical for reactive magnetron sputtering of compounds from a metallic target 共see, for instance, Ref. 50兲.


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Down to a H2 S content of 9% in the Ar–H2 S sputtering gas mixtures no phase other than MoS2 could be observed in the EDXRD spectra. The coexistence of Mo metal or Mo2 S3 phase with molybdenite phase MoS2 , which is known from the phase diagram,51 could explain the sulfur deficiency, but such phases have not been identified by XRD in crystalline MoSx films deposited by sputtering methods. It is, of course, expected that metal Mo phase appears when the H2 S content is further decreased. Most of the samples of both series were investigated additionally by RBS, and the results agree with the ERDA results within measurement uncertainty. Of the contaminants found in the films, hydrogen and carbon are the most prominent with concentrations up to 5 at. %. The other contaminants, nitrogen, oxygen, and argon, have concentrations less than 1 at. %. The dependence of the hydrogen and carbon concentrations on the ratio F H2 S /(F Ar⫹F H2 S) in the series in Figs. 8共a兲 and 8共b兲 leads to the conclusion that, in both cases, contamination stems from the process gases, most probably from H2 S, since it contains relatively large amounts of carbonoxysulfide 共COS兲. The total amount of contaminants detected in the films does not account for the large deviations 共up to 20%兲 from the MoS2 composition. The range of sulfur-to-molybdenum values found in our experiments are consistent with results of other groups.1,48,49,52 There are indications from structural investigations 共TEM and XRD兲 that this sulfur deficiency is connected to structural imperfections of the MoS2 lattice. Because of weak binding between the van der Waals 共001兲-lattice planes, crystallographic order in the c direction is disturbed, leading to turbostratic growth. This is possibly accompanied by sulfur deficiency in these van der Waals planes. Further experiments using higher substrate temperatures and lower deposition rates are necessary to check this hypothesis. In Fig. 8共c兲 the S-to-Mo ratio, determined by ERDA, is shown as a function of the substrate temperature. The stoichiometric ratio S/Mo⫽2 is only reached at temperatures below 200 °C. For higher substrate temperatures the sulfurto-molybdenum ratio decreases to 1.7, which is consistent with values in the literature.49,52 This significant sulfur deficiency can only partly be explained by the contaminants 共H,C,O兲 found in the films. Since the H2 S partial pressure is sufficiently high, one would not expect decomposition due to desorption of sulfur from the growing film. Another factor that could cause sulfur deficiency is the relatively high amount of reactive, i.e., atomic, hydrogen that is formed in the plasma from the reactive H2 S sputtering gas, and which could lead to a reduction of the growing MoSx film. The hydrogen detected seems to not be adsorbed on the film surface, which is very high for dendritic-like growing films. This was concluded because no significant desorption of hydrogen by ion bombardment was measured during ERD analysis. Therefore, the hydrogen is expected to be fixed in the volume of the crystallites, which could be another reason for the observed c-lattice strain if hydrogen species are intercalated into the van der Waals planes. This effect has been found for MoS2 powder treated in hydrogen at pressures up to 50 bar. Stoichiometries of H0.5MoS2 共i.e., 14 at. % H兲 and c-lattice expansion of ⫹23% have been reported.53 MoS2

nanotubes were found to be capable of storing 1.2 wt % hydrogen,54 and it is suggested that the storage capability is strongly dependent on the specific surface area of MoS2 , 55 which is also an important factor in our films due to the dendritic morphology. On the other hand, self-intercalation of Mo atoms in the van der Waals planes of MoS2 may explain the sulfur deficiency and c-lattice expansion. It was proposed56 that intercalation of S, W, or O atoms into van der Waals gaps of WS2 may be responsible for the increase in lattice parameter c although no proof for this assumption has been be given so far. Using ERDA, it was found that an increase of negative substrate bias voltage up to 100 V did not change the stoichiometry significantly. The S-to-Mo ratio was 1.8 –1.9 for this, which is the range expected for the substrate temperature of 450 °C and a sputtering pressure of 4 Pa chosen. IV. CONCLUSIONS

The deposition of MoSx films by reactive magnetron sputtering in Ar/H2 S atmosphere from a metallic target was investigated in situ by energy dispersive x-ray diffraction. By analyzing the characteristic peak parameters of the fluorescence and diffraction peaks of the spectra series it was possible to obtain time-resolved data for the film thickness, lattice expansion, grain size, and the texture of MoSx films. The in situ diffraction analysis reveals that under the conditions of high substrate temperature and low deposition rate strong 共001兲 texture is observed when only a small number of crystallites with 共100兲/共101兲 orientation are present. For all deposition conditions the films start to grow with 共001兲 crystallite orientation, i.e., with the van der Waals planes parallel to the substrate surface. It was found that during growth of reactively sputtered MoSx films texture crossover from 共001兲 to 共100兲 texture occurs, especially at high deposition rates and for high energetic particle bombardment of the films. High resolution cross-sectional TEM images reveal a very high dislocation density. The dislocations may be nucleation sites for 共100兲-oriented crystallites and, therefore, responsible for texture crossover. The film composition analysis by the elastic recoil detection method revealed a significant sulfur deficiency in the films of up to 20%, which we ascribe to the reducing effect of atomic hydrogen in the sputtering atmosphere. ACKNOWLEDGMENTS

The authors thank K. Harbauer, J. Hinze, K. Martin, and P. Vo¨lz for their technical help and for assisting construction with that entailed various modifications of the sputtering chamber. H. Wulff and M. Quaas 共Ernst-Moritz-ArndtUniversita¨t, Greifswald, Germany兲 are acknowledged for the x-ray reflectivity and S. Lindner for the ERDA measurements. Thanks also go to Y. Tomm and A. Ramirez for providing the MoS2 single crystal. V. Buck, Wear 114, 263 共1987兲. T. Spalvins, J. Vac. Sci. Technol. A 5, 212 共1987兲. P. D. Fleischauer, Thin Solid Films 154, 309 共1987兲. 4 H. Wise, Polyhedron 5, 145 共1986兲. 5 H. Tributsch, Ber. Bunsenges. Phys. Chem. 81, 361 共1977兲. 1 2 3


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