Growth and Characterization of 3C-SiC Films for Micro Electro

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DOI: 10.1021/cg900677c

Growth and Characterization of 3C-SiC Films for Micro Electro Mechanical Systems (MEMS) Applications

2009, Vol. 9 4852–4859

Matteo Bosi,*,† Bernard E. Watts,† Giovanni Attolini,† Claudio Ferrari,† Cesare Frigeri,† Giancarlo Salviati,† Antonella Poggi,‡ Fulvio Mancarella,‡ Alberto Roncaglia,‡ Oscar Martı´ nez,# and Vanesa Hortelano# †

IMEM-CNR Institute, Parco Area delle Scienze 37/A, 43100 Parma, Italy, ‡IMM-CNR Institute, Area della Ricerca di Bologna, Via P. Gobetti 101, 40129 Bologna, Italy, and #Optronlab group, Departamento Fı´sica de la Materia Condensada, Edificio IþD, Universidad de Valladolid, Paseo de Bel en 1, 47011-Valladolid, Spain

Received June 18, 2009; Revised Manuscript Received September 1, 2009

ABSTRACT: The growth of 3C-SiC on (001) silicon substrates by means of vapor phase epitaxy is described. The growth mechanisms are discussed with the aid of structural and morphological characterizations performed by X-ray diffraction, transmission electron microscopy, and atomic force microscopy. Raman spectroscopy was used to study the residual stress. A large shift of Raman peaks with respect to the expected values for the bulk is observed and explained by the relaxation of Raman selection rules due to lattice defects. The stress and stress gradients through the film thickness are observed and studied on micrometer-sized structures such as membranes and cantilevers. Local Raman peak fluctuations are observed on millimetersized membranes, while cantilevers show different degrees of curling depending on film thickness.

1. Introduction Silicon carbide (SiC) has gained importance as both a coating and a structural material for micro electro mechanical systems (MEMS).1,2 SiC is a wide bandgap semiconductor used for high temperature, high power applications and radiation-hard environments. The high Si-C bond energy confers a high Young’s modulus and hence mechanical toughness and a high fracture strength;3 moreover, it is chemically inert to the most corrosive and erosive chemicals and is biocompatible.4-6 More than 100 polytypes of SiC exist but the SiC cubic phase (3C-SiC) has drawn particular attention because it can be deposited on Si. Electronic devices such as metal-oxide-semiconductor field-effect transistors (MOSFETs) or diodes can be fabricated from 3C-SiC/Si thanks to epitaxial techniques and dopant implantation, so that integration between MEMS and 3C-SiC electronics is possible, enabling the fabrication of high temperatures, harsh environments, and radiationresistant devices.7,8 3C-SiC is also used as a high temperature gas sensor in the form of a Schottky diode with Pd contacts that catalyzes H2 or other gases.9-11 3C-SiC can be grown on Si substrates despite the 20% lattice mismatch and 8% thermal mismatch. This allows low cost and large area growth on Si substrates and the use of conventional etching and photolithographic processes to micromachine structures. Moreover, 3C-SiC acts as a natural etch-stop during wet etching of Si, since it is not attacked by any chemicals at low temperatures. Unfortunately, the high lattice and thermal mismatch and chemical interdiffusion12 hinder the growth of SiC on Si by creating high residual stress, and generating a highly defective layer at the interface which must be appropriately controlled for specific applications. Large tensile or compressive residual stress lead to macroscopic wafer bending,13 while variation of *Corresponding author. E-mail: [email protected]. pubs.acs.org/crystal

Published on Web 10/06/2009

stress through the film thickness may deform microstructures.14 In addition, stress can alter their resonant frequency and, consequently, their theoretically predicted behavior. MEMS are a family of technologies and integrated devices that are becoming more and more important in modern life. Some areas in which these systems are already applied are shock sensors for airbag, inkjet printers, accelerometers and gyroscopes for boats and airplanes, entertainment, healthcare instruments, communication and information technologies, biology and biosensors. MEMS offer significant advantages over hybrid systems and devices because of their small dimensions, integration of different components and low power consumption.6,15-17 Although Si is presently the most used material for MEMS fabrication, it has serious limitations for some applications, such as high temperature (T > 300 C) and/or harsh environments with corrosive chemicals and biocompatibility. Therefore, alternative materials are required for MEMS applications.7,8 In this work, we will discuss the growth and characterization of 3C-SiC/Si films by relating Raman spectroscopy to other techniques such as transmission electron microscopy (TEM) and X-ray diffraction (XRD), in order to study the presence of stress and defects of the material. Membranes and cantilevers, which are the basic building blocks for advanced MEMS devices, were fabricated and analyzed in order to study stress-related problems at the micrometer scale and stress gradient through the film thickness. 2 -. Experimental Section The SiC films were deposited on 2” p-type (001) Si substrates in a cold-wall VPE reactor, heated by induction. The substrates were etched in 10% HF solution for 60 s before introduction in the growth chamber and heated at 1000 C for 10 min to completely remove the Si native oxide layer. A H2 flow of 2000 sccm was used as the carrier gas, and the growth was performed at atmospheric pressure using silane and propane precursors (diluted 3% in H2). r 2009 American Chemical Society

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In order to relax the lattice mismatch between SiC and Si, a carburization step was performed before the growth: the temperature was lowered to 400 C, 3 sccm of propane was injected in the growth chamber, and the temperature was increased to 1100 C in about 2 min. This temperature was then held for 5 min to complete the carburization process and for a thin monocrystalline SiC layer to form on the Si substrate. Misfit dislocations are created at the interface, with an extra SiC plane every 4 Si planes and a 5:4 ratio between SiC and Si lattice constant.12 After the carburization process, the propane flow is interrupted, the temperature is raised to 1200 C and SiH4 and C3H8 are injected in the growth chamber with a C/Si ratio of 1.75. The SiC films were deposited with a growth time of 10 min and with the same C/Si ratio but using different C3H8 and SiH4 gas flows (SiH4 = 5, 2.5, and 0.83 sccm for samples A, B and C, respectively), all the other conditions being the same. Atomic force microscopy (AFM) analyses were performed in contact mode by using a Digital Instruments Nanoscope IIIa. XRD spectra were obtained on a Siemens D500 diffractometer with Cu KR radiation and Θ-2Θ geometry while the SiC film mosaic spread was measured by setting the 2Θ position to the SiC 002 peak and scanning the angle of the incidence Omega near the maximum position. TEM images were acquired using a JEOL 2200FS microscope operated at 200 kV. Æ110æ cross section specimens for TEM were prepared by sandwiching a sample between two slabs of Si. The sandwich was then mechanically ground down to 30-40 μm and subsequently thinned to electron transparency by Ar ion beam bombardment. Micro-Raman (μR) measurements were carried out at room temperature in a Labram UV-HR800 spectrometer (Jobin Ivon) attached to a microscope, exciting with an Arþ laser (514.9 nm) in a nearly backscattering configuration. A liquid nitrogen cooled CCD detector was used to detect the signal. The thickness of the SiC layers was measured from the interference fringes of the reflection spectra using a Jasco UV-vis V-530, knowing the SiC refractive index of 2.55 þ 3.41  10-4  λ-2, where λ is the wavelength of interest.18 Suspended SiC structures were fabricated by micromachining of the 3C-SiC film deposited on Si wafers: rear side bulk anisotropic Si etching was employed to fabricate SiC membranes ranging in size from 0.5  0.5 up to 1.3  1.3 mm2, and a front side micromachining process was used to obtain cantilevers and bridges, up to 300 μm long. The chemical robustness of the SiC films was tested during the membrane fabrication process because the release of the membrane was performed by Si etching in tetramethyl ammonium hydroxide (TMAH) at 90 C for 5 h with SiC film exposed to the etchant: no microcracks or holes were observed, indicating the films are of good quality for wet anisotropic TMAH bulk micromachining. Scanning electron microscopy (SEM) micrographs of the released microstructures were performed in a Zeiss Gemini 1530.

3. Results and Discussion 3.1. SiC Growth. The high lattice and thermal mismatch between SiC and Si can lead to significant wafer bending of the 3C-SiC epilayers when the growth is performed above 1300 C.19,20 To minimize this problem, the films were grown at a lower growth temperature. In common with experiments performed in horizontal reactors without a rotating susceptor,21 a thickness variation across the wafer is observed, due to the different precursors decomposition rates over the susceptor length. All the samples showed an almost linear thickness gradient along the flow direction: the thickness was measured at the wafer center and at about 5 mm from the inlet and outlet wafer borders: at the gas outlet position it was about 15% the value at the gas inlet position for all the samples. The depletion of silane along the length of a horizontal reactor chamber is a common problem and, since it is silicon (not carbon) that controls the growth rate,22 the SiH4 partial pressure in the flow direction is a very important parameter.

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Figure 1. SiC film growth rate at different positions along the wafer. The open symbols represent the thickness of sample C multiplied by 3 and 6 times. The lines are guides for the eyes.

Figure 1 shows the thickness measured at the three wafer positions given above, for samples grown with different SiH4 flow rates but identical C/Si ratios and total H2 flow rates. The SiH4 flow rate used for samples A and B are respectively 6 and 3 times the flow rate used for sample C. Given that the SiC growth rate is linearly dependent on the local SiH4 concentration,22 one would expect for samples A and B a thickness respectively 6 and 3 times the one of sample C. This is indeed not observed experimentally, since 6 times the thickness of sample C is greater than the thickness of sample A, as seen in Figure 1 (Δ symbols). This is particularly evident in the gas inlet zone also for sample B: the values measured on this sample are not obtained by multiplying 3 times the C thickness (see r symbols in Figure 1). This means that linearity between the SiH4 flow rate and the layer thickness is not observed for all the SiH4 partial pressures. These results suggest that when the SiH4 partial pressure is above a certain threshold in given flow conditions, most of the Si is not incorporated in the film, but reacts in the gas phase or is depleted at the reactor inlet before reaching the substrate. This also means that SiC growth rate cannot be increased continuously by raising the SiH4 concentration above a certain threshold. Maintaining the SiH4 flow below this threshold may help in reducing homogeneous and parasitic reactions. In order to overcome the thickness gradient caused by the SiH4 depletion along the wafer length, the flow dynamics in the growth chamber should be accurately addressed, for example, by modeling an optimal reactor geometry or by adjusting the amount of carrier gas flow. 3.2. Structural Characterization. XRD diffraction measurements show that the fwhm of the SiC (002) peak of sample C is slightly lower than that of the other two samples (about 0.50 compared to 0.60), suggesting that lower growth rates may lead to a better crystal quality. High resolution XRD performed on sample C (Figure 2a) shows the 3C-SiC (002) and (004) peaks at 41.5 and 90.1. No significant differences were found between the different positions along the flow direction. The mosaic spread of the SiC film was measured by setting the 2θ position to the SiC (002) peak and scanning the angle of incidence ω near the maximum position (Figure 2b). The measured value of about 1.4 is an indication of the good crystalline quality and corresponds to the average misorientation of the mosaic

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Figure 2. (a) High resolution-XRD of sample C at inlet, center, and outlet position. (b) Mosaic spread observed in the same film.

Figure 3. SiC layer C. (a) Typical TEM cross sectional image taken in the central area. (b) SAD pattern from the same sample. (c) High resolution TEM image of the planar defects.

domains, due to lattice defects, with respect to the (001) direction. TEM images of this sample (Figure 3a) show a SiC/Si interface free of micropipes or voids in the Si substrate, commonly seen in these heterostructures.23,24 The spots of the selected area diffraction (SAD) pattern (Figure 3b) show that the layer is monocrystalline and (001) oriented. Extra spots displaced from the matrix spots by vectors of ( 1/3 Æ111æ are visible and of appreciable intensity indicating the presence of a high density of twins. The high

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density of twins is confirmed by high resolution TEM images as the one of Figure 3c. Stacking faults are also present. Twins and stacking faults are typical defects of the epitaxial SiC system because their formation energy in SiC is extremely low.25 The high density of twins and stacking faults may even induce a small degree of polytypism, since the difference between the SiC polytypes is merely a change in the (111) plane stacking sequence. Although a contribution of a SiC hexagonal phase to the formation of the extra SAD spots cannot be excluded, its presence would not seem likely because of the low growth temperature used here. AFM images evidence a smoother film as the thickness diminishes. Figure 4 shows the evolution of the surface morphology along the flow direction in the three positions: smoother film grows downstream on the wafer, from the inlet toward the outlet position, and the root mean square (rms) roughness decreases from about 25 to 2 nm. Radmilovic26 and Fu27 investigated the dependence of 3C-SiC rms surface roughness and found that it increases with increasing thickness and, for a given thickness, it falls with increasing SiH4 concentration. Since a thickness gradient due to SiH4 depletion along the flow direction was observed, the discrimination between the two effects is not straightforward. A nearly linear dependence is observed between surface rms roughness and sample thickness (Figure 4d), suggesting a relation between these two quantities. However, more investigations are needed to discern between roughness induced by thickness increase and the one induced by the change in SiH4 partial pressure. 3.3. Raman Spectroscopy. Figure 5 shows the microRaman spectra taken on sample C at different points along the flow direction. The spectra were fitted with mixed Gaussian/Lorentzian curves to obtain the peak positions: the main 3C-SiC TO and LO lines, for this particular sample, are located at about 794.5 ( 1 cm-1 and between 965 and 970 cm-1, respectively, depending on the position. The other features observed at about 1515 cm-1 and 1710 cm-1 are ascribed to second order optical phonon modes.28 Both TO and LO peaks are sensibly broadened toward lower wavenumbers, and a broadband appears at about 900 cm-1 between the TO and LO modes. In all the samples analyzed there is no evidence of graphite bands that should be located at 1360 and 1620 cm-1.29,30 Figure 6 shows the results of the Raman peak fitting of both TO and LO modes for samples A, B, and C, together with the reference bulk peak position reported in the literature.31 The three samples show a different behavior: the position of the LO peak for samples A and B and of the TO peak for sample C is almost constant in all the wafer positions, within a range of about ( 1 cm-1. Instead, the TO peak position for samples A and B shows a similar decreasing trend, while the LO peak position for sample C shows an increasing trend. At the gas inlet position, the LO peak position of samples A and B is almost the same as reported in the literature for the bulk 3C-SiC. Raman peaks can shift because of temperature, external applied pressure, stress, stacking disorder, and coupling with electron plasmon modes.31-37 In order to investigate the cause in our samples, we have explored all these effects. Raman measurements carried out by varying the laser power showed no significant variations in the spectra; thus, the influence of heat on the peak shifts can be neglected.

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Figure 4. AFM images of three different zones (a: inlet, b: center, c: outlet) of sample C. Surface root-mean-square roughness is 25.5, 8.1, and 1.7 nm for a, b, and c, respectively. Panel d shows the relation between sample thickness and surface rms roughness for this sample.

Figure 5. Raman spectra of sample C taken in different positions along the wafer from the inlet to outlet zone.

Coupled phonon-plasmon modes are observed in hexagonal SiC polytypes: the collective free carrier excitation can interact with the LO phonons through a macroscopic electric field in order to form a coupled mode, but this is known to

Figure 6. Results of Raman peak fitting for the different samples in different zones, going from the inlet to the outlet part of the wafer. The lines are a guide for the eye. The open symbols represent the position that should have the LO peak if the shift contribution to this peak were due to stress only.

shift the LO peak position toward higher wavenumbers as the doping levels increase. Here, the shift toward lower wavenumbers excludes this effect.

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Several works describe quantitatively the biaxial stress effects due to lattice mismatch34,37 reporting a red shift with increasing biaxial compressive stress. Olego et al.32 described the TO and LO shift in dependence of the lattice mismatch Δa/a due to stress and found that the peak position (in cm-1) can be described by a linear relation: ωTO ¼ 796:5 -ð3734 ( 30ÞΔa=a ωLO ¼ 973 -ð4532 ( 30ÞΔa=a 34

Zhu et al. and Rohmfeld et al.35 successfully applied this model to describe their experimental data and, using appropriate Δa/a values, were able to explain both the ωTO and ωLO shifts. In particular, a redshift of ωTO and ωLO peaks is evidence of a tensile stress. Here, two values for Δa/a differing between 50% and 100% among themselves were obtained when applying the previous formulas, considering the values of ωTO and ωLO separately. This discrepancy suggests that there must be causes for the Raman peak shift other than biaxial stress. Moreover, such effects should also explain the different behavior of the LO and TO peak positions for the samples grown with different SiH4 flow rates. In the classical theory, the only phonon modes that give a Raman signal in a perfect crystal are the ones at the Brillouin zone center at k = 0. If the crystal is not perfect, due to the presence of nanometer-sized grains or crystal defects which interrupt the lattice periodicity, the Raman selection rules are relaxed and phonons far from the Brillouin zone (with k 6¼ 0) become Raman active, producing a broadening and/or a shift of the standard modes. Such effects have been observed in 3C-SiC nanowires,36,38 where the dimension of the crystal becomes important compared to the phonon mean free path, and in 3C-SiC with lattice defects such as stacking faults.28,33 As discussed previously, a high density of twins and stacking faults is evidenced by the TEM images, and a small degree of polytypism may be present.25 The peak broadening and the occurrence of the band between the TO and LO peaks could then be ascribed to relaxation of the Raman selection rules due to lattice disorder. Ward et al.38 analyzed in detail the effect of phonon confinement due to small-sized crystallites in SiC monofilaments, concluding that it could lead to a downshift and broadening of the Raman line shape. The peak shift and broadening effect depends on the phonon dispersion curve, and has been observed in other materials such as TiO2,39 GaN,40 GaAs,41 and Si.42 For SiC, the TO and LO phonon dispersion curves have different shapes, and in particular the TO curve is flatter (much less dependent on the wavevector) than the LO one:43,44 it is thus reasonable to expect a stronger redshift and broadening of the LO peak with respect to the TO one, as observed in our samples. Phonon confinement is generally observed in the presence of nanometer-sized grains or nanowires, but a high density of defects such as twins and stacking faults have a similar effect on the Raman spectra on SiC films. Since the TO and LO phonon dispersion curves are different, and because the TO one is flatter, it is reasonable to consider the TO peak shift less dependent on the stacking disorder and selection rules relaxation and to consider, as a first approximation, that the TO band shift is due to stress only, according to the model proposed by Olego and Cardona.32 In this framework, from Figure 6 one can infer

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that for sample C the average stress is nearly constant in all the wafer positions (the peak fitting position within an error of (1 cm-1), while for samples A and B the biaxial stress increases with decreasing thickness. Assuming that the TO position depends only on stress, one can calculate the biaxial stress, Δa/a, from the TO fitting and use it to obtain the value of ωLO for that stress. The open symbols in Figure 6 report these calculations, that is, they represent the LO peak position if the band shifted only because of stress. The observed difference between this theoretical value and the experimental values is then ascribed to the presence of phonon confinement and relaxation of the Raman k = 0 selection rule. As previously discussed, the LO peak positions derived from the fitting is almost constant in samples A and B. Given the fact that the TO line position shifts with increasing thickness, we suggest that for thicker samples (inlet position in Figure 6) the phonons selection rules become more and more relaxed (the influence of lattice defects increases) since the difference between the theoretical LO position derived from the Δa/a calculation and the experimental one increases. The influence of lattice defects (stacking faults, twins) on the Raman selection rules depends on the thickness, and in particular for increasing sample thickness the breaking of Raman selection rules become more important, a mechanism that could be related to stress relaxation induced by crystal defects. 3.4. Stress Analysis. Raman spectroscopy on the different samples, up to 3 μm thick, reveals a gradual attenuation of stress as the thickness increases. In particular, the overall stress determined by Raman measurements is tensile for the thinnest films, becoming less tensile with thickness. This suggests that tensile stress is relieved as the thickness increases, as several works point out.21,35 Zielinsky et al.19 observed that a change in the C/Si ratio in the gas phase could affect the strain in the samples: for high C/Si ratios the stress is tensile while for low C/Si ratios it is compressive. As described above the thickness variation in these samples could be due to silane depletion in the gas phase along the flow direction, so that even this effect may play a role in samples when the flow dynamics conditions are not optimal. In our samples, we observe a variation in stress along the flow direction which, according to the literature, may be ascribed to the different thickness and to a change of the C/Si ratio in the gas phase. Lattice and thermal mismatch in SiC/Si heteroepitaxy generate compressive and tensile stress, respectively. The stress due to lattice mismatch depends on the thickness, while the thermal contribution depends only on ΔT.45,46 However, lattice and thermal mismatch generate a uniform stress component, while more localized effects such as atomic diffusion, grain size variation through the film thickness, interstitial or substitutional defects may cause stress gradients (i.e., variation of the stress through the thickness).45 Moreover, interdiffusion between the substrate and the growing film or during carburization may generate stress at the local scale.12 Stress gradients in micromachining are known to produce effects such as cantilever curling, which are deleterious for MEMS applications.45,47 In SiC films, stress gradients were observed by Roper et al. and were found to depend also on postgrowth annealing conditions and possibly on oxygen contamination:48 in particular, their samples were more compressively stressed near the surface. Suspended microstructures were fabricated to study the presence of stress gradients and local variation of stress at the

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Figure 7. Images of two SiC freestanding membranes, buckled (a) and flat (b), both obtained on sample C.

Figure 9. Correlation between the TO and LO Raman spectra, obtained from peak fitting, for the freestanding membrane. Figure 8. Raman spectra obtained on the freestanding membrane.

micrometer scale. These structures were also used to assess the potential of our epitaxial SiC films for MEMS applications: excessive tensile stress may hinder the realization of freestanding structures because of the collapse of membranes and breaking of bridges. On the other hand, excessive compressive stress may induce buckling in the membranes and bridges. Furthermore, a poor quality of the 3C-SiC deposited layers may result in the impossibility to obtain freestanding microstructures. Figure 7 shows micrographs of two membranes obtained from sample C fabricated by etching the Si substrate to leave a square window of about 1.3 mm. Membranes were obtained in different wafer positions; thus, different amounts of stress could be related to the different thickness. In particular, we generally obtained “buckled” membranes (Figure 7a) when the thickness was over 1 μm and “flat” membranes (Figure 7b) with thickness of about 600-700 nm. The buckled membrane is an indication of excessive compressive stress. Raman analysis on the freestanding “flat” membrane is shown in Figure 8. The spectrum is very similar to the one taken from the SiC/Si film, the TO and LO positions being located at 795 ( 1 cm-1 and 965.4 ( 1 cm-1, respectively, depending on the position on the membrane. Figure 9 reports the results of the peak fittings of both TO and LO bands for different points along the membrane. A good correlation between the shifts of the TO and LO peaks is observed, showing that there is a small contribution of the

stress on the peak shift at the different points of the suspended structure. As was observed on the film, even in this case the LO peak has a stronger redshift than the TO one, ascribed to the relaxation of the Raman selection rules due to lattice defects. The shift toward lower frequencies indicates an overall tensile stress, coherent with the fact that the membrane is “flat” and not buckled. Moreover, the observed Raman peaks fluctuation suggests that local stress variations occur, which may be tentatively related to the presence of stress gradients at the micrometer scale. Raman analyses on buckled membranes were not performed since the deviation from the backscattering geometry is difficult to analyze. MEMS structures such as cantilevers were fabricated on sample C in order to investigate the dependence of the stress gradient in the SiC film on its thickness. The bending observed on these microstructures indicates a stress gradient in the film:49 Figure 10 presents the SEM micrographs of SiC cantilevers with a thickness of about 400, 800, and 1000 nm, showing that the deformation magnitude decreases with increasing thickness. In particular, in thicker layers the stress gradient attenuates and becomes negligible. The origin and the development of this stress gradient are not known at present and the literature on SiC is lacking about this topic. The combined results of Raman and micromachining on our samples suggest a complex behavior of the stress: the thickest membranes are buckled, indicating the presence of a compressive stress, while the Raman peaks shift toward higher wavenumbers as thickness increases, suggesting that the tensile average stress is gradually relieved, although

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were subsequently processed to fabricate micromechanical structures to assess the strain in the SiC at the micrometer scale. The samples were characterized by XRD, AFM, and TEM, to show that the films are monocrystalline with a high density of defects such as twins and stacking faults. An analysis of the growth rate indicates depletion of precursors inside the reactor. Moreover, it is suggested that SiH4 partial pressure should be kept below a certain threshold to avoid gas-phase reactions and excessive depletion of the nutrient phase. Raman analysis revealed a strong discrepancy between the standard TO and LO peak positions expected from the bulk SiC. The deviation was related to relaxation of Raman selection rules induced by defects, coupled to the shift induced by an excessive residual average stress incorporated in the layer. MEMS structures were realized from the epitaxial film and were used to assess the strain in the SiC at the micrometer scale. Raman spectroscopy was used to study flat membranes, and tensile strain with local fluctuations was observed. The deformation of the cantilevers indicated the presence of a strain gradient, which diminished as the film became thicker. An overall complex behavior and development of the stress incorporated in the film is evidenced. Acknowledgment. We acknowledge the helpful discussions with Francesca Rossi and Filippo Fabbri and the work of Carlo Mora and Tullo Besagni on XRD measurements. AFM images were obtained at the “Centro Interfacolt a Misure” of Parma University.

References

Figure 10. SEM images of cantilevers obtained from sample C with different thicknesses (400, 800, 1000 nm for a, b, c, respectively), showing different degrees of bending. The marker is 20 μm.

a compressive average stress is not observed by Raman even in the thickest films. The presence of a stress gradient through the film thickness is evidenced by the curling of the cantilevers (Figure 10a,b), and local stress fluctuations are observed by micro-Raman (Figure 9). The stress gradient is relieved as the thickness increases and becomes negligible for thickest films, as evidenced by the cantilever negligible bending (Figure 10c). As a tentative hypothesis, the high density of crystal defects, observed both by TEM and Raman spectroscopy, may contribute to the complex stress evolution with thickness. Flow dynamic conditions, local precursor partial pressure, and the C/Si ratio may also influence the stress incorporated in the film. Further analysis of both cantilevers and bridges are in progress to evaluate the magnitude of this stress gradient and the type of the average residual stress. Conclusions 3C-SiC films have been deposited on Si at atmospheric pressure with VPE technique, using a carburization procedure followed by growth using SiH4 and C3H8 at 1200 C. The films

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