Fundamental properties of a-SiNx:H thin films deposited by ICP-PECVD for MEMS applications

Applied Surface Science 284 (2013) 348–353

Contents lists available at ScienceDirect

Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc

Fundamental properties of a-SiNx :H thin films deposited by ICP-PECVD for MEMS applications D. Dergez ∗ , J. Schalko, A. Bittner, U. Schmid Institute of Sensor and Actuator Systems, Vienna University of Technology, Floragasse 7, A-1040 Vienna, Austria

a r t i c l e

i n f o

Article history: Received 28 May 2013 Received in revised form 15 July 2013 Accepted 20 July 2013 Available online 27 July 2013 Keywords: ICP-CVD Silicon nitride Deposition rate Residual stress Wet etching Infrared spectroscopy X-ray photoelectron spectroscopy

a b s t r a c t In this study, the impact of deposition conditions on the properties of amorphous hydrogenated silicon nitride (a-SiNx :H) films using an inductively coupled plasma enhanced chemical vapor deposition technique (ICP-CVD) is evaluated. Due to the large number of experiments – even when taking only the most important synthesization parameters into account such as the total pressure in the deposition chamber, the substrate temperature, the ICP power and the flow rate ratio of N2 /SiH4 – a design of experimentsbased approach is chosen. As expected, the deposition rate strongly depends on the ICP power and the N2 /SiH4 flow rate ratio, respectively. Films in the field of investigation deposited with a high flow rate of N2 labeled as Type I show relatively low mechanical stress values between −50 and +200 MPa, but exhibit a strong drift behavior toward compressive stress. Layers deposited at low nitrogen flow rates (Type II), however, yield large compressive stress and are stable as a function of time. The wet etch rate in hydrofluoric acid shows a gap of over two orders of magnitude when comparing the two a-SiNx :H types, indicating strong differences in the chemical composition. Fourier-transform infrared measurements demonstrate that in Type I films the hydrogen is mainly bonded to nitrogen, in contrast to Type II films, where Si–H bonds dominate. Surface related X-ray photoelectron spectroscopy measurements show that Type II layers have higher relative silicon content, while depth profiles yield that the oxygen content Type I films is above 10 at.%. This high oxygen content is proposed to be the result of diffusion of H2 O into the layer, causing oxidation, and, as a consequence, the drifting behavior of the intrinsic film stress. © 2013 Elsevier B.V. All rights reserved.

1. Introduction Silicon nitride (Si3 N4 ) thin films are besides silicon dioxide well established in the fabrication of silicon microelectronic and MEMS (microelectromechanical systems)-based devices due to their excellent electrical and mechanical properties in combination with a high chemical resistance and hence, have been subject to intensive research activities in the last five decades. Applications of silicon nitride and amorphous hydrogenated silicon nitride (a-SiNx :H) thin films in microelectronics and MEMS include, but are not limited, to electrical isolators [1], surface passivation layers [2], thin membranes and cantilevers [3], respectively. Especially the latter application scenario is nowadays in the focus of many research activities, as silicon nitride is regarded to be a key material for advanced atomic force microscopy (AFM) purposes. This user-case of cantilever production for AFM [3] sets a large number and partially competing requirements on the

∗ Corresponding author. Tel.: +43 15880176642; fax: +43 15880136698. E-mail address: [email protected] (D. Dergez). 0169-4332/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.apsusc.2013.07.104

mechanical (i.e. residual stress [4–6] and Young’s modulus [3]), chemical (i.e. etch selectivity) [7–9] and electrical [10] properties of these films. For high speed measurements in liquid environment for bio-medical applications small, but soft AFM cantilevers are targeted having a thickness of 300 nm or even below [11]. For sensing the deflection of the cantilever the optical approach is less advantageous. Therefore, strain sensitive elements such as nanogranular metals [12] need to be applied on the cantilever resulting with the necessary passivation layers in a multilayer system that needs to be in addition stress-compensated and provides a specific spring constant and quality factor. For the above applications vapor phase deposition techniques are commonly used to synthesize silicon nitride thin films. Physical vapor deposition (PVD) comprises radio-frequency sputtering of Si3 N4 [13] or reactive sputtering of Si3 N4 or Si targets in N2 atmosphere [14,15]. More common, however, is the application of chemical vapor deposition (CVD) techniques: low pressure chemical vapor deposition (LPCVD) [3,16] is a high temperature (typically: 650–900 ◦ C [17]) technique using silane (SiH4 ) or dichlorosilane (SiCl2 H2 ) and ammonium (NH3 ) as precursors. The integration of such high temperature processes into the process flow is often problematic; therefore alternative processes are

D. Dergez et al. / Applied Surface Science 284 (2013) 348–353

required with a lower deposition temperature. A possibility is to use catalytic CVD, where the sources gases are decomposed thermally using a heated filament [18], while the substrate temperature can remain at moderate levels. As an alternative, to compensate for the lower deposition temperature, a plasma can be ignited as an auxiliary energy source resulting in plasma-assisted or plasmaenhanced chemical vapor deposition (PECVD) [19–23]. In this case the SiH4 is often diluted in Ar/He or H2 for safety reasons. A variation of this process is the pulsed-PECVD technique [24,25]. A major drawback of conventional (capacitively coupled) PECVD techniques is the relatively high concentration of hydrogen in the synthesized films leading to a low density. This problem can be decreased by replacing NH3 by N2 or by specific post-deposition temperature-treatments, such as annealing in different gas atmospheres [26–28], under vacuum conditions or by rapid thermal annealing [29]. Although this seems to substantially lower the advantage of performing the deposition at moderate temperatures, it may still be advantageous in some cases considering the higher flexibility of the process. Other varieties of PECVD may also be used, which utilize high density plasmas (HDP-CVD) [30], as electron cyclotron resonance (ECR-CVD) [7] or inductively coupled plasmas (ICP-CVD) [2,31–33]. These deposition techniques allow a high power density, which leads to a higher degree of dissociation for the precursor gases, resulting in lower hydrogen content of the films, at moderate deposition temperatures, for e.g. solar cell surface passivation purposes [2,23,30]. The ICP-CVD process offers the further advantage of allowing the fine-tuning of intrinsic film stress by the high flexibility of the process. This study aims to clear some basic relationships between the deposition parameters and resulting properties of a-SiNx :H films synthesized with the ICP-CVD deposition technique.

2. Experimental details 2.1. Silicon nitride deposition For all experiments, double-side polished, n-type (phosphorous doped,  > 50  cm) (1 0 0)-oriented single-crystalline 4 in.-silicon wafers have been used as substrate, having a thickness of 350 ± 15 ␮m. Prior to the thin film deposition, the substrates have been rinsed with acetone and isopropyl alcohol, followed by a short dip in buffered hydrofluoric acid to remove the native oxide from the wafer surface. The deposition experiments were performed using an Oxford Plasmalab 100 System ICP-CVD (inductively coupled plasmachemical vapor deposition) downstream reactor, with an additional capacitively coupled RF plasma source connected to the substrate electrode. Both plasma sources operate at a frequency of 13.56 MHz, whereas only the ICP source has been activated during deposition. The process chamber was evacuated to a base pressure of 9 × 10−5 Pa, before nitrogen and argon were introduced together with silane (SiH4 ) via a gas distribution ring arranged next to the substrate electrode. Flow rates of the SiH4 precursor and the Ar dilution gas were kept at constant values of 48 and 12 sccm, respectively. In order to screen a wide parameter field with moderate effort, a two-level full factorial design of experiments (DoE) was performed including two additional runs with parameters set at the mid-points. The investigated process parameters were ICP plasma power PICP , total pressure pdep in the deposition chamber, substrate temperature T, and the flow rate ratio R of N2 and SiH4 . This yielded an experimental design with 18 runs in total with the deposition parameters listed in Table 1. When depositions were performed at a substrate temperature of 120 ◦ C, the wafers were clamped with a glass ring and the heat transfer was improved by applying a He backflow, whereas at 350 ◦ C the wafer clamp was removed

349

Table 1 Run list. Parameter set

T [◦ C]

pdep [Pa]

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18

120 120 120 120 120 120 120 120 350 350 350 350 350 350 350 350 120 350

0.53 0.53 0.53 0.53 1.33 1.33 1.33 1.33 0.53 0.53 0.53 0.53 1.33 1.33 1.33 1.33 0.93 0.93

PICP [W] 500 500 1000 1000 500 500 1000 1000 500 500 1000 1000 500 500 1000 1000 750 750

RN2 /SiH4 0.5 2 0.5 2 0.5 2 0.5 2 0.5 2 0.5 2 0.5 2 0.5 2 1 1

and radiative heat transfer was used. The deposition runs at both temperature levels were carried out randomly in order to reduce any potential biasing of the results. The deposition time was held constant at 15 min. 2.2. Analytical methods The thickness d and refractive index n (at  = 632.8 nm) of the deposited thin films were determined using the spectral reflectance method. The measurements were conducted with a Filmetrix F20 film thickness measurement instrument, which determines the reflectance between 200 and 1100 nm. Mechanical film stress was measured using an E+H Metrology MX-203 capacitive wafer geometry gauge, yielding the curvature of the wafer prior and after thin film deposition. For further evaluation purposes, a modified version [34] of Stoney’s formula [35] was used. Temporal evolution of the residual film stress at room temperature was monitored by re-measuring of the wafer bow. Wet chemical etching tests have been performed on each coated wafer in a 1:7 mixture of concentrated (50%) hydrofluoric acid (HF) and de-ionized water. The wet etch rates (WER) were determined by measuring the resulting step height by a Veeco Dektak surface profilometer. Fourier-transform infrared (FT-IR) measurements on a Bruker Vector 22 (Software: Opus 5.5) infrared spectroscope in transmission mode between 400 and 4000 cm−1 with a resolution of 2 cm−1 were applied to investigate the chemical composition of the thin films. The infrared spectrum of a blank silicon wafer was taken as reference. After baseline correction of the measured spectra, the absorption peaks were normalized by the film thickness and by the height of the silicon lattice absorption peak at 610 cm−1 , assuming an equal thickness of the substrates. X-ray photoelectron spectroscopy (XPS) measurements were conducted with a Thermo Scientific Theta Probe small spot XPS instrument to study the chemical composition of some selected SiNx :H films. Both depth profiles throughout the thickness of the film and surface-near scans were performed to get detailed information of the chemical bonding state. 3. Results and discussion 3.1. Deposition rate and refractive index As expected, the ICP power has a strong influence on the deposition rate rdep , since a higher plasma power level leads to an enhanced density of reactive species originating from the different precursor gases. The higher availability of these species results

350

D. Dergez et al. / Applied Surface Science 284 (2013) 348–353

Fig. 1. The refractive index n, measured at  = 632.8 nm, as a function of the deposition rate, in the case of layers deposited with a N2 flow rate of 6 sccm (black squares), 12 sccm (red circles), and 24 sccm (blue triangles). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

in higher reaction rates, not only on the surface of the substrates, but also on the reactor walls [31,32]. Independent of other deposition parameters an average rdep of 9.4 and 16.5 nm/min results at 500 and 1000 W, respectively. Another significant parameter affecting the deposition rate is the N2 flow rate [31]. It becomes immediately clear that there is a significant difference between depositions carried out at low (average rdep : 8.18 nm/min at 6 sccm) and those done at high flow rates of nitrogen (average rdep : 16.6 and 17.7 nm/min at 12 and 24 sccm of N2 ). From this fundamental difference, the layers can be grouped into two basic types. Type I denotes the films deposited with a N2 flow rate of 12 or 24 sccm, while Type II represents those synthesized with 6 sccm of N2 . In the case of Type I films the deposition rate slightly decreases at enhanced substrate temperatures from about 18.2 nm/min at 120 ◦ C down to 16.8 nm/min at 350 ◦ C in contrast to Type II films where this parameter shows no significant effect (T = 120 ◦ C: average rdep = 8.1 nm/min, T = 350 ◦ C: average rdep = 8.2 nm/min). Also variation of pdep did not result in an evident change of rdep either in the case of Type I or Type II layers. In Fig. 1, the inverse correlation between the refractive index and the deposition rate dominated in turn by the nitrogen flow rate is shown, as reported in [31]. The refractive index n in the case of Type I films is in the range of 1.96–1.98 (cf. nSi3N4 ∼ 2.0). In contrast, this material parameter is measured for Type II films between 2.1 and 2.7. For the latter samples, two further effects can be observed: there is a significant increase of the IOR at elevated substrate temperatures (n = 0.15 ± 0.03 from T = 120 to 350 ◦ C), or by increased total pressure during deposition (n = 0.37 ± 0.03 from pdep = 0.53 to 1.33 Pa,). These substantial differences indicate the presence of major compositional differences within the various Type II films.

Fig. 2. The total film stress as deposited (black squares) and after one week (red circles) as a function of deposition rate. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

and 350 ◦ C), or pdep (average: −14.3 and +18.5 MPa at 0.53 and 1.33 Pa). In the case of Type II films, both higher T (average: −835 and −1087 MPa at 120 and 350 ◦ C) and higher PICP values (average: −890 and −1032 MPa at 500 and 1000 W) induce stronger compressive stress levels in the “as deposited” state. Stress measurements up to 3 weeks after deposition where the samples are exposed to atmospheric conditions demonstrate that the level of compressive stress increases significantly at Type I films, while at Type II films this material parameter stays almost constant in the investigated time frame. The different characteristics are illustrated in Fig. 2. Basically, the drift in film stress measured 7 days after deposition is higher when the films are deposited at higher T (average  = 101 and 158 MPa for films deposited at 120 and 350 ◦ C). Fig. 3 shows the difference between the temporal evolution of the film stress for Type I and Type II samples. At the latter samples, the drift rate is highest right after deposition, becoming gradually slower with time. Due to the increase in compressive stress, it is reasonable to assume that the modifications in the films are based on diffusion processes, predominantly the penetration of humidity and oxygen into the layers originating from the ambient atmosphere at room temperature [36], and the subsequent oxidation until the saturation limit is reached.

3.2. Film stress and drifting behavior As shown in the previous section, the N2 flow rate has the strongest impact on the film stress  in the “as-deposited” state. For Type I films,  is determined between −200 and +100 MPa, demonstrating a low stress behavior, while Type II samples are under strong compressive stress ranging between −1350 and −700 MPa. The film stress of Type I films is affected in the tensile direction, when increasing either T (average : 0.8 and 58.6 MPa at 120

Fig. 3. Temporal evolution of the total film stress in the case of stable and drifting samples. The inserted lines serve as guides to the eye.

D. Dergez et al. / Applied Surface Science 284 (2013) 348–353

Fig. 4. Wet etch rate in hydrofluoric acid as a function of deposition rate, showing the N2 flow rate of 6 sccm (black squares), 12 sccm (red circles), and 24 sccm (blue triangles). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

3.3. Etch rate in wet hydrofluoric acid Fig. 4 shows the wet etch rate as a function of the deposition rate. At low values up to 10 nm/min, the material removal in HF is in the same order of magnitude and hence, very slow. Above, the etch rate increases abruptly by several orders of magnitude up to about 1000 nm/min being partially even higher than those reported for PECVD SiO2 at comparable conditions [37]. This result indicates a substantial difference in the chemical composition and the microstructure between the two types of a-SiNx :H films. For Type I layers, a higher T has a negative (∼1890 and 740 nm/min at 120 and 350 ◦ C), while a higher pdep has a positive (∼1170 and 1700 nm/min at 0.53 and 1.33 Pa) effect on the wet etch rate. These results are attributed to the enhanced adatom mobility during film growth resulting in a denser and more compact microstructure of the films and increased resistance against HF exposure [9]. Although the Type II films exhibit largely different film properties they show the same behavior, but at substantially reduced values (11.3 and 2.7 nm/min for a deposition temperature of 120 and 350 ◦ C and 10.5 and 3.4 nm/min for 0.53 and 1.33 Pa). Independent of other deposition parameters the impact of PICP is negligible on the etch rate (6.2 and 6.6 nm/min for samples deposited at 500 and 1000 W). These results suggest that the PICP has only minor influence on the chemical bonding configuration of the a-SiNx :H layers. 3.4. Chemical analyses It was found that the FT-IR spectra of Type I and Type II layers exhibit very distinct characteristics. Therefore, one typical representative for each type is displayed in Fig. 5. The most prominent differences are the presence of the Si–H stretch mode peak at about 2150 cm−1 in Type II and the absence in Type I films, and vice versa in the case of the N–H (Si–NH2 ) stretch peak at approximately 3335 cm−1 [38,39]. After normalizing the infrared absorption spectra, the straightforward way would be the deconvolution of the FT-IR spectra into the individual absorption peaks, in order to find correlation between the deposition parameters and chemical bond densities in the sample. Due to the fact, that a multitude of these peaks can occur in the absorption spectra of a-SiNx :H, with the exact quantity and position of these being unknown the aforementioned approach is regarded to be highly imprecise considering the partly diffuse

351

Fig. 5. Comparison of the FT-IR spectra of typical example layers prepared using a deposition process having a N2 flow rate of 6 sccm and 24 sccm, respectively. The graph shows the details of the Fourier infrared spectra highlighting the most prominent absorption peaks.

nature of the spectra (i.e. several absorption bands are located in close vicinity to each other). Therefore, the corresponding heights of the various prominent absorption peaks were taken as a significant indicator. This approach does not yield a highly quantifiable result regarding the densities of the different bond species in the sample. It shows, however, the trends how the deposition parameters affect the compositional changes in the a-SiNx :H thin films. The absorption peak of the single crystalline Si lattice at 610 cm−1 as well as the Si–O stretching mode absorption peak at 1108 cm−1 is detectable in both layer types. The latter is clearly more pronounced in the case of Type I films as for Type II, supporting the assumption of oxygen incorporation into these layers. In Type I films, the height of the N–H (Si–NH2 ) peak (in arbitrary units) is inversely correlated with the increase of T (average peak height: 0.41 and 0.32 at 120 and 350 ◦ C), but decreases with increasing pdep (0.33 and 0.43 at 0.53 and 1.33 Pa) and increasing PICP (0.30 and 0.47 at 500 and 1000 W). In the case of Type II films the height and position of the Si–H stretch mode peak are both influenced by the reactor pressure during deposition. The larger the pdep value, the higher is the absorption peak (average: 0.24 at 0.53 Pa and 0.33 at 1.33 Pa). On the other hand, with increasing reactor pressure the peak is shifted to lower wave numbers (average value: 2164 and 2124 cm−1 at 0.53 and 1.33 Pa). Since this shift shows no correlation with the residual film stress, it is attributed to differences in the chemical composition (see the dependence of n on pdep for Type I films), indicated by a more pronounced presence of Si–H species (i.e. Si3 –Si–H, Si2 –Si–H2 and Si–Si–H3 ) having a Si-richer configuration. In addition, the sharp H2 –Si–N2 peak characteristics in the case of Type I/II films is different; for Type I the Si–N asymmetric stretching mode peak at 902 cm−1 is dominant, while for Type II films the Si3 N4 related peak at 870 cm−1 is more pronounced (Fig. 5). But, the height of these peaks is influenced by the other deposition parameters as well (see Table 2). Increasing the substrate temperature T seems to have a positive effect on the silicon–nitrogen bond densities, while a higher pdep or PICP leads to lower Si–N bonddensities. As seen in Fig. 6, there is an inverse correlation between the heights of N–H (Si–NH2 ) and Si–N asymmetric stretch absorption peaks, which means that the more nitrogen is incorporated in the film via Si–N bonds, the lower the N–H bond density. Furthermore, these results indicate that a higher T causes an enhanced desorption rate of hydrogen from the forming films, thus resulting in a higher density of Si–N, and a lower density of N–H bonds. In contrast, increasing the total pressure results in the formation of

352

D. Dergez et al. / Applied Surface Science 284 (2013) 348–353

Table 2 Influence of the deposition parameters on the height of different silicon-nitrogen infrared absorption peaks. Peak height [a.u.]

Type I: 902 cm−1 Type II: 870 cm−1

T [◦ C]

pdep [Pa]

PICP [W]

120

350

0.53

1.33

500

1000

0.88 1.09

0.94 1.20

0.97 1.25

0.82 1.04

1.09 1.27

0.70 1.03

Fig. 6. Height of the N H absorption peak (3335 cm−1 ) as a function of the Si N Si absorption peak (902 cm−1 ) in Type I a-SiNx :H films.

layers, which are hydrogen-rich due to the higher amount of SiH4 available during the deposition process. XPS surface scans show a difference in the chemical shift of the silicon peaks in the case of Type I and Type II films (see Fig. 7). Even a straight-forward deconvolution of the peak into three Gaussian peaks associated with pure silicon, silicon nitride and silicon dioxide [40], displays some substantial differences between the two layer types. In Type I films (Fig. 7b) only a very low portion is pure silicon, the majority of the signal originates from electrons associated with Si–N bonds. In contrast, Type II samples (Fig. 7a) show a relatively large portion related to pure silicon in their Si peak. The tendency is clear: in Type I layers silicon atoms form on average more bonds with nitrogen atoms than in Type II films. This result is consistent with the dominance of different Si–N bonds in the infrared spectra of the both types of layers. The surfaces of the samples are in both cases strongly oxidized; therefore a high portion of Si is in a SiO2 bonding configuration when measuring only surface near.

XPS depth profiling revealed the chemical composition of the samples throughout the film thickness. The N/Si ratio averaged over the whole thickness of the samples was found to be between 0.57 and 0.65 for Type I, and between 0.35 and 0.52 for Type II films, meaning that both types of layers are silicon rich, having an N/Si ratio far away both from the stoichiometric ratio of 1.33 for Si3 N4 and the N/Si ratio of 1–4 in the precursor gas mixtures. These results suggest that most probably due to the higher dissociation rate of SiH4 the incorporation of Si atoms into the a-SiNx :H films is much more efficient compared to that of N atoms from a N2 precursor. The atomic concentration of oxygen in Type I layers was in the range of 11.6–17.7 at.%. Due to the fact that a relatively long time has passed between the deposition and the XPS analysis, this oxygen concentration does not show considerable variation throughout the sample thickness, except for the vicinity of the surface, where it can be as high as 27–29 at.%. In Type II films the average oxygen concentration was found to be below 2 at.%, with surface concentrations as high as 23–25 at.%. Note that these values are valid only if we consider a purely Si/N/O system. These results clearly testify to the theory of the oxidation-induced drifting behavior of the mechanical stress of Type I a-SiNx :H films (see Fig. 3). Since it is known from the FT-IR spectra that a considerable amount of hydrogen is incorporated in the layers, which is invisible to XPS, the actual atomic concentrations are different.

4. Conclusions In this work, a DoE study on a-SiNx :H thin films deposited by ICPCVD was performed varying systematically key process variables such as plasma power, substrate temperature, the back pressure in the deposition chamber and the N2 /SiH4 flow rate ratio. In the scope of this study, especially the latter parameter is of utmost importance in the deposition process, having a substantial impact on the deposition rate and on the mechanical, optical and chemical properties of the layers. Thin films deposited with a higher N2 /SiH4 flow rate ratio have low (<2) refractive indices, are deposited at higher rates (>10 nm/min), are etched extremely fast in hydrofluoric acid and yield low mechanical stress levels (−50 to +200 MPa) in the “as deposited” state, but exhibit a strong drift toward compressive

Fig. 7. XPS measurement data, showing the details of the Si absorption peak of example layers deposited with different flow rates of N2 (a: 6 sccm; b: 24 sccm).

D. Dergez et al. / Applied Surface Science 284 (2013) 348–353

stress values. In contrast, layers deposited with a lower N2 /SiH4 flow rate feature high refractive indices (>2.1), and are deposited at lower rates (<10 nm/min). They are etched very slowly in HF and are under strong compressive stress (−1300 to −700 MPa) in the “as deposited” state, but show stable stress values as a function of time. Fourier-transform infrared spectroscopy and X-ray photoelectron spectroscopy measurements show both significant differences in the chemical composition of the two types of a-SiNx :H thin films. In both sample types the hydrogen is incorporated predominantly either in N–H or in Si–H bonds. Atomic ratios of Si/N obtained by XPS measurements show that both types of a-SiNx :H films are silicon-rich, with significantly higher nitrogen concentration in films deposited with a higher N2 flow. In these samples a remarkably high amount of oxygen is present, supporting the assumption that the drift of the intrinsic film stress is directly linked to the incorporation of oxygen or humidity from the ambient into the films. Detailed analyses of the different silicon–nitrogen IR absorption bands and the silicon peak in the XPS spectra both show that an increase of the N2 flow rate in the deposition process leads not only to an increase of the nitrogen concentration in the sample, but also to a bonding configuration where the Si atoms are bonded on average with a higher probability to N atoms. All in all, it can be summarized that with a tailored fabrication process comprising the deposition parameters as well as carefully selected post deposition annealing processes the realization of a-SiNx :H films are feasible offering simultaneously key features for MEMS such as a low mechanical stress value in combination with low drift effects and a high resistance against hydrofluoric acid. Acknowledgement The research leading to these results has received funding from the European Union’s Seventh Framework Program managed by REA-Research Executive Agency http://ec.europa.eu/research/rea (FP7/2007-2013) under grant agreement no. 286146. References [1] C.H. Ng, K.W. Chew, S.F. Chu, IEEE Electron Device Letters 24 (8) (2003) 506–508. [2] S.S. Sandeep, K. Warikoo, A. Kottantharayil, Proceedings of the 38th IEEE Photovoltaic Specialist Conference 2012, 2012, p. 001102. [3] A. Khan, J. Philip, P. Hess, Journal of Applied Physics 95 (4) (2004) 1667–1672. [4] S. Hasegawa, Y. Amano, T. Inokuma, Y. Kurata, Journal of Applied Physics 72 (1992) 5676. [5] M.P. Hughey, R.F. Cook, Journal of Applied Physics 97 (2005) 114914. [6] P. Morin, G. Raymond, D. Benoit, P. Maury, R. Beneyton, Applied Surface Science 260 (2012) 69–72.

353

[7] K.B. Sundaram, R.E. Sah, H. Baumann, K. Balachandran, R.M. Todi, Microelectronic Engineering 70 (2003) 109–114. [8] K.B. Gavan, H.J.R. Westra, E.W.J.M. van der Drift, W.J. Venstra, H.S.J. van der Zant, Microelectronic Engineering 86 (2009) 1216–1218. [9] W.T. Ho, H.J. Lee, Q.Z. Zhang, A. Tan, Y.W. Lim, R. Nagarajan, Defense Science Research Conference and Expo, 2011, p. 5. [10] H. San, Z. Deng, Y. Yu, G. Li, X. Chen, Journal of Micromechanics and Microengineering 21 (2011) 125019. [11] G.E. Fantner, G. Schitter, J.H. Kindt, T. Ivanov, K. Ivanova, R. Patel, N. HoltenAndersen, J. Adams, P.J. Thurner, I.W. Rangelow, P.K. Hansma, Ultramicroscopy 106 (2006) 881–887. [12] C.H. Schwalb, C. Grimm, M. Baranowski, R. Sachser, F. Porrati, H. reith, P. Das, J. Müller, F. Völklein, A. Kaya, M. Huth, Sensors 10 (2010) 9847–9856. [13] R. Tiwari, S. Chandra, Advanced Materials Research 254 (2011) 187–190. [14] M.A. Signore, A. Sytchkova, D. Dimaio, A. Capello, A. Rizzo, Optical Materials 34 (4) (2011) 632–638. [15] J.H. Kim, K.W. Chung, Journal of Applied Physics 83 (11) (1988) 5831–5891. [16] J.G.E. Gardeniers, H.A.C. Tilmans, C.C.G. Visser, Journal of Vacuum Science and Technology A 14 (5) (1996) 2879–2892. [17] A. Stoffel, A. Kovács, W. Kronast, B. Müller, Journal of Micromechanics and Microengineering 6 (1996) 1–13. [18] S.B. Patil, A. Kumbhar, P. Waghmare, V.R. Rao, R.O. Dusane, Thin Solid Films 395 (2001) 270–274. [19] T.M. Klein, T.M. Anderson, A.I. Chowdhury, G.N. Parsons, Journal of Vacuum Science and Technology A 17 (1998) 108. [20] G.N. Parsons, J.H. Souk, J. Batey, Journal of Applied Physics 70 (1991) 1553. [21] J.A. Taylor, Journal of Vacuum Science and Technology A 9 (4) (1991) 2464–2468. [22] D. Criado, I. Pereyra, M.I. Alayo, Materials Characterization 50 (2003) 167–171. [23] R.-Y. Tsai, L.-C. Kuo, F.C. Ho, Applied Optics 32 (28) (1993) 5561–5566. [24] M. Matsumoto, Y. Inayoshi, S. Murashige, M. Suemitsu, S. Nakajima, T. Uehara, Y. Toyoshimta, Journal of Vacuum Science and Technology B 27 (1) (2009) 223–225. [25] B. Kim, S. Kim, Thin Solid Films 517 (2009) 4090–4093. [26] C. Aguzzoli, C. Marin, C.A. Figueroa, G.V. Soares, I.J.R. Baumvol, Journal of Applied Physics 107 (2010) 073521. [27] C.M.M. Denisse, K.Z. Troosr, F.H.P.M. Habraken, W.F. van der Weg, M. Hendriks, Journal of Applied Physics 60 (1986) 2543. [28] H.F. Dekkers, G. Beaucarne, H. Miller, H. Charifi, A. Slaoui, Applied Physics Letters 89 (2006) 211914. [29] S.P. Singh, P. Srivastava, S. Ghosh, S.A. Khan, G. Vijaya Prakash, Journal of Physics: Condensed Matter 21 (2009) 095010. [30] S.H. Lee, I. Lee, J. Yi, Surface and Coatings Technology 153 (2002) 67–71. [31] S.-S. Han, B.-H. Jun, K. No, B.-S. Bae, Journal of The Electrochemical Society 145 (1998) 652. [32] L. da Silva Zambom, R.D. Mansano, R. Furlan, Vacuum 65 (2002) 213. [33] A. Kshirsagar, P. Nyaupane, D. Bodas, S.P. Duttagupta, S.A. Gangal, Applied Surface Science 257 (2011) 5052–5058. [34] P.H. Townsend, D.M. Barnett, T.A. Brunner, Journal of Applied Physics 62 (1987) 4438. [35] G.G. Stoney, Proceedings of the Royal Society: London 82 (1909) 172. [36] W.-S. Liao, S.-C. Lee, Journal of the Electrochemical Society 144 (1997) 1477. [37] K.R. Williams, K. Gupta, M. Wasilik, Journal of Micrelectromechanical Systems 12 (6) (2003) 761. [38] Q. Xu, Y. Ra, M. Bachman, G.P. Li, Journal of Vacuum Science and Technology A 27 (2009) 145. [39] Z. Lu, P. Santos-Filho, G. Stevens, M.J. Williams, G. Lucovsky, Journal of Vacuum Science and Technology A 13 (1995), 607. [40] M. Matsuoka, S. Isotani, W. Sucasire, L.S. Zambom, K. Ogata, Surface and Coatings Technology 198 (2010) 2923.