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. Author manuscript; available in PMC: 2020 Jul 1.
Published in final edited form as: J Biomed Mater Res B Appl Biomater. 2018 Oct 15;107(5):1654–1661. doi: 10.1002/jbm.b.34258

Effect of oxidation on intrinsic residual stress in amorphous silicon carbide films

Felix Deku 1,#, Shakil Mohammed 2, Alexandra Joshi-Imre 1, Jimin Maeng 1, Vindhya Danda 1, Timothy J Gardner 3, Stuart F Cogan 1
PMCID: PMC6465170  NIHMSID: NIHMS993186  PMID: 30321479

Abstract

The change in residual stress in PECVD amorphous silicon carbide (a-SiC:H) films exposed to air and wet ambient environments is investigated. A close relationship between stress change and deposition condition is identified from mechanical and chemical characterization of a-SiC:H films. Evidence of amorphous silicon carbide films reacting with oxygen and water vapor in the ambient environment are presented. The effect of deposition parameters on oxidation and stress variation in a-SiC:H film is studied. It is found that the films deposited at low temperature or power are susceptible to oxidation and undergo a notable increase in compressive stress over time. Furthermore, the films deposited at sufficiently high temperature (≥ 325 °C) and power density (≥ 0.2 Wcm−2) do not exhibit pronounced oxidation or temporal stress variation. These results serve as the basis for developing amorphous silicon carbide based dielectric encapsulation for implantable medical devices.

Keywords: Amorphous silicon carbide, residual stress, a-SiC oxidation, air stability, PECVD

1. Introduction

Thin-film coatings of amorphous silicon carbide (a-SiC), have been used as a protective coating for implantable neural devices [1–6], orthopedic implants [7] and coronary stents [8–10]. The a-SiC is usually deposited by plasma enhanced chemical vapor deposition (PECVD), which allows for moderately low-temperature deposition (≤ 400 °C), high deposition rates and precise control of residual film stress. Residual stress may be intrinsic, arising from hydrogen incorporated in the film during deposition, or extrinsic, due to differential thermal expansion between the a-SiC film and the coated substrate. Although most studies have associated residual stress of PECVD a-SiC (a-SiC:H) films to deposition conditions, we are interested in the stress evolution post-deposition in response to the ambient environment as changes in the residual stress state can alter the planarity of implantable thin-film devices fabricated from a-SiC:H films[5]. Loss of planarity may prevent devices from being inserted into tissue or, if residual stresses change after implantation, there is the possibility of tissue damage or unpredictable changes in device performance.

The intrinsic stress in a-SiC:H is usually compressive arising from the incorporation of hydrogen as evident from the Si–H peak intensities in Fourier transform infrared (FTIR) spectra and the monotonic relationship between film stress and the peak intensity [11,12]. Hydrogen incorporation occurs because of incomplete decomposition of the hydrogen-containing precursors. These weakly bonded hydrogen atoms permit the carbide surface to oxidize to a few nanometers when exposed to ambient air[13]. The reaction kinetics are faster in films with high carbon and hydrogen compositions[14] and films with higher micro-void densities[15]. Previous published research on controlling residual stress has identified power density, temperature, gas ratio, pressure as well as He or Ar dilution of the reactive gas as factors influencing the residual stress in a-SiC:H films [16–21].

In the present study, the time-dependent change of residual stress in a-SiC:H films is investigated. The relationship between film oxidation and stress change, along with the effects of different deposition conditions is examined. X-ray photoelectron spectroscopy (XPS) and FTIR spectroscopy were used to understand the molecular origin of stress variation over time in these films. Evidence of a-SiC:H films reacting with water vapor and oxygen from ambient environments is presented. This study is intended to inform the selection of deposition conditions to produce low-stress and air stable a-SiC:H dielectric coatings for chronic implantable medical devices.

2. Materials and Methods

2.1. Amorphous silicon carbide film deposition

A PlasmaTherm Unaxis 790 series PECVD system with a platen area of approximately 1000 cm2 was used to deposit the a-SiC:H films. The films were deposited on 100 mm Si (100) prime wafers (University Wafer) along with double-sided polished (DSP) high resistivity Si wafers (Silicon Valley Microelectronics) for FTIR measurements. Three sets of films were deposited for the study with the deposition condition ranges selected based on earlier studies [1,11,12]. The first set was deposited at different temperatures, from 150°C to 350 °C at a fixed power density of 0.27 Wcm−2. The second set was deposited at different power densities, ranging from 0.05 Wcm−2 to 0.32 Wcm−2, at a fixed temperature of 350 °C. The chamber was maintained at a pressure of 1000 mTorr with the total gas flow fixed at 800 sccm using Ar as the carrier gas. For all depositions, a SiH4:CH4 gas flow rate ratio of 1:3 was used. The thickness of the films was maintained close to 1 µm by varying deposition time. The a-SiC:H films were exposed to air at room temperature, and the intrinsic stress in these films was monitored once every 7 days over a period of 28 days.

The third set of films were deposited at 150 °C and 350 °C, and a power density of 0.27 Wcm−2, and were exposed to various environments including Dry Air (7 ppm of H2O, ∼ 20% O2), deionized (DI) water, N2 (99.998%), and O2 (99.994%). Changes in intrinsic stress in these films was monitored for 24 hours.

2.2. Amorphous silicon carbide film characterization.

A Nanometric NanoSpec 6100 analyzer was used to measure the average film thickness from 20 different points randomly across the wafer. A Toho Technology FLX-2320-S system was used to measure the film stress. The stress values were estimated from changes in the radius of curvature of the silicon substrate due to the deposition of the a-SiC:H film using the Stoney equation[22] and have an accuracy of <2.5% of the measured stress for values above 40 MPa based on manufacturer’s specifications. FTIR transmittance spectra were recorded using a ThermoElectron Nicolet iS50 spectrometer with a deuterated triglycine sulfate (DTGS) detector at a wavelength resolution of 4 cm−1. A PHI VersaProbe II spectrometer equipped with an Al Kα monochromated X-ray source (hν = 1486.7 eV) was used to collect the XPS data. The films were pre-sputtered with 1 keV Ar ions for 10 minutes to remove a thin layer of the a-SiC:H surface prior to spectral collection. Voigt line shapes in conjunction with a Shirley background subtraction utilizing “AAnalyzer” software was used to fit the spectra[23].

3. Results

3.1. Residual stress stability

The residual stress of air-exposed a-SiC:H films was monitored over a period of 28 days. The time-evolution of the residual stress as a function of deposition temperature and power density is shown in Figure 1. Figure 1(a) shows the residual stress in films deposited at temperatures ranging from 150°C to 350 °C at a constant power density of 0.27 Wcm−2. The residual stress in the as-deposited films (at day 0) becomes less compressive as deposition temperature is increased (−230 MPa for TD = 150 °C to −100 MPa for TD = 350 °C), similar to results reported elsewhere[21]. The stress in the films deposited at 325 °C and 350 °C remained approximately unchanged over time. However, in the lower-temperature films (TD = 250 °C and TD = 150 °C), a large increase in compressive stress was observed. Between the initial measurement (day 0) and the first measurement at day 7, the magnitude of the compressive stress increased from −163 to −203 MPa (24.5% increase), and from −226 to −356 MPa (57.5% increase), for TD=250 °C and TD=150 °C, respectively. The stress became increasingly compressive as the air-exposure period was extended to 28 days. Figure 1(b) shows the stress variation in films deposited at power densities ranging from 0.05 Wcm−2 to 0.32 Wcm−2 with the platen temperature controlled to 350 °C.

Figure 1.

Figure 1.

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Residual stress in air-exposed a-SiC:H films deposited at (a) various deposition temperatures and power density of 0.27 Wcm-2 and (b) various power densities at a deposition temperature of 350 °C, monitored over a period of 28 days. Data are reported as mean±1 std. dev., n=3.

The compressive stress of as-deposited films increases with increasing deposition power (−35 MPa for 0.05 Wcm−2 to −125 MPa for 0.32 Wcm−2), similar to previous published studies[21]. The present work demonstrates that the power density at which the films are deposited also influences the changes in film stress over time. The residual stress remains relatively unchanged in the films deposited at power density equal to or greater than 0.2 Wcm−2, while more pronounced changes are observed at low-power deposition conditions.

3.2. XPS and FTIR Analysis

XPS spectra collected 28 days post-deposition are plotted in Figure 2. The Si 2p, C 1s and O 1s regions were charge compensated by shifting the data to center the O 1s peak at 532.2 eV prior to fitting the spectra. The Si 2p region (Figure 2 (a)) was fitted using doublet peaks considering the spin-orbit splitting of the orbital, but only the convolution of the doublet is displayed for clarity. The doublet peak represented by Si 2p3/2 and Si 2p1/2 are separated only by 0.6 eV and have an intensity ratio of 2:1[24].

Figure 2.

Figure 2.

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The effect of deposition temperature on the XPS spectra of a-SiC:H air-exposed for 28 days. The dashed lines represent the peak positions of (a) Si 2p, (b) C 1s and (c) O 1s.

The overall Si 2p peak position was observed at lower binding energies for higher deposition temperatures. For films deposited at TD = 350 °C and 325 °C, a single peak centered at ~100.2 eV, typical of the Si–C bond[25], was observed. Two distinct fitted peaks were observed for films deposited at 150oC and 250oC. The primary peak appears at 100.8 eV in the 250 °C film and at 101 eV in the 150 °C film. These peaks were assigned to the Si–O–C bond, as the peak position has previously been associated with SiOC3[26] or Si2OC2.2[27] bonding. The secondary peak at 102.1 eV can be assigned to either O-rich SiO3C[28] or a sub-stoichiometric SiOx phase[29].

The C 1s spectra, shown in Figure 2(b), exhibits two peaks for each film. The peak at 283.3 eV in both 350 °C and 325 °C films is due to C–Si bonding. The C–Si peak shifts to higher energies at lower deposition temperatures, which is attributed to the presence of C–Si–O bonding structures. The shift of the C–Si peak in the C 1s region corresponds to the observed shift of Si–C peaks at lower deposition temperatures in the Si 2p region. The peak at 284.7 eV is due to C–C bonding and does not shift with deposition temperature but appears to increase in intensity relative to the C-Si-O bonding as the deposition temperature is decreased.

The O 1s region of both the 350 °C and the 325 °C deposited films show very low peak intensities, as shown in Figure 2(c). The intensity of O 1s increases in the 250 °C films and further increases in the 150 °C films. Two peaks were identified for the low temperature films, with the primary peak located at 532.2 eV arising from O–Si–C bonding, and the secondary peak at 533.2 eV arising from O–Si bonding. The O–Si peaks in the O 1s region corresponds with the Si–O peaks in the Si 2p region for the low-temperature deposited films.

The film composition x, defined as the ratio of the number of silicon atoms to the total number of silicon and carbon atoms, x = Si / (Si+C) was estimated using the atomic concentration calculated from the XPS spectra. The nominal composition was approximately x=0.5, with a trend towards increasing Si concentration as the deposition temperature was increased.

XPS spectra showing the compositional changes in the film deposited at TD = 150 °C, at time points 0 (as-deposited), 7 and 28 days post deposition are shown in Figure 3. The peak intensity of the O–Si bonding in the O 1s region and Si–O bonding in the Si 2p regions increased over time. Also, the O–Si–C peak intensity in the O 1s region increased and the Si–O–C peak intensity in the Si 2p region decreased as time progressed. The peak positions of the Si 2p and the C 1s did not change relative to O 1s over the observed period.

Figure 3.

Figure 3.

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XPS of a-SiC:H deposited at TD = 150 °C and characterized after 0, 7, and 28 days of air exposure. The dashed lines represent the peak positions of (a) Si 2p, (b) C 1s, and (c) O 1s.

Figure 4 shows the infrared (IR) transmittance spectra of the silicon carbide films that were deposited on high-resistivity (3 – 5 kΩ-cm) double-side polished (DSP) silicon wafers. The changes in the spectra were monitored immediately after deposition and then 1, 2, 7 and 28 days post-deposition. The peaks observed conform to those published in the literature[30–34]. The peak at 607 cm−1 is due to the Si–Si phonon absorption from the substrate[1].

Figure 4.

Figure 4.

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Representative FTIR spectra of air-exposed a-SiC:H films deposited at various temperatures are recorded (a) after deposition, (b) 7 days post-deposition, and (c) 28 days post-deposition. The evolution in the FTIR spectra for the a-SiC:H deposited at TD=150oC over the 28-day post-deposition period are compared directly in (d). The inset in figure 4(d) shows the shift and change in intensity of the Si–H stretching band.

The most intense peak at 780 cm−1 has been reported to be the Si–C stretching mode, and the shoulder at ∼1020 cm−1 has been identified as either a Si–CH2 or a Si–O–Si stretching, or both[30–32]. The peak at ~2100 cm−1 arises from the Si–H stretching, and various C–Hn stretching modes appear at ∼2900 cm−1 [11,35]. Figure 4 (a), (b) and (c) show spectra collected for all deposition temperature (TD) films after deposition (day 0), day 7 and day 28 respectively, while Figure 4 (d) shows an overlay of the FTIR spectra for films with a TD = 150 °C as time progressed. For the as-deposited a-SiC:H films, the peak intensities for the Si–C bond were almost the same regardless of the deposition temperature, while the Si–O–Si, Si–CH3 and Si–H peaks increased with decreasing temperature. Also, the Si–H peak shifts to higher frequency as the deposition temperature decreases. The peak intensities and peak positions did not change for films deposited at higher temperatures (TD = 350 °C and TD = 325 °C) on exposure to air. The films deposited at lower temperature (especially at TD = 150 °C) exhibit a gradual decrease of the Si–H peak intensity at ∼2100 cm−1, with a peak shift to higher frequencies (inset), and a noticeable increase of the Si–O–Si peak intensities at the ∼1020 cm−1 shoulder as shown in Figure 4(d).

Examination of Figure 4(d) also reveals the growth of a broad band at ∼3425 cm−1 which has been previously assigned to the hydroxyl (O–H) group [36,37]. A plot of the Si–H peak position as a function of time for the TD= 150°C film is shown in Figure 5, including a Gaussian fit (inset) suggesting that this absorption band arises primarily from Si–H vibration with a small contribution from Si–H2 [38,39].

Figure 5.

Figure 5.

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The FTIR shift in the Si–H peak for the film deposited at TD = 150 °C is shown in the plot. (Inset) The de-convoluted peaks show Si–H and Si–H2 stretching modes

Furthermore, the area under the Si–O–Si and Si–C absorption band was integrated and shown in Figure 6 as a function of air exposure time for the TD = 150 °C film. The Si–C bonding remains constant, whereas the Si–O–Si bonding is observed to increase over time.

Figure 6.

Figure 6.

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Area of the Si–C, and Si–O infrared absorption bands for TD = 150 °C a-SiC:H as a function of air exposure time. The intensity of the Si–C band remains unchanged, while the Si–O band increased in intensity with time corresponding to observed changes in the O 1s XPS data in Fig. 3c.

3.3. Effect of Environment on Stress Change

To evaluate the relative effects of different environmental conditions on a-SiC:H films stress, a short-term study was conducted in which a set of films deposited at TD = 150 °C and 350 °C were monitored over a 24-hour period under exposure to laboratory air, dry air (7ppm of H2O), nitrogen, oxygen, and deionized (DI) water. As shown in Figure 7, residual stress of the films deposited at TD = 350 °C remained nearly unchanged over 24 hours of exposure, while films deposited at TD = 150°C changed to varying degrees depending on the environment. For all a-SiC:H films except those stored under nitrogen, a trend of increasing compressive stress is observed. The films stored under dry air and oxygen showed an increase of 17% and 25%, respectively. The film stored under DI water and air showed the highest increase in the compressive stress, as high as 32%, during the 24-hour study period.

Figure 7.

Figure 7.

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Residual stress in a-SiC:H films deposited at TD = 150 °C under exposure to various ambient environments over a period of 24 hours.

4. Discussion

As expected, the results presented show that residual stresses in as-deposited a-SiC:H films are strongly dependent on both deposition temperature and power. The results also show that the stress in a-SiC:H films deposited at low temperature (TD = 150oC) changes over time to an extent that is determined by the degree of exposure to oxygen or water. Films deposited at low temperature or low power density (TD ≤ 250 °C or < 0.2 W-cm−2 power density), as shown in Figure 1, increase in compressive stress on air exposure. Since the films for studying the effect of deposition temperature were all deposited at a high-power density of 0.27 Wcm−2, and the films for investigating the effect of power density were all deposited at 350 °C, the results indicate proper selection of both parameters is necessary for making a-SiC:H films with stable stress. Low-temperature films exhibit oxidation over time as observed from the XPS data in Figure 2. Since the XPS data were collected after 10 min of Ar-sputtering to remove surface layers, the data represent sub-surface bonding in the a-SiC:H. Similarly, the XPS data in Figure 2 show Si–O and O–Si peaks at Si 2p and O 1s regions indicating sub-surface oxidation in the low temperature deposited films. The peak intensities increase with time of air exposure (Figure 3).

From changes in FTIR and XPS spectra, oxidation of air-exposed low-temperature a-SiC:H films is ongoing over at least a 28-day period, and the oxidation corresponds to an increase in compressive residual film stress. The TD = 150 °C films also show an increase in Si-O-Si bonding and a decrease in Si-H bonding (inset of Figure 4(d)), suggesting that Si-H, from incomplete reaction of the SiH4 precursor, is being oxidized to Si-O. Since the C–C and C–Si–O peak intensities are unaffected with time (Figure 3(b)), oxidation does not appear to involve carbon. Since the FTIR transmission spectra in Figure 4, show notable intensity changes in the Si–O–Si and the O–H vibration modes at 1020 cm−1 and 3425 cm−1, respectively, it seems likely that the oxidation is occurring in the bulk of the film as well as at the surface of the TD =150oC films. While XPS in Figure 2 shows that the sub-surface of the low temperature films is predominantly composed of silicon-oxy-carbide at initial time points, FTIR reveals a strong Si–C peak (the highest intensity peak) whose band intensity remains relatively constant over time and is not affected by the growing peak at its shoulder representing the increasing Si–O–Si content (Figure 6). This result further suggests that the Si-C bonding in the TD=150oC is stable to air exposure, and the observed oxidation is predominately related to the high Si-H content of the low temperature films.

From Figure 7, residual stress in the low-temperature deposited films does not vary greatly when exposed to a N2 environment, suggesting that the residual stresses in a-SiC:H films will remain nearly unchanged in the absence of oxygen or water. In the presence of oxygen, as evident from the film stored in dry air (∼20% O2) and O2 ambient (100% O2), the rapid oxidation of the film results in a rapid increase in compressive stress; at a rate dependent on O2 concentration. A pronounced increase in compressive stress is also observed on air exposure, but less in dry air (Figure 7) implicating water vapor as well as oxygen gas in the oxidation of the low deposition temperature a-SiC:H. Finally, the largest oxidative effect and subsequent increase in compressive stress is observed for films soaked in deionized water. From these results, it is evident that both oxygen and water react with a-SiC:H to increase compressive film stresses. Similarly, reaction with oxygen and water vapor have also been shown to produce compressive stresses in other thin films such as aluminum [40] and silicon monoxide [41]. The environmental conditions investigated in this work reflect the oxygen, water vapor exposure, and liquid water the a-SiC:H films may encounter during thin-film processing and device fabrication. Residual stress changes in deionized water also suggest that stress changes may be driven by exposure to the physiological environment.

The increase in compressive stress in the low-temperature or low-power a-SiC:H films is therefore attributed to the oxidation of Si-H and the incorporation of larger oxygen or hydroxyl moieties into the a-SiC:H film. A consequence of the change in residual stress of the a-SiC:H is the loss of stress balance in thin-film structures. The lower density a-SiC:H may also not be as effective an encapsulation layer as the higher density films.

Our primary focus is the development of devices for chronic intracortical stimulation and recording. These multi-layer devices include both insulating and conductive thin films [4, 5], and will require careful balancing of film stress and film thickness to preserve a planar geometry [5]. Any deviation from planarity will likely exacerbate tissue damage and hemorrhage during insertion into target neural tissues or cause the device to deviate from its intended insertion path. Changes in a-SiC film stress after implantation may also be a concern since a stress change is expected to progressively induce curvature in the device resulting in potentially harmful tissue displacement. Recently, devices that actuate in vivo to produce a localized displacement of cortical recording electrodes have been investigated with the objective of optimizing the neuron-electrode interface [42]. Localized changes in residual stress may be one approach to inducing small localized displacements in implanted devices.

This study focused on the effects of deposition conditions on the magnitude and stability of residual stresses in a-SiC:H thin films. Films deposited under conditions that produced a stable residual stress were always observed to be under compression with a typical as-deposited stress of −80 MPa to −120 MPa (Fig. 1). This magnitude of residual stress is enough to induce curvature in a-SiC:H films, requiring a stress balancing strategy (5). Of interest for future work would be the development of deposition conditions that lead to a-SiC:H films with a more neutral residual stress that is also stable. In addition, the investigation of the effect of reactive oxygen species and other potentially damaging chemical constituents associated with the foreign body response to implanted devices should also be investigated [43].

5. Conclusions

The change in residual stress in PECVD a-SiC:H films exposed to air and wet ambient environments was investigated, and a relationship between film oxidation and stress changes identified. The increase in compressive residual stress is associated with oxidation of Si-H which is more prevalent in the films deposited at low temperature or at low power. To achieve stable low-stress a-SiC:H films, it is desirable to deposit these films at high temperature (≥325 °C) and power density (≥ 0.2 Wcm−2). These findings should serve as a guide for developing a-SiC:H based dielectric coatings for implantable medical devices where thin-film stress management and long-term stability are essential. While these results and earlier work[1] suggest an optimum deposition temperature of >300oC for PECVD a-SiC:H, alternative approaches to densifying PECVD films including the use of very high frequency (VHF) generators, multiple frequencies or pulsed plasma deposition could be investigated as a means of obtaining stable, low stress films at lower deposition temperatures.

Acknowledgements

The authors wish to thank the University of Texas at Dallas Natural Sciences and Engineering Research Laboratory cleanroom staff, especially Dr. Gordon Pollack and Mr. Scott R. Riekena, for their support in maintenance of the PECVD tool. We also want to thank Rebecca Frederick for assisting with the preparation of this manuscript. This work was conducted under NIH grant U01NS090454 awarded to Dr. Timothy J. Gardner of Boston University.

Footnotes

Conflict of interest: none declared

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