# Temperature Influence on the Deposition of Nitrogen-Doped Silicon Carbide Polycrystalline Films

Michail Gavalas, Scott Greenhorn, Frédéric Mercier, Konstantinos Zekentes

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Michail Gavalas, Scott Greenhorn, Frédéric Mercier, Konstantinos Zekentes. Temperature Influence on the Deposition of Nitrogen-Doped Silicon Carbide Polycrystalline Films. Coatings, 2025, 15 (1), pp.106. 10.3390/coatings15010106. [hal-05396544]

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Article

# Temperature Influence on the Deposition of Nitrogen-Doped Silicon Carbide Polycrystalline Films

Michail Gavalas<sup>1,2</sup>, Scott Greenhorn<sup>2,3,4</sup> (ID), Frédéric Mercier<sup>1,* ID</sup> and Konstantinos Zekentes<sup>2,3 ID</sup>

**Abstract:** Polycrystalline nitrogen-doped cubic silicon carbide (3C-SiC) thin films are grown on 2" Si wafers by a low-pressure chemical vapor deposition (LPCVD) technique with the aim for them to be used as support and active materials in microelectronic devices for neural interfaces. The effect of deposition temperature on the structural, mechanical, and electrical properties is investigated. The growth rate is varying, from 1 µm/h to 14 µm/h, along with the deposition temperature. We show that the structural and electrical properties of polycrystalline SiC are modified when changing the deposition temperature. Films with resistivity as low as (10.0 ± 0.5) mΩ·cm, a low residual stress of (–397 ± 158) MPa, and a low root mean square surface roughness of (53 ± 19) nm are achieved. Accelerated aging tests in heated phosphate buffer solution (PBS) show an etching rate of less than 1 nm/day and a steady low electrical resistivity for 77 days, indicating that the nitrogen-doped polycrystalline SiC is a chemically stable material, capable of chronic stability in a saline electrolyte.

**Keywords:** polycrystalline SiC; chemical vapor deposition; thin films; neural interfaces

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updates

Academic Editor: Mark Baker

Received: 25 November 2024

Revised: 10 January 2025

Accepted: 14 January 2025

Published: 18 January 2025

**Citation:** Gavalas, M.; Greenhorn, S.; Mercier, F.; Zekentes, K. Temperature Influence on the Deposition of Nitrogen-Doped Silicon Carbide Polycrystalline Films. *Coatings* **2025**, *15*, 106. https://doi.org/10.3390/coatings15010106

**Copyright:** © 2025 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https://creativecommons.org/licenses/by/4.0/).

## 1. Introduction

Polycrystalline SiC is one of many materials that can grow in the form of thin films by the thermal CVD process [1]. Its excellent thermal properties, radiation resistance, and chemical inertness make it appropriate for devices operating in extreme environments and thermal coatings, like automotive and aerospace parts [2], nuclear sensors [3], and chemically resistive coatings [4,5]. The low-resistive and low-stressed polycrystalline SiC is promising for high-quality MEMs fabrication, operating at high temperatures [6,7]. Its nontoxicity, along with its biocompatibility within a physiological environment, demonstrates its appropriateness for biosensing [1,8] and neural-interface applications [9]. For the latter application, and more specifically for an active (i.e., neural recording and/or stimulation) SiC neural-interface material, it is very important to have a combination of low electrical resistivity (<1 Ω·cm), low residual stress (<500 MPa), long stability in physiological solutions (many years), biocompatibility, as well as a low surface roughness in the range of 20 nm–50 nm, which is a trade-off between the requirements for an increased electrochemical signal and device processing limitations [9,10].

The purpose of our study is to determine the optimum deposition conditions to demonstrate the suitability of polycrystalline SiC as an active material for neural interfaces. LPCVD is used as the deposition technique for the polycrystalline SiC material.

<footnote n="1">¹ SIMAP, University Grenoble Alpes, CNRS, Grenoble INP, 38000 Grenoble, France; michail.gavalas@grenoble-inp.fr</footnote>

<footnote n="2">² Microelectronics Research Group, IESL/FORTH & University of Crete, 71003 Heraklion, Greece; scott-cameron.greenhorn@grenoble-inp.fr (S.G.); konstantinos.zekentes@grenoble-inp.fr (K.Z.)</footnote>

<footnote n="3">³ CROMA, University Grenoble Alpes, CNRS, Grenoble INP, 38000 Grenoble, France</footnote>

<footnote n="4">⁴ LMGP, University Grenoble Alpes, CNRS, Grenoble INP, 38016 Grenoble, France</footnote>

<footnote>* Correspondence: frederic.mercier@grenoble-inp.fr</footnote>
Thermal CVD of polycrystalline SiC is performed at low temperatures (800 °C–1500 °C) and pressures (traditionally some mbars–LPCVD) with a variety of precursors (organosilanes, chlorinated organosilanes, etc.) [11,12]. Both single-source precursors [13–18] and dual precursors are used [19–21]. In situ growth with NH₃(g) [22,23], N₂(g) [24–26], and PH₃(g) [18] results in n-type doping, with NH₃(g) the most frequently used. The deposition temperature effect on the nitrogen-doped polycrystalline SiC growth for both NH₃(g) [22] and N₂(g) [24,26] has already been studied; however, none of these reports investigated the effect of temperature on the structure, the mechanical stress, and the electrical properties of the polycrystalline SiC films simultaneously to meet the above-mentioned requirements for neural-interface active materials. In addition, the study of the chemical stability of the SiC material (see a review in [5]), and more specifically in buffer electrolytic solution (PBS: phosphate buffer saline), is limited to amorphous SiC [9,27,28] and single-crystalline SiC [29], while there are no studies yet on n-doped polycrystalline SiC material to our knowledge. The etching rate of amorphous SiC is generally in the order of some nm/day [9] (1.2 nm/day) and [27] (2.4 nm/day) in saline solutions, while single-crystalline SiC has shown a negligible dissolution rate [29]. Despite the fact that the monocrystalline SiC shows a negligible dissolution rate, the increased disorder of polycrystalline SiC may affect performance similarly to amorphous SiC.

In this work, polycrystalline SiC thin films are deposited with dual precursor (SiH₄ and C₃H₈) LPCVD, and the material's structure, mechanical stress, and electrical properties are investigated as a function of the deposition temperature with the aim to evaluate the deposited material's suitability as an active material for neural interfaces. Polycrystalline SiC nitrogen-doped films exhibiting simultaneously low residual stress, low electrical resistivity, and appropriate surface roughness are achieved. Finally, accelerated aging tests in heated saline electrolyte solution are performed to study the chemical reactivity of silicon carbide films.

## 2. Materials and Methods

### 2.1. Deposition Details

Polycrystalline SiC films were grown in an LPCVD vertical cold-wall reactor. The evacuation of gases was ensured by a rotary vane vacuum pump that allowed a primary vacuum inside the reactor (base pressure of 3·10⁻³ mbar). Silane (SiH₄) (3% in H₂(g) 99.9999%, supplied by Nippon Gases, Oevel, Belgium) was used as a Si source precursor, and propane (C₃H₈) (99.95%, supplied by Air Liquide, Chalon-sur-Saone, France) as a C source precursor. The SiH₄(g) and C₃H₈(g) flow rates were both set to 5 sccm, leading to a precursors ratio of (C/Si) (g):3. Hydrogen H₂(g) (99.9999%, supplied by Air Liquide) was used as a carrier gas with a flow rate of 400 sccm. Gaseous dichloride Cl₂(g) (99.999%, supplied by Air Liquide) was also injected in order to prevent the deposition of solid Si in the deposited films, fixed at 7.5 sccm, giving a fixed precursors ratio of (Cl/Si) (g):3. In situ n-doping was performed by injecting gaseous ammonia NH₃(g) (99.96%, supplied by Air Liquide) at 1 sccm, with a fixed precursors ratio of (N/Si) (g):0.2. For the correct electrical characterization, the films were grown on 1-micrometer-thick thermally grown SiO₂ on 2" Si (100) wafers provided by the Silicon Materials Company (Kaufering, Germany). The deposition time was set to 30 min for all investigated films. For each deposition temperature, at least four experiments are performed. Deposition temperature varied between 1050 °C and 1250 °C and the chamber total pressure during deposition was set to 20 mbar. The growth temperature was measured with a pyrometer, and the control of the heating power supply was performed manually. Note that the temperature range investigated here is similar to that investigated in [24,26], and higher than that used in [22]; however, these studies used different chemical systems (single-source 1,3-DBS(g) + NH₃(g) [22]; single-
source TMS ((CH₃)₃SiH)₃(g) + N₂(g) [24]; TCS ((Cl)3SiH)₃(g) or SiH₄(g) + C₂H₄(g) + N₂(g) [26]) to the one used in this study, which do not allow for direct comparisons. Therefore, the preliminary studies we performed showed that 1050 °C is the lower temperature required to obtain polycrystalline SiC, and below 1050 °C, nanocrystalline SiC is always obtained. Nevertheless, one sample deposited at 900 °C—with SiH₄(g), C₃H₈(g), Cl₂(g), and NH₃(g) flow rates set to 15 sccm, 10 sccm, 15 sccm, and 0.75 sccm, respectively, and fixed ratios (C/Si)₃(g):2, (Cl/Si)₃(g):2, and (N/Si)₃(g):0.25—is included in the discussion of the structural characterization and mechanical stress. Although this sample was not grown at the same precursor flow rates as the other samples, it can be indicative of the evolution of these properties at lower temperatures than the range studied.

## 2.2. Characterization Techniques

X-ray diffraction (XRD) was employed for the investigation of the crystallographic orientation and crystallite size of the deposited films. The measurements were performed in an asymmetric Bragg Brentano geometry using Cu Kα and a grazing incidence of 3°. One representative wafer for each temperature was measured by XRD. Two individual XRD measurements were performed per wafer; one in the center of the wafer and one at 2 cm from the center. The texture coefficient can be calculated by the following formula (Equation (1)) according to the model of Harris [30].

$$TC_{hkl} = \frac{\frac{I_{hkl}}{I_{0hkl}}}{\frac{1}{N}\left[\sum_{1}^{N} \frac{I_{hkl}}{I_{0hkl}}\right]} \quad (1)$$

where N is the number of peaks; and I$_{hkl}$ and I$_{0,hkl}$ are the intensity of the indices hkl of Bragg's reflection for the measured sample and a randomly oriented sample, respectively, as given in the ICDD file n° 00-029-1129. The crystallite size at each crystallographic orientation was estimated from the XRD measurements using the Scherrer equation [31]. The surface morphology was observed with scanning electron microscopy (SEM) for one film per deposition temperature. Optical profilometry was used for the measurement of the film's RMS surface roughness (with a scanned area of 173 µm × 173 µm) and also the curvature (with a scanned area of 3473 µm × 3473 µm) after the deposition, for all the deposited wafers. Five areas were measured per wafer, giving a mean value and a standard deviation of the RMS surface roughness and curvature. The residual stress was estimated from the wafer's curvature according to the model of Klein et al. [32] by considering in a first approximation the bilayer SiC/Si. The thermal stress between the film and the substrate was calculated using the thermal expansion coefficients (CTE) for cubic SiC [33], and Si [34] as a 5th-order polynomial function of the temperature. The mean values of residual stress per sample are displayed in the graphs since tensile and compressive stress can be found for a given temperature. The growth stress was calculated by subtracting the calculated thermal stress from the measured residual stress. The mean value for the film's thickness was determined by the weight difference of the sample before and after the deposition. The thickness variation, from the center to the edge of the wafer, was estimated with optical ellipsometry for two selected films; one grown at 1050 °C and one at 1100 °C. A Lorentz dispersion model with the parameters selected to match an optical index and dielectric constant of cubic SiC, n = 2.7 and k = 0.04 at 633 nm [35], was used for the optical analysis of the films. The electrical resistivity was measured by a 4-point probe at room temperature for all the deposited films. The mean value and the standard deviation of
resistivity per wafer is given for nine measured areas. The uniformity for each measured property is estimated using Equation (2).

$$ Uniformity \% = \frac{St.Deviation \%}{Average \ Value} \qquad (2) $$

Accelerated aging tests were performed by immersing the polycrystalline SiC in phosphate-buffered saline solution (PBS). Two n-doped polycrystalline SiC films, one grown at 1100 °C and one at 1250 °C, were studied. First, the samples were cleaned in acetone and rinsed with deionized water; then, a part of the sample was covered with polyimide tape as a reference. Then, the samples were placed in 200 mL of PBS solution of pH 7.4, heated at 90 °C, and removed after 14, 28, 43, and 77 days. A temperature of 90 °C was maintained through the experiment as much as possible—while the containers were sealed, there was still some evaporation, which led to crystallization (the PBS becoming cloudy) over the course of several weeks. To avoid this, the PBS was refreshed with new, freshly prepared solution every ~2 weeks in order to keep the concentration as consistent as possible. Typically, the fresh PBS was warmed for ~1 h at 90 °C to reduce the temperature change. During the solution changes, the samples were manipulated only by the protective PI tape, not by directly contacting the exposed sample surface. After removal from the PBS solution, the films were cleaned from PBS by an ultrasound bath of acetone, followed by immersion in isopropyl alcohol. The electrical resistivity of the uncovered part was measured by a 4-point probe at room temperature. Mechanical profilometry measured the step size between the covered region and the exposed region to estimate the etch rate after the final removal from PBS. Additionally, a Si substrate (280-micrometer-thick, n-doped, ρ = 2 mΩ·cm), supplied by the Siltronix Company, and a 1-micrometer-thick thermally grown SiO₂ on Si (100) wafer, from the Silicon Materials Company were also soaked for the aging tests to be compared with SiC since both materials were used previously in biomedical devices [36].

# 3. Results and Discussion

## 3.1. Deposition Rate

The deposition rate increases proportionally with the deposition temperature following the Arrhenius law (Figure 1), indicating that the growth is limited by the chemical kinetics. The growth rate value of the film grown at 900 °C has been removed since nucleation in the gas phase was observed, providing an artificially high value of the growth rate. Given the standard deviation of the growth rate values at each deposition temperature, the activation energy is estimated as (154 ± 16) kJ/mol for the deposition temperature range of 1050 °C–1250 °C. Similar values are also given for the growth of polycrystalline 3C-SiC on Si (100) (140 kJ/mol) [21] and graphite (108 kJ/mol) [16] substrates using the chemical system Si-C-Cl-H. Thickness was found to increase from the center to the edge of the wafer with an estimated thickness variation of 600 nm for a 1.5-micrometer-thick film, grown at 1100 °C, and 260 nm for a 680-nanometer-thick film, grown at 1050 °C. This is coherent with the fact that the wafer edges are hotter than the center of the wafer due to the heating system. The temperature difference is experimentally estimated to be 15 °C. For these films, the uniformity—estimated for a central wafer area of 2.7 cm and 1.2 cm diameter—is, respectively, 13% and 5% for the 1100 °C film and 10% and 5% for the 1050 °C film.
<figure><img src="image_3.png" /><figcaption>**Figure 1.** Deposition rate vs. reciprocal deposition temperature for n-doped polycrystalline SiC thin films. The graph displays the mean value and standard deviation of a series of experiments for each temperature.</figcaption></figure>

## 3.2. *Structural Characterization*

The comparison of the XRD diffractograms for the n-doped polycrystalline SiC films grown at different temperatures is shown in Figure 2. The broad peak observed for all samples between 65° and 75° is the signature of Si (004) (69.1°) coming from the substrate due to the asymmetric geometry of the XRD measurement (tilt of 3°). Note that this peak is more pronounced in Figure 2 for the lower temperatures due to the appearance of the peak at 71.8° at higher temperatures. Films deposited at 900 °C show only a broad peak from 25° to 40°. This can be attributed to the Si (111) orientation (peak position of 28.4°) or to the 3C-SiC (111) orientation (peak position of 35.6°) but certainly not from the (100) Si substrate. Nanocrystalline materials, such as Si(\_s) or SiC(\_s) or Si(\_s) + SiC(\_s), could be responsible for the appearance of this peak. From 1050 °C to 1250 °C we observe three SiC peaks: (111) at 35.6°; (220) at 60.1°; and (311) at 71.8°. The (311) and (200) (at 42°) peaks are clearly visible above 1200 °C and are possibly present for lower temperatures with less intensity. A SiC peak (222) at 76.1° is observed only for the sample grown at 1200 °C. The sample grown at 1250 °C shows an additional intense peak at 33.7°. A shoulder is also visible for the sample grown at 1200 °C at the same position. This peak could result either from the presence of hexagonal SiC (ICDD n° 00-029-1127), according to Liu et al. [37], or from cubic SiC with its high density of stacking faults, according to Gallou et al. [13]. Additional weak peaks at 28.4° and 47.6°, however, exist for all the samples except the one grown at 1100 °C. These peaks are linked to the presence of a small contribution of silicon clusters in the deposited films.

In order to obtain a better understanding of the evolution of the crystalline structure with the deposition temperature, the texture coefficient has been calculated for all samples, except 900 °C, as shown in Figure 3. The deposition temperature has a strong impact on the texture of the grown polycrystalline material. Lower temperatures produce films with (111) preferential orientation; however, a change in the preferential orientation from (111) to (220) mainly—but also to (311) and (200) at a lesser degree—is observed for films deposited at 1250 °C. Such changes in preferential orientation with the increased temperatures are also observed in previous reports for chlorinated chemistries of LPCVD [13,16,38] and could be attributed to a twin-mediated growth mechanism according to [13]. Indeed, all these studies have shown the existence of twin planes of <110> orientation [13,16,38]. This mechanism is activated for a decrease in the supersaturation and an increase in the
deposition temperature due to the enhancement of the adatoms' surface mobility and due
to the lower nucleation frequency [13].

<figure><img src="image_4.png" /><figcaption>**Figure 2.** Comparison of XRD diffractograms of n-doped polycrystalline SiC grown at 900 °C (black), 1050 °C (red), 1100 °C (blue), 1200 °C (green), 1250 °C (purple). The diffractograms are stacked on log10 scale.</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>**Figure 3.** Texture coefficient of (111)—black, (220)—green, (311)—red, and (200)—orange XRD peaks of the n-doped polycrystalline SiC vs. deposition temperature.</figcaption></figure>

The evolution of the crystallite size for the (111) orientation of the n-doped poly-
crystalline SiC is also investigated. First, it must be mentioned that the crystallite size
variation between the center and 2 cm off-center remains lower than 10%, independent
of the deposition temperature, although the crystallite size was found to increase toward
the wafer's edge. This can be explained by the highest temperature being at the edges
of the wafer, as discussed previously for the thickness variation. It appears (see Figure 4)
that the increase in the deposition temperature results in a larger crystallite size due to the
higher mobility of the adsorbed species, probably leading to a larger diffusion distance,
thus improving the crystallinity of the deposited film [16]. Since our deposition surface is
amorphous SiO2, the nucleation process will depend only on Gibbs energy and not on a
possible low-energy crystalline site [39]. Thus, at higher deposition temperatures, there is a
decrease in the saturation density of nuclei and an increase in the critical size of the grown
nucleus [40], which provides more space for islands to grow. In addition, increasing the
temperature could increase the mobility of the adsorbed species, resulting in higher surface
diffusion distances, which is in agreement with previous studies [16,21,37,38,41,42]. Above 1100°, we observe saturation and even a slight decrease in the case of 1250 °C. This can be due to the change in the preferential orientation of the crystallites (see Figures 2 and 3) or from the appearance of other peaks close to that of (111) [43]. The peak at 33.7°, attributed either to hexagonal SiC or stacking faults, is very close to that of (111).

<figure><img src="image_6.png" /><figcaption>**Figure 4.** Evolution of the average crystallite size for the (111) crystallographic orientation with the deposition temperature. The sample grown at 900 °C is indicated with a square.</figcaption></figure>

### 3.3. *Surface Morphology and Roughness*

For the deposition temperatures 1050 °C to 1100 °C, spherical-shaped grains are observed (Figure 5) on the films' surfaces. A similar morphology of spherical-shaped agglomerates is observed for the sample at 900 °C. These agglomerates are attributed to a nucleation in the gas phase close to the substrate surface. Faceted SiC grains are observed for deposition temperatures higher than 1100 °C. Our observation shows that the grain size is increasing from 1050 °C to 1250 °C. Similar results, in terms of surface morphology, have been found in previous reports for polycrystalline SiC and polycrystalline Si films [16,20,37,41,42,44]. The increase in grain size is also in agreement with the corresponding increase in the crystallite size given in our XRD results.

<figure><img src="image_7.png" /><figcaption>**Figure 5.** *Cont.*</figcaption></figure>
<figure><img src="image_8.png" /><figcaption>**Figure 5.** SEM images: surface morphology of n-doped polycrystalline SiC grown at (a) 1050 °C, (b) 1100 °C, (c) 1200 °C, (d) 1250 °C, and (e) 900 °C.</figcaption></figure>

The RMS surface roughness, determined by optical profilometry, was found to increase proportionally from 1050 °C to 1250 °C (Figure 6). The increased roughness is related to an increased grain size (Figure 5) and crystallite size (Figure 4), in agreement with [19,24,38,42]. Therefore, the growth conditions favoring large grains also favor high surface roughness. The increased RMS roughness measured for the films grown at 900 °C is probably due to the presence of agglomerates (see Figure 5e) rather than the actual roughness itself. According to the initial requirements, deposition temperatures of 1050 °C–1100 °C provide appropriate surfaces for MEMs and neural-interface fabrication, respectively. Surface roughness uniformity for 900 °C to 1200 °C was found to be 10%–40%, while for films grown at 1250 °C, the roughness shows a lower uniformity of 3%–5%. No clear evolution has been observed between the centers and the edges of the wafers.

### 3.4. Residual Stress

Figure 7 shows the evolution of residual stress according to the deposition temperature. For most experiments, it was found that the films have stress lower than 500 MPa in absolute value. The films have tensile stress, except those grown at 1100 °C, for which the stress is compressive. We observe that between 1050 °C and 1250 °C deposition temperatures, there is no clear trend between the residual stress and the deposition temperature. Residual stress seems to have a steady uniformity range (10%–40%) per wafer independent of the deposition temperature. No clear evolution was observed between the centers and the edges of the wafers.
<figure><img src="image_9.png" /><figcaption>**Figure 6.** RMS surface roughness values per wafer (red points) of n-doped polycrystalline SiC films with deposition temperature. Dotted line represents the evolution of the average RMS roughness with temperature for all the wafers (blue points). The sample grown at 900 °C is indicated with a blue square.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>**Figure 7.** (a) Residual stress per wafer (black points) and their average value per temperature (red points); and (b) calculated thermal and growth stress of n-doped polycrystalline SiC films with deposition temperature, grown on SiO₂/Si (100). The sample grown at 900 °C is indicated with a red square in both (a,b).</figcaption></figure>

To further investigate the effect of deposition temperature on the stress generated in the polycrystalline SiC films, we first calculated the thermal stress and then extracted the growth stress (Figure 7) [32]. The calculated thermal stress is compressive for all deposition temperatures since the difference between the two thermal expansion coefficients of SiC and Si is negative ($CTE_{(SiC)} - CTE_{(Si)} < 0$). As long as the thermal stress has a monotonic behavior for our system, the evolution of the growth stress is the reason behind the corresponding evolution of the residual stress with the deposition temperature. The growth stress, as calculated by subtracting the calculated thermal stress from the measured residual stress, was found to be tensile and more or less constant, except for 1100 °C, where the minimum growth stress value was observed.

In polycrystalline materials, the major mechanisms that influence the growth stress are as follows [45–47]: (1) cluster coalescence; (2) incorporation of impurities either in the grain or in the grain boundaries. The cluster coalescence mechanism induces tensile stress due to the surface free energy between the crystallites during coalescence [45,46]. The stress generation, due to the incorporation of impurities, is a more complicated process.
Inside the grains, nitrogen atoms are likely replacing C atoms in the “C sites” within the crystallite lattice during in situ doping, according to the site competition theory [48]. Theoretical [49,50] and experimental [6,15,23] studies show increasing tensile stress with the addition of nitrogen. The proposed physical reason is that the newly Si-N covalent bonds—which have a smaller covalent radius than the original Si-C bonds—cause a shrinkage in the lattice, resulting in an increase in the tensile stress. On the contrary, the incorporation of impurities at the grain boundaries could also induce compressive stress in non-energetic deposition conditions like CVD processes [46]. In our work, it is possible that the growth stress is affected more by the crystallite size for lower temperatures, assuming that the amount of nitrogen impurities is not high enough. Thus, the decrease in growth stress from 900 °C to 1100 °C is due to an enlarged crystallite size. On the contrary, for higher temperatures, the amount of nitrogen can be increased significantly, thus increasing the growth stress, which could explain the observed increase for temperatures higher than 1100 °C. Indeed, it is reported that increasing the deposition temperature leads to higher nitrogen incorporation for the 4H-SiC single-crystalline [51] and quite probably inside the grains of the polycrystalline film. The effect of incorporating nitrogen at the grain boundaries and the corresponding effect of the compressive stress generation is not observable in the present study. Thus, it would be useful to identify the average nitrogen concentration N<sub>d</sub> of the polycrystalline SiC films, both in the grains and at the grain boundaries for each deposition temperature, in order to investigate the reasons for the temperature-dependent growth stress in the future.

### 3.5. *Electrical Resistivity and Free-Carrier Density*

Figure 8 shows the evolution of the electrical resistivity for the n-doped polycrystalline SiC wafers with the deposition temperature. In our study, the lowest value achieved is 8 mΩ·cm at 1200 °C, even lower than the previously reported values [15,22,24,26]. We propose that the optimum deposition temperature in terms of reproducibility and resistivity values is 1250 °C, where there is a combination of low resistivity values and small values for dispersion. One can say that higher temperatures result in lower resistivity. Similar behavior has been reported previously [22,24,26] (Figure 8b). The authors propose that the decrease in resistivity is the result of an improved crystallinity of their films at higher deposition temperatures but without further explanation.

<figure><img src="image_11.png" /><figcaption>**Figure 8.** Electrical resistivity of n-doped polycrystalline SiC films and deposition temperature: (a) average electrical resistivity of all wafers vs. temperature; and (b) comparison of our individual values with the available state of the art (References on graph (b): Cheng et al., 2002 [24], Privitera et al., 2015 [26]).</figcaption></figure>
Various deposition experiments have been performed without introducing NH₃(g) into the chemical system, and the resulting resistivity was several orders of magnitude higher (10–10,000 Ω·cm); this indicates that the decrease in resistivity when introducing NH₃ is due to nitrogen atoms’ incorporation in acting as n-type dopants. Early results from Hall measurements on the films grown at 1100 °C showed an average carrier’s concentration of 5·10¹⁸ cm⁻³. Our values are similar to those of previous reports of n-doped polycrystalline SiC (~6·10¹⁸ cm⁻³ [23]). The average measured mobility for our films was 11.55 cm²/V·s.

Higher deposition temperatures resulted in doped polycrystalline films with a larger crystallite size and, therefore, fewer trap states located in the grain boundaries. This effect results in a higher amount of free carriers and, thus, in lower resistivity, according to the models of the dopant segregation in the grain boundaries [52] and the free carriers trapping [52,53]. In addition, as discussed earlier, higher amounts of incorporated nitrogen at higher deposition temperatures [51,54] allow for even more free carriers to be available for conductivity.

Finally, we found that electrical resistivity uniformity decreases with deposition temperature. Especially for deposition temperatures 1200 °C and 1250 °C, the resistivity has a uniformity of less than 10%. No significant evolution between the center and the edges was observed.

### 3.6. *Chemical Stability*

Mechanical profilometry measurements show no observable step between the reference area covered with polyimide tape during immersion and the uncovered polycrystalline SiC in all of the sequential measurements up to 77 days. Indeed, after removal of the polyimide tape and profilometry measurement, there is no indication of etching steps, showing a negligible etching rate in PBS solution at 90 °C. Similar behavior is observed for the sample deposited at 1100 °C and the sample deposited at 1250 °C. The same experiments conducted with Si and thermally grown SiO₂ exhibit much higher etching rates of 104 nm/day and 26 nm/day (Figure 9), respectively, in agreement with Phan et al. [29]. Therefore, the chemical stability of polycrystalline SiC is better than amorphous SiC and performs as well as epitaxial SiC in saline solutions; this indicates a denser structure for the polycrystalline SiC compared with its amorphous counterpart [28]. In addition, the grain boundaries inherent in polycrystalline SiC do not appear to affect its chemical stability in saline solutions.

<figure><img src="image_12.png" /><figcaption>**Figure 9.** (a) Etching rate for polycrystalline SiC, Si, and thermally grown SiO₂ obtained after 77 days of accelerated aging tests in heated 90 °C PBS solution. The data for amorphous SiC are taken from</figcaption></figure>
Cogan et al. [27] (154 days in heated 90 °C PBS solution, SiC film thickness of 1 µm). The data for epitaxial SiC are taken from Phan et al. [29] (8 days in heated 96 °C PBS solution with a SiC film thickness of 230 nm); (b) room-temperature electrical resistivity of n-doped polycrystalline SiC films grown at 1100 °C (blue) and 1250 °C (red) in comparison with the Si substrate (green).

The electrical resistivity of n-doped polycrystalline SiC is stable for a period of 77 days of accelerated aging tests for both films (Figure 9b). Since no damage was detected by the etching process, no broken chemical bonds or changes in the chemical-crystalline structure are expected and, thus, no changes in the electrical properties of the material. Therefore, our polycrystalline SiC is appropriate as an active material for neural-interface devices. In contrast, the n-doped Si sample showed an increase in the electrical resistivity from 2 mΩ·cm to 10⁷ mΩ·cm (14 days of aging) and 5·10⁶ mΩ·cm (28 days of aging). The intense etching results in a change in its structure and eventually in its electrical properties. Thus, it is expected that Si-based neural devices are not appropriate for long-term use.

## 4. Conclusions

Nitrogen-doped polycrystalline SiC films have been demonstrated with low electrical resistivity $ρ < 1 Ω·cm$, low residual stress $σ ≤ 500 MPa$, and a surface roughness of less than 100 nm, deposited at 1100 °C–1200 °C; these are appropriate as both an active material and support material for MEMs and neural-interface devices since they combine all of the required properties. These SiC films are chemically and electrically stable for 77 days in a heated saline solution, which makes them appropriate for long-term biomedical devices. Future work can include an investigation of the influence of crystallite and grain size on the chemical stability for potentially over 100 days. Additionally, live electrical measurements of resistivity while immersed in heated PBS at 90 °C could also be studied.

The deposition temperature plays a crucial role in the structural quality, mechanical stress, and electrical resistivity of the deposited films. It was found that the growth rate, crystallinity, and surface roughness are proportional to the deposition temperature. On the contrary, electrical resistivity was found to be inversely proportional to it. The deposition temperature has a complex effect on the evolution of growth stress, attributed to the cluster coalescence and impurity incorporation mechanisms.

Despite the fact that the deposition temperature impacts the growth stress and the electrical resistivity of the polycrystalline film, a solid explanation for the temperature dependence on nitrogen incorporation—along with its effect on the growth stress and electrical resistivity of polycrystalline SiC—is yet to be given and needs further study.

**Author Contributions:** Conceptualization, M.G., F.M. and K.Z.; methodology, M.G., S.G., F.M. and K.Z.; validation, F.M. and K.Z.; investigation, M.G., F.M. and K.Z.; data curation, M.G., F.M. and K.Z.; writing—original draft preparation, M.G.; writing—review and editing, M.G., S.G., F.M. and K.Z.; visualization, M.G.; supervision, F.M. and K.Z.; project administration, M.G., F.M. and K.Z.; funding acquisition, F.M. and K.Z. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the National French Research Agency under the project SiCNeural (ANR Grant ID: ANR-21-CE19-0054-01) and a Horizon Europe research and innovation program NANOMAT with the grant agreement no. 101091433.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Data are contained within the article.
**Acknowledgments:** The authors thank Thierry Encinas for the XRD measurements, Alexandre Crisci for his help on the ellipsometry and profilometry measurements, and Katerina Tsagkaraki for her help in the scanning electron microscopy observations.

**Conflicts of Interest:** The authors declare no conflicts of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript; or in the decision to publish the results.

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