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Article

Electrochemical Performance and Cytocompatibility of HVOF-Sprayed Cr3C2-20(Ni20Cr)-20HAp-XSi Coatings for Dental Applications

by
John Henao
1,
Oscar Sotelo-Mazon
2,*,
Rosa M. Montiel-Ruiz
3,
Carlos A. Poblano-Salas
4,
Diego G. Espinosa-Arbelaez
5,
Jorge Corona-Castuera
4,
Astrid Giraldo-Betancur
6,
Ana L. Islas-Garduño
3 and
Victor M. Zezatti
7
1
SECIHTI-CIATEQ A.C., Parque Industrial Bernardo Quintana, El Marqués, Querétaro 76246, Mexico
2
ICF-UNAM, Av. Universidad s/n, Col. Chamilpa, Cuernavaca 62210, Morelos, Mexico
3
Centro de Investigación Biomédica del Sur. CIBIS-IMSS, Calle Argentina 1, Xochitepec 62790, Morelos, Mexico
4
CIATEQ A.C., Parque Industrial Bernardo Quintana, El Marqués, Querétaro 76246, Mexico
5
CIDESI, Av. Pie de la Cuesta 702-No. 702, Desarrollo San Pablo, Querétaro 76125, Mexico
6
Cinvestav, Unidad Querétaro, Lib. Norponiente 2000, Juriquilla, Querétaro 76230, Mexico
7
Centro de Investigación en Ingeniería y Ciencias Aplicadas, CIICAp, Universidad Autónoma del Estado de Morelos, Cuernavaca 62209, Morelos, Mexico
*
Author to whom correspondence should be addressed.
Appl. Sci. 2025, 15(17), 9308; https://doi.org/10.3390/app15179308
Submission received: 2 August 2025 / Revised: 21 August 2025 / Accepted: 21 August 2025 / Published: 25 August 2025
(This article belongs to the Special Issue Surface Coatings: Materials and Techniques)
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Abstract

Biocompatible coatings are widely employed in dental applications to enhance the biofunctionality of metallic implants exposed to the aggressive oral environment. Among them, hydroxyapatite (HAp)-based and carbide-reinforced coatings have been explored due to their favorable mechanical and biological performance. In this study, Cr3C-20(Ni20Cr)-20HAp-XSi coatings were deposited using the high-velocity oxy-fuel (HVOF) technique. The coatings were applied onto commercially pure titanium substrates, with the silicon content varied between X = 0, 5, 10, and 20 wt%. To evaluate the coatings’ corrosion resistance, electrochemical techniques such as potentiodynamic polarization curves, linear polarization resistance (LPR), electrochemical impedance spectroscopy (EIS), and open circuit potential (OCP) were employed. Artificial saliva was used as the corrosive medium at 37 °C for 168 h. The feasibility of producing carbide-HAp-Si coatings with excellent corrosion resistance and cytocompatibility via HVOF was demonstrated here, although some of the tested coatings (20 wt% Si) showed reduced electrochemical stability, attributed to faster dissolution processes and associated with a thinner coating layer, as confirmed by SEM analyses. X-ray diffraction (XRD) analyses revealed the formation of new phases in the coatings during thermal spraying, including Cr2O3 and Cr7C3. Additionally, MTT assays using 3T3-L1 fibroblasts showed no significant cytotoxic effects after 24 and 72 h of exposure to some of the coatings, confirming their biocompatibility for potential dental applications.

1. Introduction

Dental implants are biomedical devices designed to be anchored into the jawbone, serving as a stable foundation for prosthetic components in the rehabilitation of missing teeth. Unlike other internal implants, they maintain direct exposure to the oral cavity, an environment characterized by constant microbial presence and mechanical challenges. This open interface increases the risk of complications such as corrosion, infection, and implant failure [1]. Therefore, ensuring both mechanical durability and biological compatibility of the implant is critical. The development of functional surface coatings plays a key role in enhancing implant performance under such complex conditions [2].
Alloys used in oral applications are typically divided into two categories: precious metal systems, such as gold-based alloys with platinum, silver, or palladium; and non-precious alloys, including titanium and its alloys, cobalt–chromium (Co-Cr), and nickel–chromium (Ni-Cr) alloys [3]. Due to cost and processing advantages, non-precious alloys are more frequently used in clinical practice [4]. Recent advancements in dental implantology have focused on enhancing the interaction between these alloys and bone tissue to promote efficient osseointegration [5]. Several metallic alloys have proven essential in this field, including titanium-based alloys such as commercially pure Ti (Grade 2) and Ti-6Al-4V, 316L stainless steel, nickel–titanium (NiTi) shape-memory alloys, and NiCr alloys (VeraBond, KeraN, VeraSoft) [6].
A comparatively lower corrosion resistance and the potential release of metal ions raise concerns regarding local biological responses of these alloys in dental environments, including allergic or inflammatory reactions [7]. To address these limitations, surface modification strategies have been developed to enhance corrosion resistance and improve biocompatibility. Such strategies include the application of ceramic-based coatings, such as oxides, nitrides, and carbides, deposited via techniques like physical vapor deposition (PVD), chemical vapor deposition (CVD), sol–gel, and thermal spray [8,9,10].
Hydroxyapatite (HAp) coatings are currently widely used to enhance the performance of metallic implants in both dental and orthopedic applications. Due to its chemical and structural similarity to the mineral phase of human bone, HAp exhibits bioactive and osteoconductive properties that promote strong interfacial bonding between the implant and the surrounding bone tissue [11]. This bonding not only improves biological fixation but also supports early-stage osseointegration, thereby reducing the risk of implant failure and promoting long-term stability [12]. A coating containing HAp may enhance the biological response while simultaneously acting as a physical and chemical barrier against corrosion [13], especially in the aggressive and variable conditions of the oral cavity.
In recent years, carbide coatings have emerged as promising materials for dental implant surface modification due to their exceptional combination of mechanical and biological properties [14,15,16]. For example, SiC-based coatings are characterized by their high hardness, wear resistance, low density, and excellent corrosion stability, even under physiological conditions. SiC coatings have shown favorable biocompatibility in various studies [17,18,19]. Experiments involving human fibroblasts, osteoblasts, and neural stem cells have demonstrated that SiC coatings are cytocompatible and in some cases, even more biocompatible than elemental silicon [17,18,20]. On the other hand, coatings based on chromium carbide (Cr3C2), particularly in combination with NiCr matrices, have demonstrated significant potential in improving the superficial mechanical properties of titanium-based implants, such as those made of Ti6Al4V [21,22,23,24]. Using thermal spray methods such as detonation spraying, Cr3C2-NiCr coatings have been successfully deposited to enhance Ti6Al4V substrate hardness, reduce wear rate, and modify surface topography [24]. These coatings have been proven to increase substrate hardness from ~28 HRC to over 34 HRC, while simultaneously decreasing wear rate by nearly 85%, which is critical in the abrasive environment of the oral cavity. Additionally, the increased surface roughness observed after coating application can be advantageous for cellular adhesion and early-stage osseointegration [25].
Recent advances in surface engineering for dental implants have explored the integration of silicon as a functional dopant in various coating systems due to its demonstrated capacity to enhance both bioactivity and electrochemical stability [26,27,28]. Silicon-doped coatings, such as Si–TiO2 layers produced via microarc oxidation on titanium substrates [27], have shown improved corrosion and wear resistance, alongside enhanced cellular adhesion and proliferation, making them particularly relevant for dental applications where implant stability and biological integration are critical. Similarly, amorphous carbon coatings doped with silicon and fluorine (a-C:H:Si:F) have been applied on NiTi alloys [28], offering strong interfacial adhesion and favorable cytocompatibility, while simplifying the deposition process by eliminating the need for multilayered systems. In the context of bioactive ceramics, the incorporation of silicon into hydroxyapatite (Si-HA) coatings has been shown to modulate the coating solubility, promote ionic substitution within the apatite lattice, and influence early-stage bone response [26]. These findings have motivated the investigation of hybrid or doped systems based on silicon-modified carbides, particularly silicon-doped carbides and composite coatings combining carbides (e.g., Cr3C2 or SiC) with bioactive or silicon-containing phases [23,29,30]. Such approaches aim to synergize the exceptional mechanical properties and wear resistance of carbide-based coatings with the biological advantages induced by silicon, offering a promising pathway toward multifunctional surfaces tailored for high-performance dental implants.
Despite the extensive research on individual coating compositions such as Cr3C2-NiCr, hydroxyapatite, and silicon-doped systems, the synergistic integration of these materials using the HVOF technique remains largely underexplored in the specific context of dental implants, where multifunctional coatings are especially needed. Thus, a critical gap exists in the development of unified, hybrid coatings that combine high hardness, wear resistance, bioactivity, and corrosion stability in a single layer tailored for the oral environment. In this context, the present study employs the high-velocity oxy-fuel (HVOF) technique for the fabrication of Cr3C2-NiCr-HAp-Si composite coatings. HVOF is a thermal spray process that uses the combustion of fuel and oxygen to generate a high-velocity gas stream capable of accelerating powder particles to supersonic speeds before impact on the substrate. This method is known for producing dense, well-adhered coatings with low oxide content and controlled crystallinity, key factors for optimizing both mechanical and biological performance [31]. This approach is motivated by previous studies [12,13,21,22,23,24,26,27,28] demonstrating the individual benefits of each component: the mechanical strength and wear resistance of carbides, the corrosion resistance of NiCr alloys, the bioactivity of hydroxyapatite, and the enhanced osseointegration potential associated with silicon incorporation. By combining these materials through HVOF processing, the aim is to achieve synergistic properties suitable for dental implant applications, where durability, biocompatibility, and chemical stability are essential.

2. Experimental Procedure

2.1. Powder and Coatings Preparation

Commercial Cr3C2-20(Ni20Cr) cermet (Diamalloy 3007, Oerlikon Metco, Westbury, NY, USA), hydroxyapatite (HAp) (Captal 30 SD, Biotal, Buxton, UK), and silicon (99% purity, Sigma-Aldrich, Darmstadt, Germany) precursor powders were used in this study. The powder mixtures were prepared using a roller mill operating at 300 rpm and 25 °C. The mixing was conducted in dry conditions under ambient atmosphere and without using milling balls. Four different powder mixtures were formulated on a weight percentage basis as follows: (1) Cr3C2-20%HAp (wt%), (2) Cr3C2-20%HAp-5%Si (wt%), (3) Cr3C2-20%HAp-10%Si (wt%), and (4) Cr3C2-20%HAp-20%Si (wt%). Prior to mixing, all powders were dried at 90 °C for 4 h to eliminate moisture. After mixing, the powders were stored in sealed containers under dry conditions until further use for HVOF spraying. Pure titanium (grade 2) was used as the substrate material for coating deposition. Titanium coupons measuring 10 × 10 × 6 mm were cut and ground using up to 1200 grit SiC abrasive paper. Subsequently, the surfaces were grit-blasted with Al2O3 particles (ANSI G-20 grade) to achieve an average roughness of approximately 5 µm. The substrates were then cleaned with ethanol prior to the spraying process.
A high-velocity oxy-fuel (HVOF) system (DJ2700, Oerlikon Metco) was employed to deposit the coatings in this study. Oxygen, air, and propane were used as oxidant and fuel gases, with flow rates of 1446 SLPM, 1785 SLPM, and 213 SLPM, respectively. Nitrogen was employed as the carrier gas at a flow rate of 2154 SLPM. The powder feed rate was fixed at 12 g/min, whereas the spray distance was maintained at 230 mm. Coating deposition was performed using a 6-axis robotic arm (KUKA, K15, Augsbourg, Germany) to control the HVOF gun movement on the sprayed surface. A list of the coatings obtained is presented in Table 1.

2.2. Structural and Electrochemical Evaluation

Scanning Electron Microscopy (SEM) and X-ray Diffraction (XRD) analyses were performed using a JEOL JSM-IT500 microscope (Jeol, Tokyo, Japan) operating at 20 kV and a Bruker D8 Advance ECO diffractometer (Bruker, Billerica, MA, USA), respectively. SEM analysis, using backscattered electron (BSE) radiation, was employed to examine the surface of the coatings. XRD was performed with CuKα radiation (λ = 1.5406 Å), operating at 30 kV and 20 mA. The XRD measurements were performed over a 20–60° 2θ range, with a step size of 0.02° and an integration time of 1 s. XRD analysis was performed to identify the phases present in the coatings.
Electrochemical measurements were carried out using a Gamry Interface 1000 potentiostat/galvanostat. A conventional three-electrode electrochemical cell was employed, in which the coating served as the working electrode, whereas a saturated calomel electrode (SCE) and a graphite rod served as the reference and counter electrode, respectively. The working electrodes were prepared by welding the coated specimens to Ni20Cr wires, which acted as electrical connectors to the potentiostat/galvanostat. To prevent contact between the corrosive solution and the Ni20Cr wire, the working electrodes were encapsulated in glass tubes and sealed with epoxy resin. Artificial saliva, maintained at 37 °C to simulate oral conditions, was employed as the corrosive medium. The precursors employed to prepare artificial saliva are presented in Table 2.
Electrochemical characterization included open circuit potential (OCP), potentiodynamic polarization (PDP), linear polarization resistance (LPR), and electrochemical impedance spectroscopy (EIS) measurements. Potentiodynamic polarization curves were recorded within a potential range of −600 to +1600 mV relative to the corrosion potential (Ecorr), using a scan rate of 1 mV/s. LPR measurements were performed in a ±10 mV potential range relative to the OCP, at a scan rate of 10 mV/min. For the EIS measurements, an amplitude of ±10 mV was applied over a frequency range of 104 Hz to 0.01 Hz. The OCP, LPR, and EIS tests were conducted continuously throughout 168 h to evaluate the electrochemical behavior of the coatings [32,33].

2.3. Cytocompatiblity Study

2.3.1. Cell Line and Culture Conditions

A 3T3-L1 fibroblast cell line (ATCC® CL-173, In Vitro S.A., Ciudad de México, Mexico) was used for cell viability assays. Cells were cultured in Dulbecco’s Modified Eagle Medium (DMEM) supplemented with 10% fetal bovine serum (FBS), L-glutamine (1 mLx100), 3X antibiotic (1 mLx100), non-essential amino acids (1 mLx100), and sodium bicarbonate (3 mLx100). The cell cultures were incubated at 37 °C in a 5% CO2 atmosphere.

2.3.2. Cell Viability Assay

3-(4,5-dimethylthiazol-2-yl)-2,5-diphenyltetrazolium bromide (MTT) was used to determine relative cell viability [34]. The 3T3-L1 fibroblasts were seeded in a 96-well plate (3 × 104 cells/well) in supplemented DMEM culture medium and incubated at 37 °C with 5% CO2 until 80% confluence was obtained. The culture medium was then removed, and the cells were treated with different concentrations (3–1000 μg/mL) of the materials, evaluated in a 1% dimethyl sulfoxide (DMSO)-DMEM solution, and left to incubate for 24 and 72 h. Then, MTT (20 μL, 5 mg/mL) and DMEM (80 μL) were added to each well, and the mixture was incubated at 37 °C for 4 h. The culture medium was removed and replaced with isopropanol (HCl 0.04 N), followed by mechanical agitation for 15 min. Finally, the optical density (OD) was determined at 570 nm in a BioTek plate reader. The results are shown as the percentage of cell viability compared to that of the untreated group (% of control).

3. Results and Discussion

Figure 1 presents the XRD patterns of the coatings in the as-sprayed condition. The diffraction peaks indicate the presence of carbide phases, primarily Cr3C2 and Cr7C3, along with a Ni-based solid solution (Ni-Cr) in the C1 coating. These findings are consistent with previously reported results for Cr3C2-20(Ni20Cr) coatings deposited by high-velocity oxy-fuel (HVOF) spraying [35]. The presence of Cr7C3 is noteworthy, as it can be formed because of partial decomposition of Cr3C2 during the spraying process. This phase transformation is attributed to the high thermal energy input, which induces carbide destabilization and reformation of secondary phases [35].
A minor presence of Cr2O3 was also identified in the XRD patterns, although the intensity of its diffraction peaks was low, suggesting that only small amounts of this phase were formed. The formation of Cr2O3 is likely due to the oxidation of chromium during the thermal spraying process, caused by its exposure to oxygen from both the combustion flame and the surrounding atmosphere [36].
These phase evolutions highlight the complex thermochemical interactions occurring during the HVOF process. Several reactions have been proposed in the literature to describe the decomposition of Cr3C2 and the formation of Cr7C3 and Cr2O3 phases during spraying (see Equations (1)–(4)) [35]:
7/5Cr3C2 → 3/5Cr7C3 + C
Cr3C2 + 5/14O2 → 3/7Cr7C3 + 5/7CO
2Cr3C2 + 13/2O2 → 3Cr2O3 + 4 CO
2Cr7C3 + 27/2O2 → 7 Cr2O3 + 6 CO
Figure 1 also shows the XRD patterns of the C2 composite coating formed by the mixture of Cr3C2-20(Ni20Cr) and (HAp). The diffractograms revealed the presence of both metallic/cermet phases derived from the carbide and crystalline HAp powders. Specifically, peaks corresponding to Cr3C2 and Cr7C3 carbides, Ni (as a solid solution with Cr), and Cr2O3 were detected. At the same time, well-defined diffraction peaks associated with crystalline HAp were also observed, confirming that the thermal spray process preserved the characteristic phases of the carbide alloy and the HAp powder.
In all coatings containing silicon, that is, the C3, C4, and C5 samples in Figure 1, the XRD patterns reveal that HAp and carbide phases are consistently observed, with no significant changes in peak position, whereas, as expected, the peaks associated with Si became larger as the Si content increased. A noticeable broadening of the diffraction peaks was observed in the 2θ range between 40° and 50°, particularly for those peaks associated with the Cr7C3 phase. This broadening is most noticeable in the peaks associated with the Cr7C3 phase and became more pronounced as the silicon content increased. Such broadening can be attributed to the coexistence of amorphous and crystalline structures within the coatings. This interpretation is consistent with reports in the literature that associate broad XRD peaks in this range with the formation of metastable amorphous phases in NiCr-based alloys, silicon being an element with a tendency to form such phases under rapid cooling conditions [36,37,38,39,40,41]. The presence of amorphous phases in these coatings can significantly influence their overall performance. On the positive side, silicon-containing amorphous phases often enhance bioactivity, as they typically participate in dissolution processes in physiological environments [42]. On the downside, amorphous phases are generally less stable than their crystalline counterparts, which may negatively impact the wear resistance of coatings, an essential property for dental applications where mechanical stress and abrasion are prevalent. In particular, the long-term mechanical stability of amorphous phases is a concern, as dissolution processes of Si-based amorphous phases may alter the coating’s mechanical behavior [43]. Although this study primarily focuses on electrochemical and in vitro performance, the influence of silicon content on the formation of amorphous phases and its subsequent effect on the coatings’ mechanical properties is a topic that deserves further investigation. In fact, future work should also address critical mechanical aspects such as adhesion strength and wear resistance. These are essential aspects for dental applications, as demonstrated in similar systems that perform well under simulated masticatory conditions [44,45,46].
Figure 2 shows the cross-sectional SEM–BSE micrographs of the coatings obtained via the HVOF process. All coatings exhibited typical microstructural features of thermally sprayed coatings, notably the formation of a characteristic lamellar structure. This layered architecture results from the successive deposition and rapid solidification of molten or semi-molten particles. It reflects the dynamics of particle melting, deformation upon impact, and solidification, which are essential for ensuring coating cohesion. Two main phases were distinguishable in the backscattered electron images, i.e., a light gray phase, corresponding to the NiCr matrix, and a darker phase, representing the ceramic Cr-based carbides (Cr3C2 and Cr7C3), typically appearing as grain-like or angular inclusions in the coatings. The metallic matrix contains chromium in solid solution, with concentrations ranging from 20 to 50 wt%, depending on the local chemical composition and thermal gradients during the spraying process [47]. In the case of the C1 coating, the microstructure revealed a heterogeneously distributed dispersion of Cr3C2 particles embedded within a crystalline NiCr matrix.
Figure 2 also shows that the microstructure of the C3 to C5 coatings exhibited combined characteristics of both NiCr and Cr3C2 phases, along with the distinct presence of ceramic particles corresponding to HAp. This composite architecture confirms the successful incorporation of HAp into the NiCr-Cr3C2 matrix, which is essential for enhancing bioactivity while preserving the mechanical strength typically provided by both the metallic and carbide phases. Figure 3 shows the SEM–BSE micrograph and corresponding EDS elemental mapping of the NiCr-Cr3C2-20 wt% HAp coating (C2) cross-section. The elemental maps reveal the presence and spatial distribution of calcium (Ca) and phosphorus (P), which are the principal constituents of HAp. These elements appear concentrated in the large dark cavities observed in the BSE image (see arrow in Figure 3), confirming the successful incorporation of the ceramic compound within the metallic–carbide matrix. Their morphology suggests that the HAp particles were partially or fully molten during the HVOF process, leading to the formation of molten and semi-molten ceramic lamellas. The HAp phase is distinguishable from the surrounding matrix, which is rich in Ni and Cr. Overall, the coating thicknesses obtained for the C1, C2, C3, C4, and C5 samples were 140 ± 10 μm, 105 ± 15 μm, 84 ± 12 μm, 65 ± 11 μm, and 35 ± 15 μm, respectively.
In addition, Figure 2 also shows that as the silicon content increased in the C3, C4, and C5 coatings, a noticeable reduction in coating thickness was observed. This effect may be associated with the role of silicon in altering the thermal dynamics during the HVOF spraying process. One possible explanation is that the simultaneous spraying of silicon, carbide, and HAp particles reduces the net energy transferred from the flame to the carbide and HAp particles. Silicon has a lower melting point (1414 °C) compared with Cr3C2 (~1930 °C) and slightly higher than that of the NiCr matrix (~1336 °C) [48]; its latent heat of fusion (1632 kJ/kg) is higher than those of HAp (342.5 kJ/kg) and Cr3C2 (529 kJ/kg) [49,50,51]. During spraying, silicon-rich particles melt more readily, requiring a considerable amount of thermal energy from the flame and competing with the energy required to melt the Cr3C2 and HAp particles. This thermal competition may reduce carbide melting efficiency and promote faster cooling, favoring the formation of metastable amorphous Si, as reported in the XRD patterns.
Figure 4 presents the SEM–BSE micrograph and corresponding EDS elemental map of the C4 coating, which consists of Cr3C2–NiCr, HAp, and silicon. The elemental distribution confirms the presence of silicon throughout the coating, with detectable concentrations in multiple regions. Silicon appears as a dispersed phase coexisting with the (Ni, Cr) metallic matrix and the ceramic phase, indicating its successful incorporation into the coating during the HVOF process, as reported in the XRD patterns. Si does not segregate into isolated regions but instead integrates with both the ceramic and metallic/carbide particles. This result is crucial for a possible dental application since dispersion of silicon within the coating could influence various functional properties, such as corrosion resistance and bioactivity, by modifying the local chemistry and cellular response at the metal–ceramic interface [52].
Figure 5 shows the surface morphology of the HVOF coatings observed by SEM. Figure 5 displays the top surface of the C1 coating, which is characteristic of carbide coatings; that is, a typical morphology composed of molten and semi-molten splats, along with fragmented particles resulting from the high-velocity impact and rapid solidification of the sprayed material. This type of morphology is common in Cr3C2-NiCr thermally sprayed coatings [36].
In contrast to C1 coating, the top surface morphologies of the C2 to C5 coatings (Figure 5) exhibit noticeable changes associated with the incorporation of HAp and silicon. While molten and semi-molten splats are still present, indicating partial or complete melting of NiCr and Cr3C2 particles, these surfaces also reveal a finer and more fragmented morphology. In particular, the appearance of smaller/irregular fragments dispersed across the surface may be attributed to the presence of unmolten or partially molten HAp particles, which is a typical feature of HAp-based coatings obtained via HVOF [53].
Figure 6 presents the potentiodynamic polarization curves of the evaluated coating systems in artificial saliva. The results reveal a similar electrochemical behavior across all samples, with corrosion potential Ecorr values ranging from –281 mV to –148 mV (see Table 3). Additionally, the corrosion current densities icorr were comparable among the coatings, with the C1 and C2 coatings showing the lowest icorr values, suggesting enhanced corrosion resistance. This improved performance may be attributed to the presence of chromium carbide (Cr3C2) within the matrix, which acts as a corrosion-resistant phase by forming stable oxide layers that inhibit further electrochemical dissolution [54]. Furthermore, the presence of the HAp phase in the C2 coating may also contribute to improved corrosion resistance due to its ceramic nature. It is important to note that the anodic branch for the C1 coating showed a small passive region between approximately 100 and 300 mV, which is likely associated with the formation of a stable protective chromium oxide film [54]. Such a layer is lost at higher polarization potentials, as the current density increases while the system is further deviated from equilibrium conditions. Also, the anodic branches of the rest of the coatings did not exhibit well-defined passive regions, indicating an ongoing active dissolution process under the imposed test conditions.
Figure 7 presents OCP measurements of the coatings evaluated in artificial saliva for 168 h. OCP is commonly used to assess the activity/nobility of material surfaces, which is directly related to the formation or absence of a passive protective layer. An increase in OCP towards more positive values generally indicates the development of a passive film on the surface, reflecting improved corrosion resistance. Conversely, stable OCP values over time suggest that the passive layer remains intact and protective [55]. In this study, the C1 coating exhibited the noblest potential, stabilizing quickly around –16 mV throughout the test duration. This behavior indicates spontaneous passivation, likely due to the saturation of chromium carbide (Cr3C2) particles at the surface, which promotes the formation of a stable passive film [56,57].
In contrast, the other coatings showed more negative OCP values. They experienced an initial drop in potential within the early hours, which is typically associated with the rupture, dissolution, or lack of formation of a protective layer [58]. Following this, fluctuations and oscillations in the potential curves were observed. Such oscillations suggest the absence of a stable, continuous passive film [59]. The presence of HAp in the coatings may explain this behavior. It has been reported that HAp undergoes a dissolution–precipitation process during short immersion times in artificial saliva, which can contribute to a potential instability of HAp-containing coatings [60]. Notably, except for the C5 coating, which stabilized around −59 mV, the rest of the coatings completed the testing with more negative potentials, grouped between −164 mV and −199 mV. The potential values observed for the C5 coating could be related to its reduced thickness (see Figure 2), which may more readily allow for electrolyte penetration through pores or cracks, allowing it to reach the titanium substrate. This hypothesis aligns with reported titanium substrate potentials under similar in vitro test conditions, reaching approximately −60.4 mV [32].
Figure 8 displays the LPR plots for the five coatings evaluated. The results indicate that the C1 coating showed the highest corrosion resistance, followed by the C2 sample. Previous studies [61] have associated the corrosion resistance of NiCr-based alloys in artificial saliva with the formation of a uniform and stable passive oxide layer. This passive layer effectively limits electrochemical degradation by preventing the penetration of aggressive ions. Among the coatings with added Si, and according only to LRP results, the C5 coating exhibited a fair corrosion resistance, possibly due to the electrolyte penetration into the substrate, as mentioned previously. It is also important to note that all coatings exhibited unstable behavior. Some studies [55,62] have shown that incorporating Si into HAp enhances its bioactivity by accelerating the dissolution–reprecipitation processes involved in apatite formation, causing this process to occur more rapidly than with HAp alone. This increase in biological activity has been attributed to various factors that produce a synergistic effect, such as a Si-substituted CaP structure and a more electronegative surface, which in turn is associated with the exchange of SiO44− for PO43−. This behavior promotes the formation of an apatite film on the coated surface immersed in saliva for long periods (>24 h) [55].
Some studies [63,64] have suggested that the presence of silicon enhances the bioactivity of implants not only by promoting the formation of an apatite layer but also by facilitating protein adsorption and improving osteoblast adhesion. These combined effects contribute to a better integration between the implant and the surrounding bone tissue, ultimately accelerating bone apposition. Therefore, lower corrosion resistance does not necessarily imply poor implant performance; instead, this electrochemical activity may be associated with both ionic interaction with the surrounding physiological body fluid and biological processes that support and enhance the osseointegration of the implant [63].
Figure 9 shows the EIS plots in Nyquist (A) and phase angle (B) formats for the different coatings evaluated in artificial saliva. The Nyquist plot (Figure 9(A1)) for the C1 coating exhibits a charge transfer resistance process, as evidenced by its tendency to form a semicircular arc, which reflects the ease or difficulty with electrons interchange at the coating/electrolyte interface. Meanwhile, the phase angle plot (Figure 9(B1)) displays a single time constant with phase values close to 90°, which is characteristic of capacitive behavior. This suggests that the coating acts as an effective dielectric barrier, impeding ionic movement and indicating the presence of a passive film. In such plots, a high and broad phase angle near 90° over a wide frequency range indicates an ideal capacitive behavior, typically associated with homogeneous and dense protective layers [65]. This electrochemical response is consistent with the dense microstructure of the C1 coating (see Figure 1), a feature commonly observed in NiCr-based alloys used in dental environments [35,66], and with the LPR results discussed previously, as it correlates with the highest resistance values (between 6656 and 16,307 Ω·cm2) determined from the Bode plot in the impedance modulus format.
Also, the C2, C3, and C4 coatings exhibited a charge transfer resistance process, as observed in the Nyquist plots (Figure 9(A2,A3,A4)). However, the resistance values obtained from the |Z| Bode plots were lower than those recorded for the C1 coating, ranging from 235 to 413 Ω·cm2 for the C2 coating (NiCr-Cr3C2 + HAp), and from 108 to 135 Ω·cm2 and 123 to 185 Ω·cm2 for the C3 and C4 coatings (including 5% and 10% Si additions), respectively. These lower impedance values indicate the reduced barrier properties and higher ionic permeability of these coatings when compared to those of the dense C1 coating. Furthermore, the phase angle plots (Figure 9(B2,B3,B4)) for C2, C3, and C4 revealed the presence of two distinct time constants, along with maximum phase angle values around 70°. As previously outlined, such a behavior is typically associated with the formation of dual-layer structures in the coating [67]. The appearance of two time constants reflects electrochemical processes occurring at different depths, one corresponding to the interaction between the electrolyte and the outer porous region, and other associated with charge transfer at the more compact inner layer. The incorporation of HAp, a material well-known for its bioactivity but also for its intrinsic porosity when thermally sprayed, contributes significantly to this behavior. In addition, the incorporation of Si into the C3 and C4 coatings appears to influence their electrochemical response, likely due to significant alterations in the coating structure compared to those of C1 and C2 [55]. As observed in Figure 2, the microstructures of C3 and C4 are noticeably more porous than that of C2, revealing a dual-phase microstructure with distinct regions corresponding to the NiCr matrix and the embedded HAp particles, which correlate with the lower impedance values, broader phase angle transitions, and dual-time constant response observed in these samples. This trend is consistent with the results of previous studies, where bioactive coatings, particularly those based on HAp, exhibited lower corrosion resistance and multiple time constant behavior due to their inherent porosity and structural heterogeneity [67].
The C5 coating exhibited a distinct electrochemical behavior (Figure 9(A5,B5)). It showed relatively higher charge transfer resistance values (432–843 Ω·cm2) in the OCP and LPR sections, associated with electrolyte infiltration reaching the Ti substrate, as previously discussed. Although higher charge transfer resistance values typically suggest better corrosion resistance; in this case, these values result from interfacial changes due to substrate exposure and the formation of corrosion products rather than from a more protective coating. In fact, this behavior is indicative of a more porous structure in comparison to the coatings with 5% and 10% Si additions (C3 and C4), as supported by the phase angle values ranging from 50° to 30° (Figure 9(B5)). Such low phase angle values are characteristic of diffusion-controlled processes, where aggressive species from the solution progressively penetrate the coating, eventually reaching the substrate. Additionally, a shift toward lower frequencies was observed in the EIS spectra, which further supports the presence of a degradation mechanism, such as coating dissolution. This process may be driven by the leaching of silicon, which may promote apatite formation in the medium, and by the reduced compactness of the C5 coating. Similar phenomena have been reported in the literature for porous or weakly adherent HAp-based coatings [68]. SEM–BSE images in this work confirm the higher structural heterogeneity and reduced thickness of the C5 coating. Moreover, an increase in the solution resistance was also observed (see Figure 9—Nyquist plot), which may be attributed to changes in the ionic strength of the electrolyte due to the dissolution of Si or HAp. This phenomenon has been previously documented in systems where the partial dissolution of bioactive phases alters the conductivity of the medium [69,70]. Table 4 summarizes the results obtained from EIS measurements.
On the other hand, the cytotoxicity of the coatings was evaluated using the MTT assay on 3T3-L1 fibroblasts after 24 and 72 h of exposure; the IC50 values were determined for each condition. As shown in Table 5, the C1 to C4 coatings exhibited IC50 values greater than 1000 µg/mL at both testing times, indicating high cytocompatibility and no significant cytotoxic effects within the tested concentration range [71]. These results suggest that the incorporation of HAp and moderate amounts of silicon (up to 10%) into the NiCr-Cr3C2 matrix does not adversely affect cell viability, at least under the experimental conditions employed here. This result is also supported by previous studies, such as in Si/Sr-co-doped HAp coatings, where Si contents around 8 wt% significantly enhanced osteogenic potential without any risk of negative cellular responses [72]. The lack of cytotoxicity is consistent with the expected biocompatibility of HAp, which is widely used in biomedical applications, and with the controlled incorporation of Si, which has been shown to promote bioactivity when maintained within optimal levels.
In contrast, the C5 coating, which contains the highest silicon content (20% Si), showed a distinct time-dependent cytotoxic response. While its IC50 value remained above 1000 µg/mL at 24 h, suggesting no acute toxicity, it decreased to 535.7 ± 67.2 µg/mL after 72 h of exposure. This reduction in IC50 indicates a delayed or progressive toxic effect, potentially resulting from the gradual release of soluble silicon species or other degradation products over time. Such species may alter the ionic environment or affect cellular metabolism, leading to decreased cell viability after prolonged exposure [71]. This observation is supported by electrochemical and morphological data presented earlier in this study. The C5 coating exhibited higher charge transfer resistance and more pronounced phase angle dispersion, indicative of a more porous and less stable surface. SEM–BSE images in Figure 2 further confirmed a lower thickness and greater heterogeneity in C5 compared with the results for the rest of the coatings, which could facilitate electrolyte penetration and increase the release of ionic or particulate matter to the surrounding medium. High concentration and dissolution of Si-rich phases have been previously reported to alter the local ionic strength, pH, and surface chemistry, all of which may contribute to changes in cell viability behavior [73]. In fact, similar dose-dependent cytotoxic effects have been previously reported, where high concentrations of Si particles reduced cell viability and triggered oxidative stress and inflammatory responses [74], supporting the fact that excessive Si metal content can shift the surface response from being bioactive to cytotoxic.
From a practical standpoint, coatings like C2 to C4 appear most suitable for clinical use, as they balance bioactivity with structural stability. These coatings may contribute to extended implant lifespan by promoting early osseointegration while resisting electrochemical degradation. The cytocompatibility results demonstrate that the C1 to C4 coatings are suitable candidates for dental applications involving direct contact with saliva, soft tissues, or bone, as they maintain high cell viability across all tested concentrations and testing times. This behavior is linked to their bioactive nature, resulting from the synergistic effects of HAp and silicon incorporation. It is worth noting that the electrochemical tests were conducted for 168 h to simulate early-stage exposure in an oral environment, but long-term in vivo conditions may involve more complex and prolonged degradation processes. Based on the current findings, coatings such as C1 and C2, with dense microstructures, high charge transfer resistance, and capacitive behavior, are expected to offer prolonged protection against corrosion. However, coatings with higher silicon content (e.g., C5) exhibited signs of structural heterogeneity and electrolyte infiltration, which could accelerate long-term degradation via mechanisms such as increased ion diffusion, localized corrosion, or matrix dissolution. The porous nature and moderate phase angles observed in C3 and C4 suggest the potential for sustained bioactivity but also raise the possibility of gradual microstructural breakdown over time. Future studies involving extended immersion times, cyclic loading, or in vivo models would be essential to validate the long-term stability and degradation pathways of these coatings.

4. Conclusions

This study evaluated the electrochemical and biological performance of Cr3C2-20(Ni20Cr) and Cr3C2-20(Ni20Cr)–20HAp–XSi coatings deposited via the high-velocity oxy-fuel (HVOF) technique. The coatings were applied onto commercially pure titanium substrates, with the silicon content varied between X = 0, 5, 10, and 20 wt%. To evaluate corrosion resistance, electrochemical techniques including potentiodynamic polarization curves, linear polarization resistance (LPR), electrochemical impedance spectroscopy (EIS), and open circuit potential (OCP) measurements were employed. Artificial saliva was used as the corrosive medium at 37 °C for 168 h. The results revealed that the C1 and C2 coatings exhibited the best corrosion resistance. The C1 coating showed a dense and homogeneous microstructure, which resulted in a high charge transfer resistance (>6000 Ω·cm2), and phase angles close to 90°, indicative of an ideal capacitive behavior. The C2 coating, composed of a NiCr matrix reinforced with Cr3C2 and HAp, also showed enhanced corrosion resistance due to both the protective oxide-forming matrix and the bioactive behavior provided by the HAp. The results indicated that the coating with 20 wt% Si exhibited the poorest electrochemical performance, attributed to faster dissolution processes and associated with a thinner coating layer, as confirmed by SEM–BSE evaluation. Interestingly, the presence of Cr-based phases such as Cr2O3 and Cr7C3 in the coatings produced favorable results in the electrochemical and cytocompatibility assays. Specifically, the C3 and C4 coatings, containing 5% and 10% silicon, respectively, and the Cr-based phases, exhibited lower impedance values and more porous microstructures. Their electrochemical response was characterized by two distinct time constants and moderate phase angles (~70°). These features reflect their bioactive nature, likely due to the synergistic effect of HAp and Si. Conversely, the C5 coating (20% Si) displayed the lowest overall electrochemical performance. Despite presenting moderately high charge transfer resistance (432–843 Ω·cm2), the coating exhibited signs of electrolyte infiltration, high structural heterogeneity, and significant dissolution behavior. Biological evaluation by MTT assays using 3T3-L1 fibroblasts revealed that the C1 to C4 coatings exhibited no significant cytotoxic effects, maintaining IC50 values above 1000 µg/mL at both 24 and 72 h, thus confirming their excellent cytocompatibility. On the other hand, the C5 coating showed a time-dependent reduction in cell viability, with an IC50 value decreasing to 535.5 µg/mL after 72 h. This behavior may be associated with the progressive release of soluble Si species or degradation products that alter the ionic microenvironment and impact cellular activity. In summary, the C2 to C4 coatings are promising candidates for dental applications, particularly those in contact with bone tissue, due to their combined electrochemical reactivity and bioactive potential. These coatings maintain high cell viability while enabling dissolution–precipitation mechanisms that may support osseointegration.

Author Contributions

Conceptualization, J.H.; Methodology, J.H., O.S.-M., R.M.M.-R. and J.C.-C.; Validation, D.G.E.-A.; Formal analysis, O.S.-M. and V.M.Z.; Investigation, J.H., O.S.-M. and R.M.M.-R.; Resources, C.A.P.-S., A.G.-B. and A.L.I.-G.; Data curation, A.G.-B. and V.M.Z.; Writing—original draft, J.H. and O.S.-M.; Writing—review & editing, C.A.P.-S., D.G.E.-A. and A.G.-B.; Supervision, A.L.I.-G. All authors have read and agreed to the published version of the manuscript.

Funding

This research received no external funding.

Acknowledgments

The authors gratefully acknowledge the financial and institutional support provided by the Secretaría de Ciencia, Humanidades, Tecnología, e Innovación (SECIHTI) through the “Investigadores por México” program under project number 848, which made this work possible. Special thanks are also extended to the Laboratorio Nacional de Proyección Térmica (CENAPROT) for the use of its facilities and for providing technical assistance during the experimental phase of this study.

Conflicts of Interest

The authors declare no conflict of interest.

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Applsci 15 09308 g001
Figure 1. XRD patterns from the as-sprayed coatings.
Figure 1. XRD patterns from the as-sprayed coatings.
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Figure 2. SEM micrographs from the cross-section of the evaluated coatings in artificial saliva at 37 °C after 168 h.
Figure 2. SEM micrographs from the cross-section of the evaluated coatings in artificial saliva at 37 °C after 168 h.
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Figure 3. SEM micrographs from the cross-section and element mapping of the C2 coating. The arrow indicates large dark cavities where calcium (Ca) and phosphorus (P) are present.
Figure 3. SEM micrographs from the cross-section and element mapping of the C2 coating. The arrow indicates large dark cavities where calcium (Ca) and phosphorus (P) are present.
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Figure 4. SEM micrographs from the cross-section and element mapping of the C4 coating, showing the Si-rich zones.
Figure 4. SEM micrographs from the cross-section and element mapping of the C4 coating, showing the Si-rich zones.
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Figure 5. SEM micrographs from the surface of the as-sprayed coatings.
Figure 5. SEM micrographs from the surface of the as-sprayed coatings.
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Figure 6. Potentiodynamic polarization curve plots of the coatings evaluated in artificial saliva at 37 °C.
Figure 6. Potentiodynamic polarization curve plots of the coatings evaluated in artificial saliva at 37 °C.
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Figure 7. OCP plots of the coatings evaluated in artificial saliva at 37 °C during 168 h.
Figure 7. OCP plots of the coatings evaluated in artificial saliva at 37 °C during 168 h.
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Figure 8. LPR plots of the coatings evaluated in artificial saliva at 37 °C during 168 h.
Figure 8. LPR plots of the coatings evaluated in artificial saliva at 37 °C during 168 h.
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Figure 9. EIS plots in (A) Nyquist and (B) phase angle format of the (1) C1, (2) C2, (3) C3, (4) C4, and (5) C5 coatings evaluated in artificial saliva at 37 °C over 168 h.
Figure 9. EIS plots in (A) Nyquist and (B) phase angle format of the (1) C1, (2) C2, (3) C3, (4) C4, and (5) C5 coatings evaluated in artificial saliva at 37 °C over 168 h.
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Table 1. HVOF coatings obtained in this work.
Table 1. HVOF coatings obtained in this work.
Sample IDCoating CompositionCr3C2-20(Ni-20Cr)
(wt%)		HAp
(wt%)		Si
(wt%)
C1Cr3C2-20(Ni-20Cr)10000
C2Cr3C2-20(Ni-20Cr)-20HAp80200
C3Cr3C2-20(Ni-20Cr)-20HAp-5Si75205
C4Cr3C2-20(Ni-20Cr)-20HAp-10Si702010
C5Cr3C2-20(Ni-20Cr)-20HAp-20Si602020
Table 2. Chemical composition of artificial saliva.
Table 2. Chemical composition of artificial saliva.
ComponentsContent
(g/L)
NaCl0.600
KCl0.720
CaCl2-2H2O0.220
KH2PO40.680
Na2HPO4-12H2O0.856
KSCN0.060
NaHCO31.500
Citric acid0.030
Table 3. Electrochemical parameters obtained from polarization curves.
Table 3. Electrochemical parameters obtained from polarization curves.
Coating SampleEcorr
(mV)		icorr
(mA/cm2)		Ba
(mV)		Bc
(mV)
C1−2814.78 × 10−2920.1497
C2−1481.84 × 10−2803.7300.6
C3−2369.35 × 10−21365.5776.8
C4−2131.31 × 10−11452.51047.4
C5−2361.13 × 10−11044.11017.2
Table 4. Electrochemical parameters obtained from EIS curves.
Table 4. Electrochemical parameters obtained from EIS curves.
Coating SampleRct
(Ω·cm2)		Phase Angle
(°)
C16656 to 1630770 to 80
C2235 to 41370 to 75
C3108 to 13565 to 70
C4123 to 18564 to 68
C5432 to 84325 to 55
Table 5. Results of cell viability (IC50) for the evaluated coating samples at 24 and 72 h.
Table 5. Results of cell viability (IC50) for the evaluated coating samples at 24 and 72 h.
Coating SampleIC50 (µg/mL)
24 h72 h
C1>1000>1000
C2>1000>1000
C3>1000>1000
C4>1000>1000
C5>1000535.5
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Henao, J.; Sotelo-Mazon, O.; Montiel-Ruiz, R.M.; Poblano-Salas, C.A.; Espinosa-Arbelaez, D.G.; Corona-Castuera, J.; Giraldo-Betancur, A.; Islas-Garduño, A.L.; Zezatti, V.M. Electrochemical Performance and Cytocompatibility of HVOF-Sprayed Cr3C2-20(Ni20Cr)-20HAp-XSi Coatings for Dental Applications. Appl. Sci. 2025, 15, 9308. https://doi.org/10.3390/app15179308

AMA Style

Henao J, Sotelo-Mazon O, Montiel-Ruiz RM, Poblano-Salas CA, Espinosa-Arbelaez DG, Corona-Castuera J, Giraldo-Betancur A, Islas-Garduño AL, Zezatti VM. Electrochemical Performance and Cytocompatibility of HVOF-Sprayed Cr3C2-20(Ni20Cr)-20HAp-XSi Coatings for Dental Applications. Applied Sciences. 2025; 15(17):9308. https://doi.org/10.3390/app15179308

Chicago/Turabian Style

Henao, John, Oscar Sotelo-Mazon, Rosa M. Montiel-Ruiz, Carlos A. Poblano-Salas, Diego G. Espinosa-Arbelaez, Jorge Corona-Castuera, Astrid Giraldo-Betancur, Ana L. Islas-Garduño, and Victor M. Zezatti. 2025. "Electrochemical Performance and Cytocompatibility of HVOF-Sprayed Cr3C2-20(Ni20Cr)-20HAp-XSi Coatings for Dental Applications" Applied Sciences 15, no. 17: 9308. https://doi.org/10.3390/app15179308

APA Style

Henao, J., Sotelo-Mazon, O., Montiel-Ruiz, R. M., Poblano-Salas, C. A., Espinosa-Arbelaez, D. G., Corona-Castuera, J., Giraldo-Betancur, A., Islas-Garduño, A. L., & Zezatti, V. M. (2025). Electrochemical Performance and Cytocompatibility of HVOF-Sprayed Cr3C2-20(Ni20Cr)-20HAp-XSi Coatings for Dental Applications. Applied Sciences, 15(17), 9308. https://doi.org/10.3390/app15179308

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