Corrosion Properties and Performance of Nanostructured Multilayered Chromium–Amorphous Carbon Coatings on HS6-5-2 Steel

Abstract

Magnetron-sputtered coatings consisting of multiple alternating layers of chromium and amorphous carbon (Cr/a-C)ml were deposited on HS6-5-2 steel with an intermediate chromium layer by varying deposition rates. Three series of coatings, S1, S2, and S3, with thicknesses of 1.74, 1.15, and 1.14 μm and average chromium contents of 89.3, 66.0, and 59.7 wt.% Cr, respectively, were obtained. Open-circuit potential, cyclic potentiodynamic measurements, and electrochemical impedance spectroscopy were used to characterize their corrosion resistance in 3.5% NaCl. The surfaces were observed with optical and scanning electron microscopy before and after the corrosion tests, and changes in the elemental composition were monitored by energy-dispersive spectroscopy. The protective properties of coatings from series S2 and S3 are similar and significantly better than those of S1. They are characterized by a corrosion current below 1 μA cm^–2 and a stable passive state up to over 0.9 V_Ag/AgCl. The coatings have cathodic behavior towards the substrate, and when the coatings are damaged, galvanic corrosion causes deep pits. Coatings deposited at lower rates and with higher carbon content demonstrate significantly enhanced corrosion resistance in 3.5% NaCl. All three series of Cr/(Cr/a-C)ml@HS6-5-2 exhibit identical corrosion behavior after compromising the coatings’ integrity.

1. Introduction

Physical vapor deposition (PVD) coatings play a critical role in surface engineering, significantly enhancing the performance and durability of tools and components in industries such as aerospace, automotive, and manufacturing. These coatings, typically composed of nitrides, carbides, or carbonitrides like chromium nitride (Cr-N), chromium carbide (Cr-C), or chromium carbonitride (Cr-C-N), improve steel substrates by reducing friction, enhancing thermal stability, and increasing corrosion resistance in diverse environments, including high-temperature, aqueous, and biological settings [1,2,3,4,5]. The protective capabilities of PVD coatings make them essential for applications requiring resilience under harsh mechanical and chemical conditions. However, growth-related defects such as pores and pinholes often compromise their effectiveness, allowing corrosive agents to penetrate and attack the less noble substrate material. Achieving robust corrosion protection requires coatings with high corrosion resistance and minimal porosity, properties governed by careful material selection, surface pretreatment, and optimized deposition parameters. Quantitative evaluation of corrosion performance remains challenging due to the complex interplay of the microstructural, chemical, and electrochemical factors influencing coating behavior.

Physical vapor deposition involves vaporizing material from a solid source in a vacuum environment and depositing it as a thin, adherent film on a substrate. Magnetron sputtering, for example, provides precise control over coating composition and structure, enabling the production of films with tailored crystallinity—ranging from amorphous to polycrystalline or nanocomposite—and thicknesses from nanometers to micrometers. These coatings enhance substrate performance by reducing friction in sliding contacts, which is critical for cutting tools and mechanical components, and by increasing hardness to resist mechanical deformation and abrasive wear. Additionally, PVD coatings maintain structural integrity at elevated temperatures, a vital property for high-speed machining, and form protective oxide layers to inhibit corrosion in aggressive environments. However, defects like porosity and pinholes create pathways for corrosive species, initiating localized corrosion. Optimizing deposition conditions, such as substrate temperature, bias voltage, and gas pressure, alongside surface pretreatment techniques like ion etching or polishing, can minimize these defects and enhance coating performance.

Chromium carbon (Cr-C)-based thin films are particularly promising for applications requiring a balance between mechanical robustness and corrosion resistance [6,7,8,9,10,11,12,13]. These films form a passive chromium oxide (Cr₂O₃) layer under various conditions, including high-temperature exposure in air [14,15,16], supercritical water [17], aqueous solutions [18,19,20,21], and chloride-containing non-aqueous media [22]. This passive layer acts as a barrier, reducing electrochemical activity and inhibiting further corrosion. The corrosion resistance of Cr-C films is strongly influenced by their composition, particularly the carbon content, and microstructure, which can be tailored through deposition techniques like magnetron sputtering. Depending on deposition conditions, Cr-C films may exhibit amorphous structures, which offer uniform properties but potentially lower hardness; polycrystalline structures, which provide enhanced mechanical strength but may have vulnerable grain boundaries; or nanocomposite structures, consisting of nanocrystalline carbide grains embedded in an amorphous carbon matrix that combines high hardness with low friction [20,21,23]. Despite their potential, the relationship between composition, morphology, and corrosion resistance remains underexplored. For instance, higher carbon content can enhance lubricity but may destabilize the passive oxide layer, reducing corrosion protection [24]. Systematic studies are needed to clarify how deposition parameters, such as sputtering power, gas ratio, and substrate temperature, affect film structure and performance in corrosive environments. Importantly, the influence of these parameters is strongly substrate-dependent and represents a gap that has not yet been sufficiently investigated. This article aims to bridge this gap.

The HS6-5-2 high-speed steel (HSS, 1.3343, BDS EN ISO 4957:2018 [25]) is widely utilized in demanding applications due to its optimized chemical composition and exceptional mechanical properties. This steel contains 5.50–6.75 wt% tungsten, which imparts superior wear resistance, hot hardness, and secondary hardness, making it ideal for high-speed cutting tools. Among W-based HSS grades, HS6-5-2 offers the best combination of wear resistance, red hardness (the ability to retain hardness at elevated temperatures), and toughness [26]. It also includes 0.65–0.80 wt% carbon, which facilitates the formation of hard carbides, such as tungsten carbide (WC) and chromium carbide (Cr₇C₃), contributing to its high hardness and cutting performance. Chromium, which is present at approximately 5 wt%, enhances hardenability, toughness, and corrosion resistance by forming a thin oxide layer, though this is insufficient for aggressive environments. Vanadium, typically at 1–2 wt%, improves temper resistance, secondary hardening, and wear resistance by stabilizing the microstructure and forming fine vanadium carbides [26]. These properties make HS6-5-2 suitable for cutting tools like drills, milling cutters, and taps, where high hardness and red hardness are critical, as well as for cold-forming tools, such as punches, dies, and cold extrusion rams, which require high tensile strength and wear resistance to withstand compressive forces. Additionally, the steel is used in precision components like bearings and gears, where durability under cyclic loading is essential [27].

Despite its mechanical strengths, HS6-5-2 steel exhibits poor corrosion resistance, particularly in humid, chloride-rich, or marine environments. It is susceptible to general rusting, where uniform corrosion occurs in the presence of moisture, forming iron oxides that degrade the surface. Steel is also prone to pitting corrosion, a localized attack that creates small cavities due to the breakdown of weak protective oxide films. In chloride-containing environments, such as marine settings or areas exposed to de-icing salts, pitting is exacerbated by the formation of micro-galvanic cells, where small anodic pits corrode rapidly relative to the larger cathodic matrix [28]. The presence of carbide-forming elements, such as tungsten and chromium, contributes to micro-galvanic corrosion due to electrochemical potential differences between carbides and the iron matrix. These carbides act as cathodes, accelerating localized corrosion at the matrix–carbide interface [29]. The corrosion susceptibility of HS6-5-2 steel limits its use in aggressive environments unless protected by coatings. Despite the widespread use of this steel and its well-documented low corrosion resistance, there are no reported studies addressing the application of nonreactive magnetron sputtering protective coatings or the evaluation of their corrosion behavior. From this standpoint, nanostructured multilayer Cr/(Cr/a-C)ml coatings offer a viable strategy to mitigate these limitations by acting as an effective barrier with high corrosion resistance while simultaneously reducing surface friction and wear. To achieve this functionality, the deposition parameters of the multilayer coatings must be optimized to ensure strong adhesion to the substrate and minimal coating porosity.

In this study, the influence of the deposition rate of a (Cr/a-C)ml coating on the corrosion properties of the system ((Cr/a-C)ml on HS6-5-2) in a model corrosion environment of 3.5% was evaluated. The most suitable multilayer deposition conditions in terms of their resistance to general and pitting corrosion were determined using the open-circuit potential, cyclic potentiodynamic polarization, and electrochemical impedance spectroscopy in combination with microscopic observations and elemental analysis. The results are compared with the corrosion behavior of the HS6-5-2 substrate.

2. Materials and Methods

Substrate blanks with dimensions of 25 × 25 × 4 mm were produced by single commercially available pre-manufactured workpieces. To validate the true elemental composition, X-ray fluorescence (XRF) analysis was performed on 5 spots using a Bruker S1 TITAN 800 handheld spectrometer (Bruker Corporation, Billerica, MA, USA). The obtained results are shown on

Table 1. Measured chemical composition of HS6-5-2.

	Si	P	Ti	V	Cr	Mn	Fe	Co	Ni	Cu	Mo	W
Average	0.88	0.03	0.01	2.08	4.56	0.28	80.79	0.60	0.18	0.11	4.81	5.65
AveDev	0.031	0.004	0.000	0.087	0.136	0.004	0.276	0.016	0.004	0.004	0.027	0.058

.

2.1. Preparation and Deposition of Coatings

The preparation of the substrate surfaces before magnetron sputtering is essential for the adhesion of the coatings. Before coating deposition, both steel substrates were prepared by successive grinding with silicon carbide abrasive papers from grade 200 to grade 1600. Polishing was carried out with a 5 µm diamond polishing agent (Struers GmbH, Willich, Germany). After polishing the surfaces were cleaned in an ultrasonic bath with deionized water and degreased in an ultrasonic bath with the Deconex^® cleaning reagent (4% dissolved in deionized water) at 60 °C for 10 min, followed by subsequent washing in deionized water, drying by air blowing, and final drying for 1 h in a vacuum dryer at 100 °C. All this preparation of the substrates was carried out up to 1 h before they were placed in the chamber of the PVD system to begin the magnetron deposition process [30].

The nanostructured multilayer Cr/(Cr/a-C)ml coating was deposited onto the substrates from HS6-5-2. The coating process was carried out in a proprietary magnetron sputtering PVD system (BG67500B1) [31]. The magnetron sputtering system has an octagonal chamber with three unbalanced magnetrons equipped with rectangular targets and a single-axis rotating table for substrate mounting. This PVD system works with rectangular targets with dimensions of 360 × 102 × 9 mm. Due to the significantly lower speed of magnetron sputtering of carbon compared to the speed of sputtering of chromium, the coatings were obtained by the simultaneous sputtering of three targets, two of which were carbon with a purity of 99.99%, and the third target was chromium with a purity of 99.8% [30]. During coating deposition, the substrates were placed at distance of 60 mm from the targets on a rotating table, as shown on Figure 1.

The procedure for depositing a multilayer coating consisting of a chromium (Cr) adhesion sublayer followed by a nanolaminate structure of Cr and C (Cr/C) is shown in Figure 2, with the aim of enhancing the adhesion, hardness, and tribological properties typical of metal-doped amorphous carbon films.

2.1.1. Process for Series 1

The deposition sequence involves the following precisely controlled steps in the PVD chamber.

(1)

Substrates are loaded with a fixed target-to-substrate distance of 60 mm, thus optimizing the flux of sputtered species for uniform deposition.

(2)

The air inside the chamber is evacuated to a base pressure of 4 × 10⁻³ Pa, thereby removing residual gases and contaminants to prevent incorporation into the growing film.

(3)

Substrates are heated to 225–230 °C for 1 h while initiating turntable rotation at 2 rpm. This rotation ensures coating uniformity on complex geometries, and the elevated temperature promotes adatom mobility for denser film growth.

(4)

A 15 min plasma cleaning step is performed in an Ar + 20% H₂ mixture at 3–4 Pa and a −900 V substrate bias, with heating temporarily disabled. This step removes surface oxides and hydrocarbons via a combination of physical sputtering by Ar⁺ ions and chemical reduction by atomic hydrogen, enhancing subsequent adhesion without excessive substrate etching.

(5)

Heating is resumed to maintain the substrates at 225–230 °C.

(6)

A 15 min metal ion cleaning phase occurs, with the Cr target powered at 1000 W, the Ar pressure at 3.1 × 10⁻¹ Pa, and a −900 V bias. Sputtered Cr⁺ ions provide high-energy bombardment, effectively removing residual contaminants and promoting interfacial mixing for superior adhesion—a technique known as metal ion etching, which is superior to inert gas cleaning for oxide-prone substrates.

(7)

The pure Cr sublayer is deposited for 8 min at a Cr target power of 1800 W, a −90 V bias, and an Ar pressure of 2.5–2.6 × 10⁻¹ Pa (maintained constant hereafter). The reduced bias minimizes resputtering, allowing controlled growth of a metallic Cr interlayer (typically 100–500 nm thick based on standard rates).

(8)

Carbon targets are ramped to 250 W for 1 min, beginning the transition to the composite layer.

(9)

Over 10 min, the carbon target power is increased to 750 W, while the Cr target power is reduced to 700 W, and the substrate temperature is lowered to 160 °C. This ramping creates a functionally graded Cr-to-C interface, reducing thermal and mechanical mismatches, while the temperature drop limits graphitization in the carbon phase.

(10)

Stabilized parameters are held for 180 min, yielding the bulk Cr-doped carbon layer with controlled metal incorporation for optimized hardness and low friction.

(11)

Power supplies and mass flow controllers are shut off for 20 s, followed by gate valve closure.

(12)

Substrates cool down to 120 °C in Ar + H₂ at 1 kPa to prevent thermal shock and oxidation.

2.1.2. Variations in Series 2 and 3

Series 2 and 3 replicate the above protocol, with the sole modification in step 9: the Cr target power is ramped up to 175 W (Series 2) or 135 W (Series 3), as shown in

Table 2. Power of the chromium target power.

Series	S1	S2	S3
Cr power	700 W	175 W	135 W

. This reduction lowers the Cr sputtering rate proportionally to the power level in magnetron mode, decreasing the Cr layer’s thickness and the Cr content in the layers. Concurrently, diminished plasma heating and ion bombardment reduce the rise in substrate temperature. Powers below 135 W risk discharge instability due to insufficient ionization.

These adjustments systematically vary the Cr layer thickness and doping in the carbon matrix, as shown on Figure 2, enabling the study of composition-dependent properties such as hardness, residual stress, and tribological performance—common traits in optimizing metal-doped amorphous carbon coatings.

This refined procedure exemplifies unbalanced magnetron sputtering practices, where a sequence of processes incorporates substrate cleaning, metal ion etching, and graded interlayers to achieve robust, functional (Cr/C)ml coatings with applications where wear-resistance of components is needed.

2.2. Layer Thickness Measurement

The coating thicknesses of the different samples in the series were determined in accordance with EN ISO 26423:2016 [32]. Measurements were performed using a KaloMAX II calotester (BAQ GmbH, Braunschweig, Germany). Experiment data was measured by following the test parameters summarized in

Table 3. Test-specific parameters for determination of coating thickness.

Test Parameters	Value
Ball diameter	30 mm
Contact load	0.54 N
Rotational speed of ball	200 rpm
Composition of the abrasive slurry	1 µm diamond paste in ethanol
Slurry feed rate	1 drop/20 s
Duration of the test	20 s

.

2.3. Electrochemical Measurement

The electrochemical tests were carried out in a three-electrode cell for flat electrodes. The surface area of the working electrode exposed to the corrosive environment was limited to 1.0 or 0.1 cm². A silver chloride Ag/AgCl/3.0 M KCl electrode (E_Ag/AgCl = 0.210 V_SHE) and a Pt plate with a surface area of 2 cm² are used as the reference and counter electrodes, respectively. All potential (E) values reported in this work are given with respect to the Ag/AgCl reference electrode.

Control of the electrochemical parameters and acquisition of the corrosion system response were performed using an Autolab PGSTAT302N potentiostat/galvanostat equipped with an FRA32M module and controlled by the NOVA 2.1.8 software (Metrohm, Herisau, Switzerland).

The influence of (Cr/a-C)ml layers on the corrosion resistance of the HS6-5-2 steel substrate was investigated in a 3.5 wt.% NaCl solution. The sequence of the electrochemical tests performed and the settings of their main parameters are presented in Figure 3. Prior to testing, surfaces were degreased with acetone before being brought into contact with the corrosive media. Variation in the open-circuit potential (OCP) with respect to time (t) was recorded for 10 min or longer until the condition dE/dt < 1 µV·s^–1 was met. Electrochemical impedance spectroscopy (EIS) was conducted with a potential amplitude of 10 mV_OCP over a frequency range from 10⁵ to 10⁻² Hz after stabilizing the initial potential.

After a 10 s delay, cyclic voltammetry (CV) was initiated in the anodic direction from −0.7 V_Ag/AgCl, which is the cutoff when either of the two conditions is reached: (i) reaching a current density of 1 mA cm^–2 or (ii) reaching a potential of 1 V_Ag/AgCl. Then the polarization direction was automatically changed to the cathodic direction, and the measurement continued until the reverse branch intersected with the forward branch of the polarization curve. The repassivation potential (E_rp) was determined at this intersection point. The degradation of the protective layer after the applied polarization was assessed by a second EIS test performed after 10 min of potential stabilization or until the condition dE/dt < 1µV·s^–1 is met.

2.4. Microscopic Observation and Surface Characterization

Observations of the surfaces of (Cr/a-C)ml and the native HS6-5-2 steel before and after the corrosion test were carried out using an IM-3MET Metallurgical Microscope (OPTIKA, Ponteranica, Italy) with a digital camera C-B10+ (10 MP) and by a SEM (“EVO MA10 Carl Zeiss”, Oberkochen, Germany) at 20 kV with an integrated energy-dispersive (EDX) X-ray detector system (“Bruker”, Berlin, Germany).

3. Results and Discussion

3.1. Thickness of the (Cr/a-C)ml Coating

The results for the Cr sublayer thickness and the total coating thickness, measured by the calotest method for the different samples, are summarized by series and presented in

Table 4. Result from the measured coating thickness.

Series	S1	S2	S3
Coating thickness, µm	1.74 ± 0.03	1.15 ± 0.03	1.14 ± 0.03
Bi-layer period $\Lambda$, nm	3.8 ± 0.1	3.2 ± 0.1	3.1 ± 0.1
Number of layers	360

. The Cr sublayer in all three series was deposited during step 7, and the measurements indicate a consistent thickness of approximately 0.432 µm.

Once the total coating thickness is known, the thickness of the nanolaminate (Cr/C)ml structure, $T_{nl}$, can be readily determined by subtracting the Cr sublayer’s contribution. As previously defined, the deposition time for the nanostructured multilayer Cr(Cr/C)ml coating was 180 min, with a substrate table rotation speed of 2 rpm.

Based on the thickness of the nanolaminate Cr(Cr/C)ml structure ($T_{nl}$) and considering that, at each revolution of the table, samples pass in front of the Cr and C targets, resulting in the deposition of 360 individual layers $(N$), the average bi-layer period ($\Lambda$) can be calculated as

$$\Lambda =T_{ml}/N$$

(1)

This approach provides an average estimation of the bi-layer period for the deposited multilayer architecture, without considering the exact contribution of the structural layers.

For the present study, it can be concluded with a sufficient level of significance that, when transitioning from the maximum rate to the minimum Cr deposition rate, the thickness of the Cr layer varies by approximately 0.7 nm, corresponding to a 21.9% contribution to the bi-layer period. This clearly demonstrates the dominant role of the Cr flux in controlling the growth kinetics of the metallic sublayers. Complete isolation of the Cr metal by fully interrupting its deposition is, however, not feasible under the present process. A coating deposited exclusively from the C source neither sustains stable growth nor achieves comparable thickness, indicating that Cr plays a critical role in promoting thinfilm continuity and growth efficiency in the multilayer system.

3.2. Electrochemical Tests

Figure 4a presents selected OCP–time dependences. The attained steady-state values of the OCP, as well as the direction of its deviation, are qualitative indicators of electrochemical interactions between metal surfaces and the corrosive environment. The values of OCP exposure in 3.5 wt.% NaCl are shown in

Table 5. Corrosion parameters extracted from the CV dependences of the multilayer Cr(Cr/C)ml coating on the HS6-5-2 substrate.

Sample	OCP (VAg/AgCl)	Ecorr (VAg/AgCl)	Jcorr (µA·cm–2)	Epitt (VAg/AgCl)	Erp (VAg/AgCl)
S1	−0.367 ± 0.045	−0.348 ± 0.059	2.156 ± 0.914	0.268 ± 0.025	−0.335 ± 0.013
S2	0.173 ± 0.016	−0.138 ± 0.005	0.229 ± 0.025	0.956 ± 0.013	−0.363 ± 0.015
S3	0.176 ± 0.036	−0.105 ± 0.051	0.456 ± 0.170	0.936 ± 0.033	−0.386 ± 0.011
HS6-5-2	−0.407 ± 0.019	−0.429 ± 0.027	17.179 ± 0.424	−0.240 ± 0.100	−0.418 ± 0.008

and Figure 5a. For series S2 and S3, their open-circuit potentials remain nearly unchanged and stable, exhibiting similar OCP values within the range of 0.10 to 0.25 V. These values are approximately 400 mV more positive than those corresponding to the coatings from series S1.

The more negative values and the unstable evolution of the OCP for S1 can be regarded as an indication of inferior corrosion protection. Nevertheless, all three series exhibit more positive OCP values compared to the HS6-5-2 substrate, which classifies them as cathodic coatings.

Selected polarization curves shown in Figure 4b were used to derive the main corrosion–electrochemical parameters such as corrosion potential (E_corr), corrosion current density (J_corr), pitting potential (E_pit), and repassivation potential (E_rp). The electrochemical data obtained with the standard error is summarized in Table 5.

The average, maximum, and minimum values of these parameters for each series of the (Cr/a-C)ml coatings, as well as for the HS6-5-2 substrate, are presented in Figure 5.

The corrosion potential of series S1 shows values of approximately −0.348 V, which are intermediate between those of the HS6-5-2 substrate (−0.430 V) and the E_corr values for series S2 and S3 (−0.138 V and −0.105 V, respectively). The corrosion potentials of series S2 and S3 are approximately 300 mV more negative than the corresponding OCP values. These values are more likely related to minor modifications of the passive films from cathodic polarization than to prolonged exposure to the corrosive environment. This hypothesis was confirmed by additional OCP measurements, which showed no significant changes in potential values over period of 8 h compared to those recorded during the first 10 min.

The corrosion rate, expressed as corrosion current density, showed the highest average value for series S1, although in some measurements, values comparable to those of the other two series were also recorded. This behavior is characteristic of coatings with an inhomogeneous composition and/or structure and a strong dependence of the corrosion process on the type of surface defect present. In this respect, the most stable performance was observed for series S2, for which the corrosion current density remained below 0.3 μA·cm⁻².

The S2 and S3 polarization curves exhibit a profile typical of passive metallic systems, starting with a cathodic branch characterized by diffusion limitations of the oxygen reduction half-reaction, followed by an abrupt reversal of the working electrode polarity and a transition to a low-slope anodic branch. The relatively slow and linear increase in current during anodic polarization from −0.12 V up to nearly 1 V is an indication of the presence of a stable passive layer limiting the diffusion-controlled reaction of anodic metal dissolution.

In contrast, the S1 coatings exhibit unstable behavior, with potential fluctuations recorded at about 200 mV of anodic polarization (vs. E_corr) and at the anodic branch of the curve with a stepwise current increase. The pitting potential, a characteristic of resistance to pitting corrosion, for the S1 samples is reached at approximately 0.27 V. For series S2 and S3, the passive state is only disrupted upon reaching relatively high positive polarization values close to 1 V (Figure 4b). These high E_pit values confirm the good protective properties of the multilayer structure of the (Cr/a-C)ml coatings in series S2 and S3.

At such anodic polarization levels, chloride ions still manage to accumulate locally and hinder the maintenance of the passive state; in isolated defective regions, they penetrate beneath the multilayer structure and initiate substrate corrosion, which is detected as an increase in current.

Upon decreasing anodic polarization (the reverse scan of the CV curves), the current density remains high, at approximately 1 mA·cm⁻². This observation is an indication of (Cr/a-C)ml-layer disruption and the existence of a localized corrosion process into the HS6-5-2 substrate. Such assertion is further supported by the E_rp values, which, for all three series of (Cr/a-C)ml coatings, are very close to the E_corr of HS6-5-2 and fall within the range of −0.407 V to −0.345 V (Figure 4b and Figure 5a).

Electrochemical impedance spectroscopy (EIS) was employed to investigate the protective properties of the (Cr/a-C)ml layers and the evolution of the corrosion process after anodic polarization. Figure 6a–c present Nyquist and Bode plots of representative spectra for each of the investigated series, as well as for the HS6-5-2 substrate before CV tests. The coatings of series S1 exhibit the lowest impedance (Figure 6c), with predominantly resistive behavior manifested as an incomplete semicircle (Figure 6a) and a shift in the phase angle to below 70° in the frequency range of 10³–10² Hz (Figure 6b). Moreover, in the low-frequency range, the interface shows ohmic behavior, and the phase shift for S1 tends to 0°. Such behavior is typical of metallic passive electrodes [33]. A decrease in impedance indicates that the solution has penetrated the coating. The Nyquist plots of the S2 and S3 (Cr/a-C)ml coatings show a steep linear section at low frequencies (Figure 6a). This type of impedance spectrum is characteristic of a blocking electrode with a small deviation from the ideal capacitive behavior [33]. Bode plots demonstrate that the impedance of S2 and S3 is dominated by capacitance even in the low-frequency range, as the phase deviation is still above 70° (Figure 6b).

Corrosion testing under applied anodic polarization significantly impacted the shape of the Nyquist plots and led to similar behavior of all three (Cr/a-C)ml coating series (Figure 6d). Two distinct semicircles are clearly observed: the high-frequency semicircle has a small diameter and corresponds to the response of the protective layers, whereas the low-frequency semicircle characterizes the corrosion process of the substrate. At the lowest frequencies, i.e., at the end of the large semicircle, a tail appears, which can be associated with diffusion limitations during corrosion in less accessible regions such as deep pits or crevices beneath the protective layers [34]. After anodic polarization during the CV tests, the behavior of the metal systems is entirely resistive, with the phase shift not reaching 55° (Figure 6e). This is an indication that the dominant processes at the metal–electrolyte interface are those limited by charge transfer processes and are not limited by diffusion in the barrier layers. However, although the corrosion resistance of all three series is significantly deteriorated, their impedance remains about one order of magnitude higher than that of the HS6-5-2 substrate (Figure 6f).

The equivalent circuit used to simulate the reaction at the metal–electrolyte interface before the corrosion tests is R_s(C_c[R_c(R_ctQ_dl)]) and is presented in Figure 7a. In this circuit, R_s is the solution resistance, C_c is the capacitance of the insulating defect-free coating, and R_c is the transfer of ionic species through the defects in the coating. Interaction at the interface HS6-5-2/(Cr/a-C)ml coating is described with the parallelly connected constant phase element of the electric double layer Q_dl and the charge transfer resistance R_ct.

The behavior of the interface after corrosion destruction of the coating is described by R_s(Q_c[R_p(R_ctQ_dl)]) and is presented in Figure 7b. In it, for simplicity, the entire circuit of Figure 7a is represented as a constant phase element of the coating (Q_c). Parallel to it, resistance (R_p) inside the pits in the multilayer is added, which is the result of difficulties with the transport processes through the electrolyte and loose corrosion products in the pits. In this case, the elements Q_dl and R_ct represent the charge of the electric double layer and the charge transfer resistance, respectively, at the metal–electrolyte interface at the pit’s bottom.

The data calculated by the first equivalent circuits are listed in

Table 6. The data deduced from EIS spectra simulated using the equivalent circuits.

Before cyclic polarization test

Sample

Rs, Ω

Rc, Ω

Rct, MΩ

Qdl, µS·sn·Ω−1

ndl

Cc, nF

χ2

S1

136.1 ± 1.7

259.2 ± 35.0

0.310 ± 0.04

1.69 ± 0.14

0.652 ± 0.013

175.0 ± 13.09

0.0023

S2

131.0 ± 0.9

142.9 ± 15.3

36.15 ± 3.7

1.70 ± 0.24

0.805 ± 0.024

291.0 ± 29.47

0.0002

S3

141.1 ± 2.3

234.1 ± 24.1

40.64 ± 1.6

1.36 ± 0.11

0.785 ± 0.052

319.6 ±35.8

0.0006

HS6-5-2

151.7 ± 1.3

12.1 ± 1.3

0.032 ± 0.03

28.93 ± 0.44

0.65 ± 0.022

283.9 ±2.06

0.0007

After cyclic polarization test

Rs, Ω

Rc, Ω

Rct, kΩ

Qdl, µS·sn·Ω−1

ndl

Qc, µS·sn·Ω−1

nc

χ2

S1

122.7 ± 2.1

1014 ± 52.0

13.3 ±1.7

203.3 ± 11.4

0.824 ± 0.029

100.6 ± 8.7

0.480 ± 0.027

0.0012

S2

137.4 ± 7.1

1035 ± 65.0

15.4 ± 2.6

34.9 ±7.9

0.706 ± 0.032

221.0 ± 13.7

0.942 ± 0.041

0.0009

S3

127.0 ± 8.8

649.5 ± 29.1

6.57 ± 2.10

551.4 ± 18.2

0.937 ± 0.026

124.1 ± 22.9

0.575 ± 0.019

0.0005

HS6-5-2

125.6 ± 3.4

113.6 ± 0.04

0.90 ± 0.12

1366 ± 21.2

0.982 ± 0.041

535.3 ± 32.1

0.979 ± 0.021

0.0001

. The (Cr/a-C)ml coatings have a one-order-of-magnitude lower capacitance than HS6-5-2, which is consistent with the insulating nature of the surface. The lower R_c of the bare substrate can be identified within the natural passive layer on HS6-5-2.

The EIS results for the chemical resistance of the coatings expressed as R_ct confirm those obtained from the OCP and CV tests. Before the polarization tests, the R_ct values for series S2 and S3 are close (36.15 MΩ and 40.64 MΩ, respectively), while the corresponding values for series S1 and the substrate are 0.31 MΩ and 0.032 MΩ. Once the multilayer is broken, the R_ct values for all three series converge and are in the range of 6.57 kΩ to 15.3 kΩ, while the capacitance of the double layer increases significantly.

3.3. Microscopic Observation

The surfaces of the S1, S2, and S3 samples were examined prior to any corrosion tests. Optical microscopy at magnifications of 200x and 500x showed uniformly distributed surface defects (Figure 8a–c). The smaller defects (around or below 10 μm) appear as slight depressions (pinhole), while the larger ones correspond to localized material accumulations formed during deposition. No traces of corrosion products from the HS6-5-2 substrate (colored spots in the orange–brown range, typical of iron compounds) were detected anywhere on the surface, indicating good protective properties of the multilayer coatings under atmospheric conditions. Imperfections observed via optical microscopy are most likely microparticles sputtered during the deposition process.

SEM observations provide significantly less useful information in terms of the defect distribution of the native multilayer. The surfaces appear much more homogeneous compared to optical observations (Figure 9a–c). Isolated defects such as slight depressions or accumulations up to 5 μm in size were seen over large areas. It can be assumed that the corrosion process develops intensively at the substrate–coating interface. Thus, in the substrate with the initial pinhole defects, corrosion products with increased volume accumulate mechanical stress that leads to cracking and detachment of small flakes from the (Cr/a-C)ml coating above them. Such delamination of the multilayer is primarily observed on the corroded surface of series S1.

Various localized surface damages were observed on all samples after corrosion testing (Figure 9d–f). Corrosion damages in the samples of all three series appeared primarily as deep pits with relatively regular circular mouths with diameters of approximately 50–150 μm. Irregularly shaped and larger pits were less common in S2 and S3 samples, but they resulted in detachment of the (Cr/a-C)ml coating in the S1 population or occasional coalescence of neighboring pits.

3.4. EDX Analysis

Table 7. The results (in wt. %) of EDX analysis for (Cr/a-C)ml coatings of S1, S2, S3, and HS6-5-2 before and after corrosion tests.

Sample		Fe	V	Cr	C	W	Mo	Others	Cr/C
HS6-5-2		77.75	1.34	4.45	3.87	6.22	4.96	Si 0.20; S 0.07;	1.15
S1	before test	3.87	–	89.32	6.81	–	–	–	13.12
	uncorroded	4.11	–	89.12	9.72	–	–	–	9.17
	corroded	–	3.45	8.04	5.22	6.58	7.21	O 18.95; S 11.82; Si 0.98;	1.46
	deposits	–	7.04	5.01	9.21	31.52	19.67	O 3.07	0.54
S2	before test	9.15	–	66.09	22.68	0.48	0.43	O 1.19; Si 0.10	2.94
	uncorroded	8.61	–	65.48	24.34	1.33	1.62	O 1.58	2.69
	corroded	18.49	5.08	8.56	11.02	20.33	16.99	–	0.78
S3	before test	10.08	0.08	59.71	28.46	0.13	0.35	Si 0.12; O 1.08	2.10
	uncorroded	9.76	0.04	60.18	28.39	0.04	0.31	Si 0.08; O 1.21	2.12
	corroded	8.44	0.09	57.82	30.08	0.39	0.34	Si 0.36; O 1.68	1.92
	deposits	6.72	–	65.01	22.91	–	–	Si 0.73; O 3.98	2.84

presents the weight percentages of the main elements measured on the (Cr/a-C)ml coatings on HS6-5-2 before and after corrosion tests in 3.5% NaCl. Due to the localized nature of the corrosion damage, analyses were performed in three specific zones: an uncorroded surface (1), areas with deposited corrosion products (2), and exposed surfaces of shallow pits (3).

Analysis of the coatings prior to corrosion testing shows an increase in Fe content in the order from S1 to S3, which most likely correlates to decreasing coating thickness. The decrease in Cr content and the corresponding increase in C content along the same series are associated with the deposition conditions. This also explains the significant differences in the Cr:C ratio, which ranged from 9.7 for S1, to 2.9 for S2, and to 2.1 for S3. For comparison, the Cr:C ratio in the substrate is 1.2.

EDX analyses confirmed that coating compositions in non-corroded areas remained the same as compositions of untreated surfaces. The Cr:C ratio for series S1 and S2 decreased slightly, with average values of 9.2 and 2.7, respectively, while for series S3, it remained unchanged at 2.1.

The contents of V, W, and Mo increased significantly in the areas with deposits, (by 3.25–5 times) compared to the HS6-5-2 substrate. The most pronounced increase was observed in series S1 (4–5.25 times), while the Cr content remained close to that of the substrate.

Significant enrichment of the main alloying elements (V, Cr, W, and Mo) and high oxygen content were confirmed in the most severely attacked zones where the (Cr/a-C)ml coating is completely destroyed. Consequently, Fe was preferentially dissolved, leading to the formation of an oxide layer composed of elements more resistant to anodic polarization. In some deeper regions, elevated sulfur content was detected in combination with Mo (series S1). It could be assumed that the presence of non-metallic MoS inclusions in the substrate compromises the adhesion of the (Cr/a-C)ml layer. These sites most likely act as local concentrators of the dissolution corrosion process rather than as accumulations formed during corrosion. Sulfur traces are relatively rare and do not play a dominant role in the progression of corrosion processes on the (Cr/a-C)ml coating of HS6-5-2.

Based on the results of the three corrosion tests, microscopic observations, and EDX analysis, it is evident that all three coatings exhibit better corrosion resistance compared to the uncoated HS6-5-2 substrate. The S1 series samples, which had the highest rate of deposition, show inferior corrosion resistance compared to the S2 and S3 samples. Among the three series, this multilayer coating exhibits the most negative values for characteristic potentials (OCP, E_corr, and E_pit), the highest corrosion current density, and the lowest impedance.

This result may seem somewhat surprising, considering that the S1 coatings have the highest chromium content and the greatest thickness. However, based on the microscopic observations and, above all, the observed undermining of the coating and peeling of the flakes from the steel surface, it can be concluded that the higher-chromium structure is more stressed. The redistribution of internal stresses to the coating/substrate interface and to surface imperfections most likely initiates pits.

4. Conclusions

A variety of electrochemical methods was used to investigate the influence of the deposition conditions on the corrosion resistance of multilayered chromium/amorphous carbon Cr/(Cr-aC)ml coatings on the HS6-5-2 substrate. The correlation of electrochemical results, microscopic surface observations, and elemental composition analysis demonstrated that

• The Cr/(Cr-aC)ml coatings exhibit cathodic behavior relative to the HS6-5-2 substrate in a neutral chloride environment (3.5% NaCl).
• When deposited at a lower rate, coatings S2 and S3 have higher carbon content and Cr:C ratios of 2.9 and 2.1, respectively. These series demonstrate significantly higher corrosion resistance in 3.5% NaCl compared to series S1 for which this ratio is 13.1.
• The S2 and S3 coatings exhibited corrosion potential values approximately 0.3 V more positive and corrosion current densities about ten times lower than the corresponding values for the S1 coating. Under anodic polarization, all three coatings develop pitting corrosion, with S2 having the most positive pitting potential of 0.96 V_Ag/AgCl.
• Once the integrity of the multilayer coatings was compromised, the behavior of all three systems became identical.
• The reduced corrosion resistance of the coating with the highest chromium content (89 wt.%) is attributed to higher internal stresses within the layer, which promote the accumulation of defects permeable to the corrosive environment and reduce adhesion at the interface with the HS6-5-2 substrate.

Multilayered nanostructured Cr-C coatings are capable of providing the desired corrosion protection of HS6-5-2 steel when properly designed and deposited. Improving the corrosion resistance of nanostructured multilayer coatings can be achieved by coating deposition process optimization and via close control of the chromium growth rate and coating composition.