An Electrochemical Investigation of the Pitting Corrosion of TZM Alloy in Chloride Solution

Abstract

In this study, cyclic polarization (CP) measurements were conducted on the molybdenum-based titanium–zirconium–molybdenum (TZM) alloy in 3.5% NaCl solutions under varying pH conditions, and the results were compared with those of pure molybdenum. No passivity breakdown was observed during cyclic polarization in acidic and neutral chloride solutions. The surface film formed on the TZM, and pure Mo samples displayed a dual-layered structure, comprising an inner layer of p-type semiconductivity and an outer layer of n-type semiconductivity. The defect density of the n-type layer ranged from 7.5 × 10¹⁷ to 7.5 × 10¹⁹ cm⁻³, while the p-type layer had a carrier density ranging from 2 × 10¹⁸ to 9 × 10¹⁹ cm⁻³. The pure molybdenum samples demonstrated lower passive current densities, lower charge carrier densities, and higher impedance than the TZM alloy. The lower corrosion resistance of TZM alloy could be attributed to the higher dislocation density, which acted as short-circuit paths for Mo diffusion, and the presence of carbides that exhibited a microgalvanic effect. Overall, this study clarified that the localized corrosion reported in the literature was not due to the breakdown of the passive layer but may be linked to the heterogeneous microstructure.

1. Introduction

Molybdenum is used as an alloying addition for high-temperature applications and as a refractory metal due to its high melting point (2623 °C) and high elastic modulus (353.3 GPa) [1]. Molybdenum is also an essential alloying addition in stainless alloys, such as duplex stainless steels and nickel-based alloys, for enhancing pitting corrosion resistance [2]. TZM is a widely used molybdenum alloy strengthened by the addition of small quantities of titanium (0.5 wt%), zirconium (0.08 wt%), and carbon (0.02 wt%). The atomic radii of Mo, Ti, Zr, and C are 145, 140, 155, and 70 picometers, respectively [3]. Ti forms a continuous solid solution in Mo above its hcp to bcc transition temperature (882 °C), while the maximum solid solubility of Ti in Mo at low temperatures is about 13 at% [4]. The higher solid solubility of Ti in Mo may be attributed to the smaller size of the Ti atom (−3.45%) compared to the Mo atom. Zr exhibits limited solid solubility in Mo. The terminal solid solubility is ~10 at% at 1880 °C, and the solubility may be less than 0.5 at% at room temperature [4]. The equilibrium binary phase diagram indicates the formation of Mo₂Zr-type Laves phase with a cubic symmetry and space group of Fd-3m [5]. Ti has little solubility in the Laves phase Mo₂Zr. However, Ti and Zr have continuous solid solubility in low-temperature hcp and high-temperature bcc structures. Carbon has minimal solid solubility in Mo, Ti, Zr, and their solid solution combinations [5]. Titanium and zirconium have a strong affinity for carbon and form fcc-structured TiC, ZrC, and (Ti, Zr)C carbides (space group: Fm-3m). These MC carbides render a high creep strength by pinning the grain boundary sliding, subgrain migration, and dislocation movements [6]. Carbon could combine with Mo and form three types of carbides, such as hcp Mo₂C (space group: P6₃/mmc), hexagonal MoC (space group: P6-m2), and η-MoC (space group P6₃/mmc) [5]. The formation of metastable Mo₂C along the grain boundaries of TZM alloys has been observed, despite MC-type carbides having a higher thermodynamic driving force [5,6]. This behavior is attributed to the sluggish diffusion kinetics of Ti and Zr, while carbon has a lower kinetic barrier and quickly diffuses in the bcc Mo matrix and forms Mo₂C [7]. In addition to TiC, ZrC, and Mo₂C, oxides of Zr and Ti were also reported on the grain boundaries of TZM [4]. The formation of oxide depends on the manufacturing process of the TZM alloy. Oxides are not reported in the vacuum-melted alloys, whereas the powder metallurgy route reveals the presence of oxide particles, such as MoO_xTiO_z and ZrO_x [4]. It is noted that the presence of secondary phases could compromise corrosion resistance due to the microgalvanic effect. When the chemical potentials of the alloy matrix and secondary phases are different, the phase having a lower chemical potential is preferentially dissolved by corrosion [8]. The chemical potential difference is attributed to the preferential partitioning of the alloying elements to the particular phases [9]. Microgalvanic corrosion typically initiates at the phase boundaries due to the difference in surface potentials [10]. The transition-metal carbides generally exhibit higher galvanic potentials than the alloy matrix, leading to microgalvanic corrosion attack [11]. The carbide phases present in the TZM alloy could induce localized attack at the matrix–carbide interface due to the microgalvanic effect.

Maday et al. [12] investigated the surface film properties of TZM using X-ray photoelectron spectroscopy (XPS) and correlated them with the electrochemical behavior. The TZM samples were immersed in oxygenated and hydrogenated water at 250 °C. A less protective duplex layer comprising inner MoO₂ and outer MoO(OH)₂ phases was observed in oxygenated conditions. An inner tetravalent Mo oxide/hydroxide and an outer hexavalent Mo oxide were observed in hydrogenated conditions, which was more protective [12]. Hu et al. [13,14] reported the corrosion behavior of TZM alloy in different chloride-containing media as a function of chloride concentrations in the range of 0.5–1.5 mol/L [13] and as a function of pH in 3.5% Cl⁻ [8]. These authors reported the destruction of passivity by chloride ions [14] and the formation of pitting around second-phase particles [13]. Ge et al. [15] investigated the effect of the secondary phase on the corrosion behavior of TZM by scanning Kelvin probe force microscopy and showed that the potential of a secondary phase particle was 73.4 mV more positive than the adjacent matrix. This potential difference led to galvanic coupling and the initiation of pits at the particle/matrix interface [15]. Tuncay and Ozyurek [16,17] investigated the role of alloying additions of Ti and Zr on the microstructure and corrosion properties of TZM alloys in dilute chloride (0.0012 g NaCl/L) solutions. Increasing the Ti content from 0.4 wt% to 0.55% increased the corrosion resistance, which was attributed to a decrease in the Mo₂C carbide content by fixing the carbon with Ti [16]. These authors also reported pitting at the regions where discontinuity in the passive film occurred. The increase in Zr alloying addition from 0.06 wt% to 0.09 wt% increased the corrosion resistance [17]. The maximum solid solubility of Zr in Mo at room temperature was about 0.1%. If the Zr content was higher than 0.1%, it resulted in the precipitation of ZrO_x particles, which affected the corrosion behavior [17]. Pitting along grain boundaries was reported in TZM alloys with 0.06 wt% Zr. Overall, several investigators reported pitting in TZM alloys in chloride-containing solutions.

It is well established that the enhanced pitting corrosion resistance of stainless steels and nickel base alloys is attributed to the alloying addition of molybdenum. On the other hand, molybdenum-based TZM alloy, which contains more than 98 wt% Mo, is reportedly prone to pitting corrosion in dilute chloride solutions at room temperature [13,14,15,16,17]. This observation questions the actual role of Mo as an alloying element in rendering resistance to localized corrosion. It could be argued that the observed pitting resistance of the stainless alloys was due to the formation of a Cr-rich oxide film that was doped with molybdenum ions, rather than just molybdenum [18]. However, several studies have pointed out the preferential enrichment of elemental or partially oxidized Mo at the actively corroding surface, and this enrichment reduces the kinetics of dissolution [19,20,21]. Furthermore, it is well established that a predominantly Mo-containing nickel base alloy, HASTELLOY B-3 alloy (UNS N10675), exhibits superior resistance to HCl, with a maximum Cr content of only 1.5 wt%.

This work uses cyclic polarization studies to investigate whether pitting occurs on the TZM alloy in 3.5% NaCl solutions under different pH conditions at room temperature. The 3.5% NaCl solution closely simulates the salt content of seawater, but the chloride concentration is much higher than that observed in the physiological condition [22]. The acidic pH condition simulates the environment typically observed at the tips of the pits or cracks. In contrast, the alkaline condition could represent the environment that exists on the exterior surface of the pit or crack wall, where the cathodic reaction occurs [23]. Furthermore, the 3.5% NaCl could potentially be used for electrochemical micromachining of Mo alloys [24]. Most investigations on the pitting of molybdenum alloys reported in the literature consisted of only single polarization runs. Conducting cyclic polarization helps understand the initiation and growth behavior of pits from the hysteresis loops. The semiconducting properties of the surface layer formed on the TZM are analyzed using electrochemical impedance spectroscopy and Mott–Schottky plots. The surface films are studied using confocal Raman microscopy. The results are compared with the pure molybdenum samples to understand the role of alloying elements on the corrosion behavior.

2. Experiments

Metallographic samples of 10 mm × 10 mm were cut from 0.25-inch-thick TZM and pure Mo sheets (H.C. Starck Solutions, Newton, MA, USA) using a diamond wheel cutter. One surface of the sample was soldered to an insulated copper wire for electrical connection, and then it was cold-mounted using an acrylic resin (Pace Technologies, Tucson, AZ, USA) by exposing only one surface. The exposed surface was metallographically polished down to 1200 grit. The polished samples were taken for electrochemical testing. The mounted samples were ground using 120, 240, 400, 600, 800, and 1200 grit SiC papers (Allied High Tech Products, Inc, Cerritos, CA, USA) for optical microscopy. Each stage of the grinding process would remove the scratches induced from the previous stage by rotating the sample surface perpendicular to the grinding direction. Once grinding was complete with the 1200 grit SiC paper, polishing the samples on a cloth polishing pad was conducted. For polishing, 1 µm alumina with 3 µm suspended diamond particles was used as the first stage, followed by 1 µm alumina with 1 µm suspended diamond particles. The last polishing stage was 0.3 µm alumina with 0.1 µm suspended diamond particles. After polishing, the samples were cleaned with deionized water and soap to remove any impurities. After air drying, the samples were etched using Murakami’s etchant (10 g of K₃Fe(CN)₆ + 10 g KOH dissolved in 100 mL of water). The etching time was ~1 min and was carried out by pipetting the etchant onto the surface of the sample and swabbing with cotton. The etchant was removed with 3% hydrogen peroxide and cleaned with methanol. Microstructural images were taken using an Olympus PMG-3 optical microscope. Grain size measurements were performed by an intercept method by following the ASTM standard E112 [25], and the average grain sizes of the samples were reported.

For transmission electron microscopy (TEM) sample preparation, samples were ground down and polished to ~200 μm thick. Discs with a 3 mm diameter were punched out and further thinned down in a Fischione Instruments Twin-jet electropolisher model 110. The chemical etchant was 13% sulfuric acid solution to 87% methanol, by volume. Dry ice was used to lower the temperature of the mixture to −40 °C to reduce the evaporation of the sulfuric acid solution during electropolishing. The samples were jet-polished until a small hole formed; the samples were then characterized soon thereafter to reduce oxide formation on the surface of the samples. A PIPS II model 695 argon ion mill (Gatan, Inc, Warrendale, PA, USA) was used to further thin and remove contaminants from the sample surface. TEM was performed using a JEOL 2010J (EOL USA, Inc., Peabody, MA, USA), at an accelerating voltage of 200 kV. A LECO LM100 Vickers microhardness tester (LECO Scientific Instruments, St. Joseph, MI, USA) was used to conduct a hardness analysis of the samples with a load of 500 g. More than twenty readings were taken for each sample, and the average values were reported.

Open-circuit potential (OCP), linear polarization, cyclic polarization, electrochemical impedance spectra (EIS), potentiostatic, and Mott–Schottky tests were performed on TZM samples. Tests were conducted using a potentiostat (Gamry Interface 1000, Gamry Instruments Inc., Warminster, PA, USA). The test solutions were 3.5% NaCl + 0.1 M HCl (referred to as pH 1 solution; the measured pH was 1.14), 3.5% NaCl without any other addition (referred to as pH 5.5 solution; the measured pH was 5.33), and 3.5% NaCl + 0.1 M NaOH (referred to as pH 13 solution; the measured pH was 12.7). The electrolyte was open to the atmosphere and not purged with nitrogen or other inert gases. Therefore, the pH of the neutral solution was low due to the dissolved carbon dioxide from the atmosphere, which was converted to carbonic acid by the reaction: CO₂ + H₂O → HCO₃⁻ + H₃O⁺. Reference electrodes were prepared with silver wire anodized in a saturated KCl solution, creating a silver chloride coating. The agar gel prepared by cooking in a saturated KCl solution was inserted into the end of an eyedropper, with the saturated KCl poured over, and the silver chloride-coated silver wire went through both. A platinum spiral counter electrode was prepared by inserting a platinum wire into an eyedropper and twisting the end into concentric circles.

OCP tests were conducted for 1800 s to equilibrate the sample after immersion in the test solution. After OCP, polarization resistance testing was conducted to measure the polarization resistance. The parameters for these tests were ±25 mV versus OCP and a scan rate of 0.1667 mV/s. Cyclic polarization tests were performed at a potential scan rate of 0.1667 mV/s by scanning the potential from −0.25 V versus OCP in the positive direction to an apex potential of 1.6 V_Ag/AgCl. The forward scan was reversed when the current density reached 5 mA/cm² or the apex potential was reached. During the forward scan, the potential at which the current increases monotonically is referred to as the transpassive potential (E_TP). The rise in current could signify either a breakdown of the passive film or an oxygen evolution reaction, depending on the potential. If the passive film breakdown occurred during the forward scan and pits were initiated, then the reverse scan would not trace the current profile of the forward scan, but it would cross the forward scan at a lower potential than the transpassive potential. A ‘positive’ hysteresis would be recorded because of the increased time for the pits to repassivate. The crossover potential of the reverse scan is referred to as the pitting protection potential (E_PP).

The passivation region of each sample was used as the range for test potentials in the following tests. Electrochemical impedance spectroscopy (EIS) was then performed in combination with potentiostatic tests to investigate the impedance due to the passivation behavior of the sample surface. First, an EIS test was performed on a sample with an initial frequency of 1000 Hz, a final frequency of 0.01 Hz, and a scan rate of 5 points/decade. The potential at which each sample was tested varied based on the passivation region in cyclic polarization tests and the chemical solution in which the sample was immersed. Afterward, a potentiostatic test was performed for 30 min, at a potential matching the earlier EIS test. This was followed by another EIS test using the same values as the previous EIS. Finally, Mott–Schottky tests were performed to provide information concerning the semiconductor behavior on the passivated surface. Parameters used were an initial voltage of 0.1, a final voltage of −0.5, a voltage step of 0.02, an AC voltage of 15, and a frequency of 100 Hz. All the tests were performed twice using three different samples. The reported data are an average of six test runs. Only representative plots are included in this publication. All the potentials are reported with reference to the Ag/AgCl reference electrode.

Raman spectroscopy was conducted on the corrosion-tested samples using a confocal Raman microscope (HORIBA XploRA Plus, Horiba Ltd., Kyoto, Japan), equipped with a 532 nm, 100 mW, Class 3B internal laser, and a CCD detector (Model: Syncerity). The surface spectra of the corrosion-tested specimens were acquired using an x100 VIS DF objective, 600 (750 nm) grating, and a 50% laser filter. The confocal hole and slit were maintained at default values of 100 µm and 50 µm, respectively, during the acquisition. The spectrum acquisition and analysis were performed using LabSpec6 software (version 6.7.1.10), facilitating data visualization and peak identification. The identified peaks were compared with references in the literature to assign possible vibrational functional groups in the corrosion layer. Based on these comparisons, potential compounds corresponding to the respective peaks were assigned using the closest possible peak matches.

3. Results and Discussion

TZM alloy consisted of 0.08 wt% Zr and 0.5 wt% Ti as the primary alloying additions, along with 0.02 wt% C. Trace amounts of N, O, Fe, Ni, and Si were also present but were not evaluated. The oxygen content could be approximately 0.04 wt%, which affected the formation of secondary phases as Mo-Ti-Zr-O compounds. Figure 1a,b show the optical microstructures of the exposed surfaces of the pure Mo and TZM samples, respectively. The pure Mo sample exhibited a relatively equiaxed fine-grained structure. The TZM sample revealed an elongated grain structure in the rolling direction.

Table 1. Average grain size and Vickers microhardness of the samples.

Alloys	Orientation	Ave. Grain Size (µm)	Aspect Ratio	VHN
Pure Mo	L-ST	50 ± 4	2.1	233 ± 8
TZM	L-ST	43 ± 17	9.8	279 ± 15

summarizes the results of grain size and the Vickers hardness of the samples. The Mo samples exhibited a hardness value of approximately 230 kg/mm², a typical value widely reported in the literature. The average grain size of the Mo was 50 ± 4 μm, with an aspect ratio of 2.1. The average grain size of TZM samples was 43 ± 17 μm, with a high aspect ratio of 9.8, indicating that the rolled microstructure had not fully recrystallized to yield an equiaxed grain structure. The hardness of TZM was higher than that of pure Mo due to a finer and elongated grain structure.

The presence of fine TiC and ZrC carbide particles (not resolved in the optical microstructures) and the possible presence of Mo₂C particles along the grain boundaries could be attributed to the finer grain size, which in turn increased the hardness. Figure 2a,b display the high-resolution microstructures of the Mo and TZM samples obtained using transmission electron microscopy. Cellular dislocations entangled in the form of subgrain structures and cell walls were observed in both samples. The Mo sample exhibited an almost equiaxed subgrain structure, while the TZM sample displayed elongated subgrain structures. The dislocation density appeared higher in the TZM than in the Mo sample. The darkly shaded particle-like structure on the left side of Figure 2b could be a carbide particle; however, neither elemental analysis nor selected area electron diffraction data were obtained to support this.

Figure 3a shows the cyclic polarization results in pH 1 solution. The E_corr values of Mo and TZM were −30 and −235 mV_Ag/AgCl, respectively. The more active corrosion potential of TZM could be attributed to the enhanced binding of hydrogen at dislocation sites and carbide particles. It is well known that dislocations and secondary phase particles act as trapping sites for hydrogen. These traps could be reversible or irreversible, depending on the binding energies with hydrogen. Recently, Faye and Szpunar [26] showed that TiC had a higher trapping strength for hydrogen than other carbides, such as NbC and VC. The hydrogen adsorption may have occurred during immersion in the acid solution and during cathodic polarization. The presence of dissolved oxygen (no deaeration was performed) in the solution and no significant hydrogen trapping on the Mo samples drove the E_corr to a nobler potential. Since the corrosion potential was high, the Mo sample exhibited only a narrow window of passivation potential. The median passive current densities of the Mo and TZM samples were 1.1 × 10⁻⁶ and 6.4 × 10⁻⁶ A/cm², respectively. The recorded passivation current density of the Mo sample was lower than that of TZM.

However, the passive current density could not be directly compared to demonstrate the ability to maintain the passive condition. The anodic current recorded during the polarization had two components: (i) the anodic dissolution of Mo and/or oxidation of low valent surface oxide; and (ii) the oxidation of the adsorbed hydrogen atoms on the Mo surface. Falkenberg et al. [27] demonstrated that the observed current was mainly due to a diffusion-controlled oxidation of hydrogen absorbed in the metal during the cathodic activation. Therefore, the higher passive current density observed in the TZM could be attributed to the higher amount of hydrogen adsorbed, which also shifted the E_corr to a more active potential. The transpassive potential at which the current increased sharply could not be determined from the plots shown in Figure 3a because the passivation current varied across the entire potential region. The transpassive potential was determined by intersecting two tangents of the forward scan of the polarization curve. The extrapolated transpassive potentials of the Mo and TZM samples were 140 and 180 mV_Ag/AgCl, respectively. During the reverse scan, the current decreased significantly just above the transpassive potential, resulting in negative hysteresis. This implied that there was no localized breakdown of the passive film. During the transpassive condition, the low-valency oxide film could have oxidized further without any breakdown. During the reverse scan, the highly oxidized film acted as a barrier, slowing down further oxidation. This process resulted in a decrease in the current density during the reverse scan. When the potential decreased below 0.2 V, a cathodic loop was observed. This observation indicated that a barrier film was intact on the samples, which increased the corrosion potential to 0.2 V Ag/AgCl. Therefore, scanning the potential below 0.2 V showed cathodic polarization behavior. The cyclic polarization of the Mo and TZM samples revealed that no localized corrosion occurred on the samples due to the breakdown of the passive film. The passive film remained intact, as indicated by the cathodic polarization loop above the transpassive potentials.

Figure 3b shows the cyclic polarization (CP) behaviors of the Mo and TZM samples in the pH 5.5 solution. The E_corr values of the Mo and TZM were −244 and −456 mV_Ag/AgCl, respectively, and the median passive current densities were 3.7 × 10⁻⁶ and 3.4 × 10⁻⁵ A/cm², respectively. The extrapolated transpassive potentials were 80 and 160 mV_Ag/AgCl, respectively. The reverse scan did not reveal a hysteresis loop associated with passivity breakdown or pitting corrosion in this pH condition. A cathodic loop was observed closer to the transpassive potential of the Mo sample. In contrast, a much lower potential was required to attain a cathodic polarization during the reverse scan for the TZM sample. This observation indicated that the Mo sample formed a more stable barrier film than TZM, resulting in a more noble equilibrium potential. The passive current density during the reverse scan of the TZM was higher than the passive current density of Mo during the forward sweep. Therefore, it is clear that the Mo formed a more stable passive layer than the TZM.

Figure 3c shows the CP plots of the Mo and TZM in pH 13 solution. The E_corr values were −517 and −494 mV_Ag/AgCl for the Mo and TZM samples, respectively. The anodic polarization increased the current density monotonically until −0.2 V_Ag/AgCl, beyond which a limiting current behavior was observed. The reverse scan traced the forward scan current profile, indicating no passive film growth occurred during polarization.

Table 2. Electrochemical parameters derived from polarization measurements.

Sample	Cathodic Exchange Current Density, i0c, A/cm2	Cathodic Tafel Slope, V/Decade	Anodic Exchange Current Density, i0a, A/cm2	Anodic Tafel Slope, V/Decade	Corrosion Current Density, icorr, A/cm2	Corrosion Potential, Ecorr, VAg/AgCl
3.5% NaCl + 0.1 M HCl (pH: 1)
Mo	3 ± 2 × 10−21	0.12 ± 0.04	6 ± 5 × 10−10	0.22 ± 0.04	0.6 ± 0.4 × 10−6	−0.14 ± 0.13
TZM	6 ± 2.3 × 10−13	0.19 ± 0.02	2 ± 8 × 10−8	0.48 ± 0.06	6 ± 4 × 10−6	−0.23 ± 0.015
3.5% NaCl (pH: 5.5)
Mo	3.2 × 10−17	0.2 ± 0.09	5 ± 3 × 10−10	0.15 ± 0.03	0.6 ± 0.4 × 10−7	−0.26 ± 0.04
TZM	1.1 × 10−10	0.25 ± 0.04	2 ± 2 × 10−6	0.6 ± 0.06	0.8 ± 0.2 × 10−6	−0.41 ± 0.04
3.5% NaCl + 0.1 M NaOH
Mo	3 ± 1 × 10−8	0.32 ± 0.03	6 ± 1 × 10−15	0.08 ± 0.01	6.5 ± 1.3 × 10−6	−0.52 ± 0.005
TZM	5 ± 0.5 × 10−8	0.35 ± 0.02	8 ± 1 × 10−15	0.06 ± 0.02	7 ± 1 × 10−6	−0.51 ± 0.006

lists the electrochemical parameters derived from the polarization plots. Overall, the cyclic polarization behaviors of Mo and TZM unambiguously revealed that no localized breakdown of the passive film and associated pitting corrosion occurred in either acidic or neutral chloride solutions. The polarization behavior in the alkaline chloride solution did not show the formation of a barrier film on the Mo and TZM samples. Therefore, pitting due to passivity breakdown may not be relevant in alkaline chloride solutions.

The exchange current density for the anodic reaction was calculated based on the following electrochemical reactions, depending on the pH conditions [28]:

Acid:        Mo → Mo³⁺ + 3e⁻;                     E = −0.2 + 0.0197 log [Mo³⁺] V_SHE

(1)

Acid:        MO³⁺ + 3H₂O → MoO₃ + 6H⁺ + 3e⁻;     E = −0.623 − 0.1182pH − 0.0197 Log [Mo³⁺] V_SHE

(2)

Neutral: MO³⁺ + 2H₂O → MoO₂ + 4H⁺ + e⁻;       E = 0.311 − 0.2364pH − 0.059 Log [Mo³⁺] V_SHE

(3)

Base:    Mo + 4H₂O →MoO₄²⁻ + 8H⁺ + 6e⁻;      E = 0.154 − 0.0788pH + 0.0098 Log [MoO₄²⁻] V_SHE

(4)

The linear portion of the anodic polarization plot was extrapolated until it intercepted the redox potential of the representative anodic reaction. The corresponding current value was reported as the anodic exchange current density [29]. A concentration of 10⁻⁶ mole/l was assumed for the dissolved species [28]. The exchange current density of the cathodic reaction was determined by extrapolating the linear portion of the cathodic polarization plot until it intercepted the redox potential of the oxygen reduction reaction. Since the electrolyte was not deaerated, the oxygen reduction reaction was considered the predominant cathodic reaction. The hydrogen reduction reaction was not considered for the exchange current density values. The corresponding cathodic reaction is given by [28]:

Acid:    4H⁺ + O₂ + 4e⁻ → 2H₂O;                E = 1.23 − 0.059 pH V_SHE

(5)

Alkaline: 2H₂O + O₂ + 4e⁻ → 4OH⁻;              E = 0.4 − 0.059 log(a_OH⁻) V_SHE

(6)

The corrosion process of pure Mo and TZM alloy could be described by the anodic processes as given by the reactions (1)–(4) and the cathodic reaction of oxygen reduction [28]. In addition to these reactions, the corrosion of TZM alloy involves microgalvanic-induced corrosion due to the difference in the chemical potentials of carbides and the matrix. The exchange current densities reported in Table 2 do not capture the microgalvanic effect. Therefore, as an anonymous reviewer of this paper pointed out, no solid conclusion could be drawn from these results. Further studies are required.

Figure 4 shows the current density versus time (I-t) plots during the potentiostatic (PS) conditioning of the TZM samples in different pH conditions. As described in the experimental section, the applied potential corresponded to the median potential in the passivation range of the pH solution. The current density was the highest in the alkaline chloride solution and lowest in the neutral solution. A decay behavior was observed, indicating the possible formation and growth of a surface layer. The current decay behavior could be expressed as:

I(t) = I₀tⁿ

(7)

where I₀ is the instantaneous current recorded at the time of potential application (t = 0), t is the time, and n is the exponent of the time or the slope of the log I(t) versus log time plot.

Table 3. Summary of the potentiostatic test results.

	PH 1		Ph 5.5		ph 13
	TZM	Pure Mo	TZM	Pure mo	TZM	Pure Mo
Iinitial (A) (t = 0)	8.1 × 10−5	2.4 × 10−5	6.4 × 10−5	1.7 × 10−5	4.5 × 10−3	3.8 × 10−3
Ifinal (A) (t = 1800 s)	7.8 × 10−5	2.0 × 10−5	2.6 × 10−5	1.3 × 10−5	3.1 × 10−3	1.5 × 10−3
Slope (n)	−0.1	−0.29	−0.23	−0.45	−0.07	−0.23

summarizes the results of the potentiostatic passivation plots of the TZM and Mo samples. Both samples showed current decay behaviors. The slope n corresponds to the current decay kinetics during the initial period (t < 100 s). It is noted that Mo exhibited faster kinetics in film formation than TZM. Overall, the current density was higher in the TZM than in the Mo in all the pH solutions under the potentiostatic conditions.

Figure 5a,b show the EIS results of Mo and TZM after the potentiostatic (PS) passivation condition in an acidic chloride solution as Nyquist and Bode plots, respectively. The EIS tests were performed before and after the PS tests. The impedance results after the PS tests were consistently higher than those before the PS tests. Therefore, only the EIS results after the PS test are presented for comparison. Both samples exhibited a single time constant, as shown in Figure 5b. The impedance of TZM was much lower than that of Mo at higher frequencies, but it increased rapidly at lower frequencies. The impedance modulus of TZM at low frequencies was almost equal to that of Mo. This can be attributed to the higher interfacial capacitance of the TZM sample compared to that of Mo. Higher capacitance may be associated with increased surface roughness and/or a narrow space charge layer and a defective surface layer.

Figure 6a,b show the Nyquist and Bode plots in neutral chloride solution. In that case, the Mo sample showed marginally higher impedance than TMZ at lower and higher frequencies. The impedance values were similar at intermediate frequencies for both samples. A single time constant was observed in this solution, also implying that the electrochemical interface can be modeled using a single RC-loop electrical equivalent circuit, also known as Randle’s circuit. However, a close observation revealed that a diffusion-controlled impedance could occur at much lower frequencies (f < 0.025 Hz), but the measurement stopped at 0.01 Hz. The Warburg-type behavior indicated that the surface layer formed on the Mo sample could be of a diffusion-limiting type, providing enhanced protection. The impedance of the surface layer formed in Mo was an order of magnitude higher in the neutral chloride solution than in the acidic solution.

Figure 7a,b show the Nyquist and Bode plots in the alkaline chloride solution, respectively. Interestingly, the impedance modulus in the alkaline condition was an order of magnitude lower than that in the acidic condition. The EIS results supported the CP results, showing no protective surface layer formed in the alkaline solution. The TZM sample exhibited a single time constant, with a possible adsorption loop observed at low frequencies. Therefore, the potential surface layer formed on the TZM in the alkaline solution could be associated with the adsorption of OH⁻ and Cl⁻ ions that exhibited current decay behavior during the potentiostatic tests. The Mo sample revealed two time constants, including a Warburg-like component. Since the impedance was low, the diffusion-limited behavior may be associated with a defective or porous surface layer, reprecipitated on the surface due to local supersaturation of the dissolved species.

The Mott–Schottky results for TZM are presented in Figure 8. The Mo samples also exhibited similar trends under each pH condition. Charge carrier densities of the surface layer were calculated from the slopes of the Mott–Schottky (M-S) plots and summarized in

Table 4. Average charge carrier densities (cm⁻³) of the surface layers of TZM and Mo samples.

	Charge Carrier Density, Number of Carriers Per cm3
	pH 1		ph 5.5		ph 13
	TZM	Pure Mo	TZM	Pure mo	TZM	Pure Mo
n-type	2.9 × 1019	7.5 × 1017	7.4 × 1019	4.8 × 1019	5.5 × 1019	2.3 × 1019
p-type	0	0	2.0 × 1018	2.5 × 1018	9.1 × 1019	1.5 × 1019

. The positive slope of the M-S plots indicated n-type semiconductivity, implying that oxygen vacancies were the majority charge carriers. Here, cation interstitials were ignored because Mo cations would be too large to be accommodated as a defect in the lattice. The negative slope in the M-S plot indicated p-type semiconductivity, implying that cation vacancies were the major charge carriers. Since both n-type and p-type behaviors were observed in neutral and alkaline conditions, it was presumed that a dual-layered surface film formed on the samples. The charge carrier density of the film formed on the Mo sample was lower than that of TZM in all the pH conditions, which aligned with the EIS and CP results. In most cases, the charge carrier density of the p-type film was lower than that of the n-type film. The n-type charge carrier density was in the order of 10¹⁹ cm⁻³ (with one exception), indicating that the film could be described as MoO_3-x, with x~0.001–0.001. For the p-type layer, the film could be described as Mo_1-xO₃, where x~0.0001–0.001.

Figure 9a,b show the optical images of the Mo and TZM surfaces after cyclic polarization in the acidic chloride solution. The Mo surface displayed a dark, bluish-green color, while the TZM sample showed yellow with purple streaks. The bluish-green color of the pure Mo samples is associated with the reduced state of the MoO₃, possibly Mo₄O₁₁, indicating impurities or defect states [30]. The purple streak indicated the possible presence of tetravalent Mo ions, and the yellow color showed the oxidation of the Mo(IV) to Mo(VI). Figure 9c shows the Raman spectra of the Mo and TZM samples after polarization in the acidic chloride solution. The surface layer consisted of MoO₃-type phases. Several reports pointed out the MoO₃ peaks occurring at 89, 100, 116, 129, 159, 197, 216, 247, 285, 293, 334, 366, 376, 380, 473, 667, 823, and 996 cm⁻¹ [31,32,33]. In this investigation, the pure Mo peaks at 115, 125, 154, 196, 218, 221, 247, 288, 345, and 473 cm⁻¹ were assigned to the MoO₃ phase. B_2g, B_3g, and A_g/B_1g translational rigid MoO₄ chain modes (116, 129, and 159 cm⁻¹) lowered the Raman shifts, possibly due to longer bond lengths because of growth stress in the oxide. The peaks at 571, 730, and 744 cm⁻¹ are associated with Mo₄O₁₁ [34]. MoO₂ and Mo₄O₁₁ phases showed common peaks at 203, 229, 345, 498, 570, and 744 cm⁻¹. Mo₄O₁₁ could be considered an oxygen-deficient MoO_3−x phase, with x = 0.25. Therefore, the oxide phases on the Mo and TZM surfaces after exposure to acid chloride solution were predominantly defective MoO₃ with a minor amount of MoO₂. The major peaks of MoO₂ were 203, 229, 363, 460, 496, 569, and 743 cm⁻¹ [33].

Figure 10a,b illustrate the surface conditions of the Mo and TZM samples after the CP tests in the neutral chloride solution (pH: 5.5). The distinct color shades on the Mo surface could be associated with varying orientations of grains or different crystallographic planes of the grains oxidized at various levels. The TZM surface displayed a uniformly colored surface film. The surface film showed a cracked-clay morphology due to drying of the hydrated film. This observation indicated that the TZM sample formed a thicker surface film than the pure Mo sample. Figure 10c shows the Raman spectra of the surface films of the samples tested in the neutral chloride solution. The Raman peaks were sharper and more intense than those observed in the acidic chloride solution. The phase contents were like those of the acidic chloride solution. High electrochemical impedance values were recorded in this condition, which could be correlated to a thicker surface film.

Figure 11a,b show the surfaces of Mo and TZM after testing in the alkaline chloride solution. The CP results did not show any growth of a passive layer. However, the polished surface was covered by a film due to the accumulation of corrosion products. The pit-like features were not deep and could be removed by polishing with 1200 grit emery paper. Since no passivation was observed, these pits were not due to the breakdown of the passive film but rather caused by localized attack at the microstructural heterogeneities. Figure 11c shows the Raman spectra of the samples after corrosion testing in the alkaline chloride solution. In this case also, a MoO₃-type layer was observed. Additional peaks were associated with the oxychloride compounds of molybdenum. The broadened and weak peaks indicated that the film formed on the TZM was highly defective and thin.

The corrosion potential of the pure Mo sample was influenced by the pH of the environment and described by the relation:

E_corr = −2.65 − 40.15pH mV(Ag/AgCl)

(8)

The slope of −40.15 mV/pH is less than the ideal slope required for using the Mo as a pH electrode; however, a strong linear correlation was observed in the pH range of 1–13. On the other hand, the TZM samples exhibited a pH-dependent corrosion potential in pH 1 and 5.5 solutions, with a slope of −49 mV/pH. However, at pH 13, the corrosion potential did not show a pH dependency. The pH dependency is influenced by the strong bond between Mo and H on the surface [27]. Overall, the TZM samples exhibited more active corrosion potentials, which could be attributed to the higher binding of hydrogen to an increased number of dislocations and TiC carbides, as discussed in the earlier sections. The TZM showed lower corrosion resistance than the pure Mo. The higher exchange current densities for cathodic and anodic reactions on the TZM than those of pure Mo increased the corrosion current despite the marginally steeper Tafel slopes. The exchange current densities indicated the catalytic nature of the TZM samples. Heterogeneous microstructures, such as TiC, ZrC, and Mo₂C, as well as other inclusions, could have served as preferential sites for both anodic and cathodic reactions. The exchange current densities were calculated by considering the equilibrium potential of pH-dependent Mo oxidation as an anodic reaction, and oxygen reduction as a cathodic reaction. The source of hydrogen could come from reactions (9) and (10) involving point defects of the passive layer and water molecule by following the hypothesis of King and Freund [35]:

$${Mo}_{{Mo}_x{O}_y}+O_{M_x{O}_y}+H_{2}O \leftrightarrow \left({HO}^{*}{V}_{Mo}^{″}{O}^{*}H\right)+{Mo}_{aq}^{2+}+2e^{-}$$

(9)

$$\left(H{O}^{*}{V}_{Mo}^{″}{O}^{*}H\right) \to O^{*}{\left(H_2\right)″}_{Mo}{O}^{*}$$

(10)

Incorporating H₂O into the Mo_xO_y oxide layer results in one oxygen in the lattice, implicitly creating one Mo vacancy (V_Mo) and thus making the oxide a p-type semiconductor. In this case, the divalent state of the Mo is assumed for simplicity. It is reported that Mo²⁺ is highly unstable [36]. The reaction could be extended to other valent states of Mo with appropriate charge and mass balances. Here, V″_Mg is double negatively charged, and OH^* is single positively charged, which makes the compound defect (HO^•V″_MgO^•H) a neutral one. This linear defect is considered highly unstable and undergoes decomposition into hydrogen molecules that occupy the cation vacancy sites. The H₂ molecules are weakly adsorbed at the cation vacancy sites and evolve out, leaving behind cation vacancies. During anodic polarization, hydrogen is oxidized, increasing the anodic current density. However, this hypothesis requires a thorough investigation.

The corrosion density of the Mo samples reported in this work matches previously reported values. For example, Badawy and Al-Kharafi [37] reported corrosion current densities of 0.654, 1.53, and 3.39 μA/cm² in pH 2, 7, and 12 solutions. The values obtained in this study were 0.22, 0.4, and 8.1 μA/cm² in the pH 1, 5.5, and 13 solutions. Hu Ping et al. [14] reported the corrosion current densities of TZM as 3.96, 5.46, and 14.9 μA/cm² in acidic, neutral, and alkaline chloride solutions, which are of the same order of magnitude as those determined in this study, as listed in Table 2. The exchange current densities for Mo and TZM alloy have not been widely reported. Gad-Allah and Abd-El Rehman [38] reported an exchange current density of 1.9 × 10⁻¹⁸ A/cm² for Mo in an acid chloride solution at 30 °C. Our study reported a value three orders of magnitude lower than the literature-reported value.

The passivation of Mo is a widely debated subject mainly due to the multivalent state of Mo species. Some authors suggested that the passivation was not due to the formation of a 3D film but by chemisorption of oxygen [24,39]. The widely accepted form of passive film is insoluble MoO₂, which is further oxidized in the transpassive region to soluble MoO₃ with HMoO₄⁻ as the dissolution product [40]. Since MoO₂ has a monoclinic lattice structure with good electrical conductivity, EIS and Mott–Schottky analyses may not help characterize the protective MoO₂ layer. Furthermore, Raman spectra could not unequivocally identify MoO₂ because of the overlapping peaks of the Mo₄O₁₁ phase [32]. The presence of a dual-layered structure with a MoO₂ inner layer and a MoO(OH)₂ outer layer was proposed [37]. The Mott–Schottky results of this study also support the concept of a dual-layered structure. The top layer exhibited n-type semiconductivity, and the bottom layer showed p-type conductivity. MoO₃ and MoO₂ are considered n-type semiconductors [41,42]. However, the mixed MoO_x structure shows p-type semiconductivity and is widely used in thin-film solar cell applications [43]. Furthermore, the hydration of the oxide layer also renders the film p-type by creating cation vacancies as described in reactions (9) and (10). These layers could act as ion-selective barriers and resist corrosion [27,40].

The central focus of this study was to determine whether pitting occurred in Mo alloys in chloride solutions due to the breakdown of the passive film or through other mechanisms. The overall results showed that the transpassive behavior was not localized, unlike what is typically observed in stainless steels. This conclusion is based on the cyclic polarization results that showed no positive hysteresis. During the reverse scan, the current density continued to decrease below the forward scan values, resulting in a cathodic loop. When tested in acidic and neutral chloride solutions, the surface analysis of corrosion-tested samples showed no pitting. However, the localized attack was observed in the alkaline solution, but no stable passive film formed based on the CP results. Therefore, the localized attack can be attributed to the microgalvanic effect because of the microstructural heterogeneity.

4. Conclusions

Based on the electrochemical measurements performed on pure Mo and TZM alloy samples in chloride solutions under three different pH conditions (pH 1, 5.5 and 13), the following conclusions were drawn:

• Cyclic polarization measurements on Mo and TZM alloy samples in acidic and neutral chloride solutions revealed the formation and growth of a protective film. The increase in the current in the transpassive region indicated further oxidation of the surface film but not the breakdown of the passive layer.
• No passivation was observed in the alkaline chloride solutions on the Mo and TZM samples.
• No localized corrosion was observed in the acidic and neutral chloride solutions, but alkaline solutions exhibited a localized attack due to the microstructural heterogeneity of the TZM sample.
• Pure Mo showed better corrosion resistance than TZM in all pH conditions.
• The Mott–Schottky results in the neutral chloride solution suggested the presence of a dual-layered passive film with a p-type inner layer and an n-type outer layer. The charge carrier densities varied from 7.5 × 10¹⁷ to 9.1 × 10¹⁹ cm⁻³, depending on the material and the pH level. The TZM alloy showed a higher defect density than the Mo.
• The pitting corrosion observed in Mo and Mo alloys was not due to the localized breakdown of the passive film but rather to the microgalvanic effect resulting from the microstructural heterogeneity.