Thermal Corrosion Properties of Composite Ceramic Coating Prepared by Multi-Arc Ion Plating

Abstract

In this study, a NiCr/YSZ coating was applied to a γ-TiAl surface using multi-arc ion plating technology to enhance its high-temperature performance and explore the mechanisms of high-temperature oxidation and thermal corrosion. The thermal corrosion properties of the γ-TiAl matrix and NiCr/YSZ coating were investigated at 850 °C and 950 °C using a constant-temperature corrosion test in a 75% Na₂SO₄ + 25% NaCl mixture. The results indicate that after 100 h, the thermal corrosion weight gain of the coating samples was 70.1 mg/cm² at 850 °C and 118.2 mg/cm² at 950 °C. At these temperatures, sulfide formation on the surface increases, leading to a loose and porous surface. After 100 h of high-temperature corrosion at 850 °C, the primary oxidation product on the surface of the coating was tetragonal-ZrO₂. At 950 °C, Y₂O₃, which mainly acts as a stabilizer in YSZ, reacted with Na₂SO₄, resulting in the continuous consumption of Y₂O₃. This reaction caused a substantial amount of tetragonal-ZrO₂ to transform into monoclinic-ZrO₂, altering the volume of the ceramic layer, which induced internal stress, crack propagation, and minor spallation. A continuous and dense internal thermally grown oxide (TGO) layer effectively impeded the diffusion of molten salt substances and oxygen, thereby significantly improving the thermal corrosion resistance of the thermal barrier coating.

1. Introduction

Due to its high specific strength, elastic modulus, and ability to maintain sufficient rigidity and strength at high temperatures, γ-TiAl is frequently used in high-temperature applications, such as marine gas turbines and aircraft engines [1,2]. However, at temperatures exceeding 650 °C, the free energy of TiO₂ and Al₂O₃ becomes comparable, facilitating selective oxidation during high-temperature oxidation. Consequently, a mixed oxide film composed of TiO₂ and Al₂O₃ forms on the surface. The incorporation of a significant amount of TiO₂ inevitably results in a porous structure in the mixed oxide film, which fails to effectively impede oxygen diffusion. As a result, the γ-TiAl alloy exhibits insufficient resistance to oxidation above 650 °C. Moreover, the fuels used in these engines often contain impurities, such as Na and S, which can adhere to the surface of TiAl alloys and cause thermal corrosion due to the alloy’s low resistance to such corrosion [3,4].

Rather than developing entirely new high-temperature corrosion-resistant material, applying a protective coating on the surface of the base alloy is a more practical approach [5]. With the aid of diverse processing techniques, scholars have successfully developed a wide range of protective coatings, thereby significantly enhancing the longevity of the substrate [6,7]. Common high-temperature protective coatings include diffusion coatings, overlay coatings (such as MCrAlY coatings), and thermal barrier coatings [8,9,10,11,12]. Among them, thermal barrier coatings are considered one of the most advanced and promising options due to their ability to extend engine life, improve engine power, and reduce fuel consumption.

Based on their working environment and performance requirements, thermal barrier coatings are mainly categorized into three structural systems: double-layer, multi-layer, and gradient. Double-layer thermal barrier coatings are the most widely used and have the most mature preparation technology. They typically consist of an MCrAlY (where M is Ni, Co, or Ni+Co) intermediate bonding layer and a Y₂O₃-stabilized ZrO₂ (YSZ) ceramic thermal insulation layer [13,14,15,16].

Plasma spraying and electron beam physical vapor deposition (EB-PVD) are the two main techniques used for preparing the thermal barrier coatings [17,18]. Plasma spraying involves the continuous impact, spreading, and solidification of molten droplets, resulting in a flaky coating with numerous pores and low thermal conductivity. This method often incorporates large amounts of slag and foreign pollutants in the coating. At high temperatures, corrosion or oxidation can occur at these defects, which can initiate cracks and ultimately cause coating failure [19,20]. In contrast, EB-PVD offers strong chemical bonding between the coating and the substrate, resulting in high bonding strength. However, it suffers from insufficient bulk density of the material due to rapid evaporation and shadowing effects [18,21,22].

Multi-arc ion plating technology is recognized for its high deposition efficiency, uniform and dense coatings, excellent conformal coverage, controllable coating preparation temperature, sufficient atomic diffusion, strong metallurgical bonding between the coating and substrate, and high bonding strength of the coating [23,24]. Recent advancements in multi-arc ion plating technology have enabled the development of various high-temperature protective coatings on the surface of γ-TiAl alloys, including aluminum coatings, chromium coatings, alloy coatings such as MCrAlX (M = Ni, Co; X = Y, Hf), oxide coatings such as Cr₂O₃ coatings, and composite coatings such as NiCrAlY/Al [25,26,27,28,29].

In this study, a NiCr/YSZ coating was applied to a γ-TiAl surface using multi-arc ion plating. The thermal corrosion resistance of the coating was evaluated by analyzing the thermal corrosion kinetics curves, surface and cross-sectional morphology, and corrosion products at different temperatures.

2. Experiment

A NiCr/YSZ coating was applied to cast γ-TiAl alloy (composition shown in

Table 1. Chemical composition of γ-TiAl alloy (wt. %).

Ti	Al	V	Nb	Cr	N	C	O
50.64	46.50	1.50	1.00	0.20	0.05	0.10	0.01

; 14 × 14 × 5 mm³) using multi-arc ion plating. After grinding and polishing, the surface roughness of the substrate was Ra ≤ 0.8. The multi-arc ion plating equipment used in this experiment was produced by Beijing Taikono Company. A NiCr alloy target (mass fraction: 3:2, purity: 99.99%) was used as the source material to prepare the NiCr alloy bonding layer, whereas a ZrY alloy target (92:8 mass fraction) was used to prepare the YSZ coatings. The oxygen-to-argon flow ratio was maintained at 1:10, and the chamber pressure was set to 0.7 Pa. The process parameters were as follows: a bias voltage of −200 V, an arc current of 80 A, a working pressure of 0.7 Pa, and a chamber temperature of 400 °C.

Thermal corrosion tests were conducted using a salt coating method. The substrate and coated samples were preheated with an alcohol lamp. A brush dipped in saturated Na₂SO₄ aqueous solution was then used to apply a salt coating of 3.5 mg/cm². The thermal corrosion tests were performed in an SX-49 box furnace (Tianjin Test Instrument Co., Ltd., Tianjin, China) at 850 °C and 950 °C for 100 h. The samples were removed every 10 h, immersed in boiling water for 30 min, and then dried and weighed after removing the salt film. Subsequently, the salt coating was reapplied using the original method to complete one thermal corrosion cycle.

During the tests, the weight gain of the samples was measured using an electronic analytical balance (Shanghai Heaven Precision Instrument Co., Ltd., FA-1004; accuracy 0.1 mg, Shanghai, China) [30,31]. The microstructure and composition of the samples before and after thermal corrosion were examined using scanning electron microscopy (SEM; FEI Quanta200, Shanghai, China) combined with energy-dispersive X-ray spectroscopy (EDS; OXFORD Xplore30, Shanghai, China). The phase structure of the substrate and coating was analyzed using X-ray diffraction (XRD; Bruker D8-ADVANCE, Beijing, China), with the following test parameters: a tube voltage of 50 kV, a tube current of 40 mA, a scanning range of 10–100°, a Cu Kα radiation source, a step size of 0.02°, and an acquisition time of 8 s per step.

3. Results

3.1. Microstructure and Surface Morphology of Coatings

Figure 1a shows an SEM image of the NiCr/YSZ coating obtained by multi-arc ion plating at 5000× magnification. The image reveals droplets with a diameter of approximately 2–4 μm that have solidified and attached to the surface of the coating. This results in surface unevenness, which is likely attributed to the high bias voltage, ion bombardment, and droplet impact. Although this unevenness is a typical feature and a major drawback of multi-arc ion plating, the overall appearance of the coating remained smooth and compact, with no visible defects, such as holes or cracks.

The EDS analysis results in Figure 1b indicate that the coating surface consists of a YSZ ceramic layer. The atomic mass ratio of Y to Zr on the coating surface was approximately 5:66, which is higher than the target mass ratio of 8:92. This discrepancy is ascribed to the segregation of components due to varying atomic sputtering rates. Specifically, Zr ions on the surface of the target initially underwent oxidation, which affected the sputtering yield of Zr ions and consequently reduced the amount of Zr in the coating.

Figure 2 shows the cross-sectional microstructure of the NiCr/YSZ coating, which is primarily composed of two layers: an 8 μm YSZ ceramic layer on the surface and a 3 μm NiCr adhesive layer between the YSZ ceramic layer and the γ-TiAl substrate. The NiCr coating was flat, compact, and free of significant defects. The inner side of the YSZ thermal barrier coating was tightly bonded to the NiCr layer without any pores or holes. However, the outer surface of the YSZ layer was uneven due to ion and metal droplet bombardment during ion sputtering. Overall, the intermediate NiCr layer between the coating and substrate was flat, compact, and free of gaps, holes, and impurities.

The XRD pattern (Figure 3) shows that ZrO₂ in the YSZ thermal barrier coating was predominantly composed of a tetragonal phase (t-phase), with some cubic phase (c-phase). Pure ZrO₂ was stable at temperatures around 1443–2643 K. However, the addition of Y₂O₃ stabilized the t-phase and restricted phase changes, with minimal impact on the thermal conductivity of ZrO₂. The dense YSZ contained numerous oxygen vacancies, replacement atoms, and other point defects, which stabilized the t-phase even at room temperature. Consequently, the t-ZrO₂ in the YSZ thermal barrier coating remained stable and resistant to failure, indicating enhanced thermal insulation performance and high-temperature oxidation resistance of the coating.

3.2. Thermal Corrosion at 850 °C

3.2.1. Thermal Corrosion Kinetics Curve

Figure 4 shows the thermal corrosion kinetics curve for the γ-TiAl substrate and the NiCr/YSZ coating when exposed to a 75% Na₂SO₄ + 25% NaCl mixture at 850 °C. The curve exhibits two distinct stages: an incubation stage and an acceleration stage. During the initial thermal corrosion stage, both samples experienced relatively low weight gain, resulting in a smoother kinetic curve. This stage, lasting for approximately 30 h at 850 °C, represents the incubation period of thermal corrosion. In the acceleration stage, the weight gain of the samples continuously increased over time. Consequently, the curve exhibits a linear progression. The weight gain also increased with temperature.

The NiCr/YSZ-coated samples exhibited slower weight gain compared to the γ-TiAl substrate, and the rate of weight gain for the NiCr/YSZ-coated sample remained extremely low. After 100 h of thermal corrosion, the weight gain of the γ-TiAl matrix was 141.4 mg/cm², which is significantly higher than that of the NiCr/YSZ coating. The weight gain of the NiCr/YSZ-coated sample (Figure 4) was 70.1 mg/cm², which is only 49.5% of that of the γ-TiAl matrix. This indicates that the NiCr/YSZ coating substantially enhancesd the thermal corrosion performance of γ-TiAl when exposed to the Na₂SO₄/NaCl mixture at 850 °C.

3.2.2. Surface Morphology after Thermal Corrosion

Figure 5 shows the surface micromorphology of the NiCr/YSZ-coated samples after thermal corrosion in Na₂SO₄/NaCl mixture at 850 °C for 0, 10, 40, and 100 h. As depicted in Figure 5b, after 10 h of thermal corrosion, the surface of the NiCr/YSZ coating became loose and porous, and some areas exhibited exfoliation and corrosion pits, with small, dispersed lumpy grains. At this stage, the coating was still in the incubation phase of thermal corrosion and exhibited high corrosion resistance. Figure 5c illustrates that after 40 h of thermal corrosion, the coating surface became looser, with more pronounced peeling. Some areas show reticular corrosion holes, where small particles were mixed. Figure 5d shows that after 100 h of thermal corrosion, the coating surface was completely cracked.

3.2.3. Cross-Sectional Microstructure after Thermal Corrosion

Figure 6a shows a cross-sectional micrograph of the NiCr/YSZ coating after 100 h of thermal corrosion at 850 °C. The image reveals that the coating tends to thicken due to the increased volume of oxide sulfides generated on the surface (Figure 6b). The YSZ coating was relatively loose compared to the non-corroded thermal barrier coating, with increased porosity along with cracks and voids that facilitate oxygen diffusion into the interior via these loose cracks. Consequently, the thermally grown oxide (TGO) layer inside the coating also became thicker, measuring approximately 3 μm.

3.2.4. Thermal Corrosion Products

Figure 7 presents the XRD pattern of the NiCr/YSZ-coated samples after thermal corrosion in a Na₂SO₄/NaCl mixture at 850 °C for 100 h. After thermal corrosion, the predominant phase on the surface of the samples remained tetragonal, with some cubic phase. No monoclinic phase ZrO₂ diffraction peaks were observed. This suggests that Y primarily exists as Y₂O₃ after thermal corrosion, providing good thermal corrosion resistance. In thermal barrier coatings, the transformation of tetragonal (t) to monoclinic (m) phase ZrO₂ can increase stress within the ceramic layer and expand its volume. Therefore, the content of m-phase after thermal corrosion can indirectly reflect the thermal corrosion resistance of the coatings. The absence of a t → m transition in the NiCr/YSZ coating indicates that Y₂O₃ in the ceramic layer stabilizes ZrO₂. However, Y₂O₃ is not detected in the XRD pattern due to the low Y content.

3.3. Thermal Corrosion at 950 °C

3.3.1. Thermal Corrosion Kinetics Curve

Figure 8 shows the thermal corrosion kinetics curve of the γ-TiAl substrate and the NiCr/YSZ coating in the Na₂SO₄/NaCl solution at 950 °C. Similar to the previous case, the thermal corrosion of both samples followed a two-stage process, with an incubation stage and an acceleration stage. The incubation stage lasted 20 h, during which both samples experienced minor weight loss, resulting in a relatively smooth kinetics curve. Compared with thermal corrosion at 850 °C, the incubation stage of the two samples at 950 °C was shorter due to the higher temperature, which accelerates oxidation and vulcanization on the surface. During the acceleration stage, the weight gain rate of the γ-TiAl matrix sample continued to increase, with the curve showing a linear progression that accelerated after 80 h. This acceleration is likely due to the exfoliation of the surface coating and rapid oxidation of the newly exposed internal surface. In the initial 40 h, the weight gain of the coated sample (Figure 8) was obviously slower than that of the substrate, with the weight gain rate remaining very low. However, after 40 h, the weight gain rate of the NiCr/YSZ coating significantly increased, indicating a decrease in its thermal corrosion resistance over time. After 100 h, the thermal corrosion weight gain of the γ-TiAl matrix was 170.3 mg/cm², whereas the weight gain of the NiCr/YSZ-coated sample was 118.2 mg/cm², which is approximately 69.4% of that of the γ-TiAl matrix.

3.3.2. Surface Morphology after Thermal Corrosion

Figure 9 shows the surface morphology of the NiCr/YSZ coating after corrosion at 950 °C for 10, 40, and 100 h in the Na₂SO₄/NaCl solution. Figure 9a reveals that after 10 h of high-temperature corrosion, cracks formed around the large particles on the surface of the NiCr/YSZ coating. These metal particles, primarily composed of t-phase ZrO₂ and Y₂O₃, are highly uneven. This results in the transformation of some ZrO₂ from t-phase to m-phase in the absence of Y₂O₃ stabilization, causing volume expansion of m-phase ZrO₂. This expansion caused surface cracking under stress.

Figure 9b demonstrates that after 40 h of high-temperature corrosion, corrosion products began to accumulate on the surface of the NiCr/YSZ coating. The coating became loose and porous, with flaking in some areas, forming corrosion pits. Additionally, small grains were distributed over the surface. This period indicates the end of the incubation stage, during which the coating still exhibited high thermal corrosion resistance. XRD analysis shows significant corrosion of Y₂O₃ into Y₂(SO₄)₃ in the ceramic layer, resulting in t → m phase transformation. This transformation caused crack expansion and an increase in surface volume due to increased internal stress, resulting in a fluffy and porous coating layer.

Figure 9c shows that after 100 h of high-temperature corrosion, the number of corrosion products on the NiCr/YSZ coating gradually increased. Strip-shaped crystalline substances, several microns long and several hundred nanometers wide, appeared on the surface of the coatings. Such morphologies have been observed in some previous studies, but their formation mechanisms are not well understood [32,33,34,35]. Compared with the thermal corrosion morphology at 850 °C, the corrosion products at 950 °C were larger and more pronounced, with more noticeable pores between the particles.

3.3.3. Cross-Sectional Microstructure after Thermal Corrosion

Figure 10 displays the cross-sectional micromorphologies of the NiCr/YSZ coating after thermal corrosion in Na₂SO₄/NaCl at 950 °C for 100 h. It is clear that a 6 μm thick TGO layer appeared on the NiCr/YSZ coating, indicating that the coating undergoes high-temperature oxidation during thermal corrosion. The surface of the YSZ layer appeared loose and porous due to the reaction between Y₂O₃ and SO₃⁻, forming Y₂(SO₄)₃. This reaction leads to the transformation of t-ZrO₂ into m-ZrO₂ in the absence of Y₂O₃ stabilization. During this transformation, ZrO₂ crystal grains expand, creating a loose and porous surface that facilitates the diffusion of molten salt from the surface to the interior.

At high temperatures, the dissolved oxygen in the molten salt diffuses into the interface between the ceramic layer and the bonding layer, generating TGO through high-temperature oxidation of the thermal barrier coating. After 100 h of thermal corrosion at 950 °C, crisscrossed cracks appeared in the ceramic layer of the NiCr/YSZ coating, which signifies that the thermal barrier coating has started to fail. At this stage, the ceramic layer undergoes a t → m phase transformation, causing crack expansion due to excessive internal stress.

3.3.4. Thermal Corrosion Products

Figure 11 presents the XRD pattern of the NiCr/YSZ-coated samples after various thermal corrosion durations. The pattern indicates that most t-phase ZrO₂ is converted into m-phase ZrO₂, and Y₂(SO₄)₃ is formed as a corrosion product. The ratio of t-phase to m-phase varies with the duration of thermal corrosion. Generally, the proportion of m-phase is calculated by analyzing the phase transformation of ZrO₂ adjacent to the diffraction angle of 2θ = 30° using the following formula [36,37]:

$$\eta ={\displaystyle \frac{I_m\left(\overline{1}11\right)+I_m\left(111\right)}{I_m\left(\overline{1}11\right)+I_m\left(111\right)+I_t\left(111\right)}}\times 100\mathrm{\%}$$

(1)

According to Figure 11, the main phases on the surface evolve with corrosion time. The diffraction peak of m-phase ZrO₂ appeared and gradually intensified with longer thermal corrosion times and higher temperatures. A weaker diffraction peak of Y₂(SO₄)₃ began to appear at 40 h of thermal corrosion and intensified further at 100 h, indicating the formation of trace amounts of this corrosion product.

Figure 12 compares the calculated volume fraction of the m-phase ZrO₂ in the NiCr/YSZ coating after thermal corrosion at different temperatures. The volume fraction of the m-phase was 2.25% after 10 h of thermal corrosion at 950 °C. By extending the corrosion time to 100 h, the volume fraction of m-phase increased to 38.5%, indicating significant t → m transformation over the extended period.

4. Discussion

The thermal corrosion process in the NiCr/YSZ coating is complex due to its intricate structure and composition. In a thermal corrosion environment, the reaction between the corrosive medium and the coating is accompanied by high-temperature oxidation. Therefore, failure and exfoliation of the NiCr/YSZ coating after thermal corrosion may be attributed to various factors, including internal stresses generated by the TGO layer during high-temperature oxidation, thermal stress resulting from significant temperature fluctuations due to mismatched thermal expansion coefficients between the coating and the substrate, and additional stress generated by phase changes in the ceramic layer during thermal corrosion.

The reaction between Y₂O₃, which stabilizes t-ZrO₂, and the molten salt contributes to the transformation from t-phase to m-phase in the ceramic layer. The stress induced by this phase transformation is the main cause of failure of the NiCr/YSZ coating [36,38,39,40]. In this study, the thermal corrosion mechanism of thermal barrier coatings was primarily analyzed from the perspective of phase transformations. The thermal corrosion process can be divided into the following three stages:

• Initial stage: As the temperature rises, molten salt forms a film on the surface of the NiCr/YSZ coating. The molten salt infiltrates the coating through cracks and pores, increasing the contact area between the molten salt and the coating.
• Intermediate stage: The molten salt reacts with Y₂O₃, generating corrosion products and gradually consuming Y₂O₃. This leads to the transformation of some t-ZrO₂ into m-ZrO₂. During the t → m transformation, the volume expands, leading to the formation and propagation of cracks. This allows the thermal corrosion reaction to penetrate further into the coating, thereby accelerating the thermal corrosion process.
• Final stage: The internal stress from the t → m transformation, combined with the stresses from TGO and mismatched thermal expansion, can result in the formation of cracks between the metal matrix, bonding layer, and ceramic layer. These stresses eventually lead to coating failure and exfoliation.

In the high-temperature corrosive environment, molten Na₂SO₄ can be considered to contain an alkaline component (Na₂O) and an acidic component (SO₃), i.e.,

$${\mathrm{N}\mathrm{a}}_2{\mathrm{S}\mathrm{O}}_4\leftrightarrow {\mathrm{N}\mathrm{a}}_2\mathrm{O}+{\mathrm{S}\mathrm{O}}_3$$

(2)

As the above reaction proceeds, the concentrations of Na₂O and SO₃ increase until equilibrium is reached. The equilibrium constant for the reaction can be expressed as follows:

$$K={\displaystyle \frac{C_{\left({\mathrm{N}\mathrm{a}}_2\mathrm{O}\right)}\times C_{\left({\mathrm{S}\mathrm{O}}_3\right)}}{C_{\left({\mathrm{N}\mathrm{a}}_2{\mathrm{S}\mathrm{O}}_4\right)}}}$$

(3)

where K represents the equilibrium constant and C denotes the material concentration. Since K is related to the ambient temperature, it has a certain value. Due to the cracks and pores formed during the multi-arc ion plating process, as well as the inherent porosity of the ceramic layer, a significant amount of Na₂SO₄ can penetrate into the coating. When the concentration of alkaline Na₂O reaches a certain level, it reacts with zirconia in the ceramic layer to form spinel, i.e.,

$${\mathrm{N}\mathrm{a}}_2\mathrm{O}+{\mathrm{Z}\mathrm{r}\mathrm{O}}_2\to {\mathrm{N}\mathrm{a}}_2{\mathrm{Z}\mathrm{r}\mathrm{O}}_3$$

(4)

The rate of Reaction (4) is affected by the concentrations of both Na₂O and Na₂ZrO₃. A high Na₂O concentration and a fairly low Na₂ZrO₃ concentration is conducive to the forward reaction in Reaction (4). The concentration of Na₂O is derived from Equation (3), whereas the concentration of molten salt, C(Na₂SO₄), basically remains constant, as do the concentrations of Na₂O and SO₃. The maximum Na₂O concentration, determined by a small equilibrium constant K, is low, causing Reaction (4) to proceed at an extremely slow rate.

As the reaction progresses, the concentration of the corrosion product (Na₂ZrO₃) continuously increases, which also influences the corrosion rate. Because the amount of Na₂ZrO₃ is very small, it is not detected in the XRD pattern. During the initial stage of thermal corrosion, Na₂ZrO₃ is dissolved in the molten salt at high temperatures and later crystallizes and precipitates in the thermal barrier coating upon reaching saturation. Although Reaction (4) causes only minor damage to the thermal barrier coating, the crystallization of the corrosion product is extremely detrimental due to the negligible amount of ZrO₂ participating in the reaction. However, the precipitation and growth of Na₂ZrO₃ in cracks and between coating layers exacerbate crack expansion and coating layer exfoliation.

When exposed to SO₃ at high temperatures, Y₂O₃, which stabilizes t-ZrO₂, also reacts to form Y₂(SO₄)₃, as follows:

$$3{\mathrm{S}\mathrm{O}}_3+{\mathrm{Y}}_2{\mathrm{O}}_3={\mathrm{Y}}_2{\left({\mathrm{S}\mathrm{O}}_4\right)}_3$$

(5)

As Reaction (5) proceeds, Y₂O₃ is continuously consumed, leading to a sustained transformation of ZrO₂ from the t-phase to the m-phase. This transformation is accompanied by a 3%–5% volume expansion, which increases internal stress between the ceramic layers and causes the formation and propagation of cracks. When the volume change reaches a certain level, the coating begins to peel off. In addition, similar to the formation of Na₂ZrO₃, the growth of Y₂(SO₄)₃ crystals at the cracks and between the coating layers imposes further stress on the ceramic layer.

During the preparation of thermal barrier coatings by multi-arc ion plating, metal particles are the first to fall off in a high-temperature environment. This results in the formation of cracks and pores around the exfoliated particles, which create channels for the diffusion of molten salt into the coatings. This increases the contact area between the molten salt and the ceramic layer, thereby triggering corrosion. The reaction produces Na₂ZrO₃ and Y₂(SO₄)₃ crystals. These corrosion products grow within the cracks and crystallize between the coating layers, generating thermal stress due to the formation and expansion of cracks in the ceramic layer. The generation of Y₂(SO₄)₃ crystals continuously consumes Y₂O₃, causing ZrO₂ to undergo a t → m phase transformation. This phase transformation and the associated volume change increase the internal stress within the ceramic layer. Additionally, during thermal corrosion, a TGO layer is formed between the bonding layer and the ceramic layer due to high-temperature oxidation. At high temperatures, the mismatch in thermal expansion coefficients between NiCr and YSZ increases internal stress. Under the combined effects of various stresses, the NiCr/YSZ coating eventually deteriorates and flakes off.

5. Conclusions

A NiCr/YSZ coating was applied on a γ-TiAl surface using multi-arc ion plating technology to enhance its high-temperature performance, and the mechanisms of high-temperature oxidation and thermal corrosion were comprehensively examined. The oxidation products on the surface of the NiCr/YSZ-coated samples after 100 h of high-temperature corrosion at 850 °C in a 75% Na₂SO₄ + 25% NaCl solution were mainly t-ZrO₂, without any phase transformation. However, at 950 °C, although the main component remained t-ZrO₂, a significant portion transformed into m-ZrO₂. This was because Y₂O₃, the stabilizer present in YSZ, reacted with Na₂SO₄ during thermal corrosion, continuously consuming Y₂O₃, which promoted the transformation from t-ZrO₂ to m-ZrO₂. The resultant internal stress from the volume change in the ceramic layer finally led to the formation and expansion of cracks. The NiCr/YSZ coating exhibited good thermal corrosion resistance after 100 h of corrosion at both temperatures, but the surface film layers presented varying degrees of damage. This damage included coating exfoliation at the edges and different degrees of crack expansion on the surface. The metal particles on the coating surface gradually flaked off in the high-temperature environment, affecting the thermal corrosion performance of the coatings. The thermal corrosion of barrier coatings was accompanied by high-temperature oxidation, leading to the formation of a TGO layer at the interface between the bonding layer and the ceramic layer. A continuous and dense TGO layer effectively decelerated the diffusion of molten salt and oxygen, significantly enhancing the thermal corrosion resistance of the thermal barrier coatings.