Microstructure and High-Temperature Oxidation Behavior of Cold-Sprayed CoNiCrAlY Coatings Deposited by Different Propellent Gases

Abstract

CoNiCrAlY coatings are widely used as typical high-temperature materials to enhance the surface performances of nickel-based superalloy materials, which can be used as bond coats in thermal barrier coatings and abradable seal coatings. In this study, high-pressure cold spray technology was used to deposit CoNiCrAlY coatings on nickel-based superalloy substrates. The microstructure characteristics and oxidation behaviors of CoNiCrAlY coatings prepared by different gas types and cold spray parameters were systematically investigated. EBSD analysis showed that the deformation of the helium coating was more distinct, and the grain size of the coating fabricated by helium was smaller than that by nitrogen as seen from the grain morphology. The high-temperature oxidation results showed that the coating oxide film thickness varied parabolically with time for both coatings after 500 h isothermal oxidation at 800 °C, 900 °C, and 1000 °C, and the oxidation rate of the coating after heat treatment was lower than that of the as-sprayed coating. In addition, the shrinkage of the aluminum reservoir inside the coating, the element diffusion rate, and the amount and type of oxide generation on the surface all affected the oxidation process. Additionally, the helium coating had a lower oxidation growth rate and better oxidation resistance. Therefore, cold spray can be an alternative way to fabricate high-quality CoNiCrAlY coatings.

1. Introduction

Nickel-based superalloys have high strength, creep resistance, and high-temperature stability, making them widely used as aero-engine compressors and turbines [1,2,3,4,5]. However, these alloys alone can no longer meet the extremely harsh working environment requirements. The demand for higher gas turbine engine performance has led to the development of thermal barrier coating (TBC) systems that are applied to engine thermal components. Typically, a TBC consists of an underlying MCrAlY bond coat with an yttria partially stabilized zirconia ceramic top coat [6,7,8]. MCrAlY alloys (M = Ni, Co, or both) are used in the bonding layer of TBCs to improve the oxidation and hot corrosion resistance of the nickel-based superalloys, prolonging the service life of the components [9].

At present, the preparation of MCrAlY coatings is mainly accomplished through conventional thermal spraying techniques, such as vacuum plasma spraying/low-pressure plasma spraying (VPS/LPPS), air-plasma spraying (APS), and high-velocity oxygen fuel spraying (HVOF). However, these methods introduce a high-temperature heat source, which inevitably causes phase transition and oxidation, seriously affecting their overall performance [10,11,12,13,14]. Compared with conventional thermal spraying, supersonic cold spraying (also known simply as “cold spraying”) has certain advantages, such as low heat input, low coating porosity, and low oxidation levels within the coating.

Cold spraying technology is a new type of coating processing technology developed in recent years. Its working principle is that powder particles pass through a convergent–divergent nozzle in a solid state and are accelerated to supersonic speed (500–1200 m/s). Upon impact onto the substrate surface, the particles have high kinetic energy, and adiabatic shear instability (ASI) occurs at the interface between the particles and the substrate through severe plastic deformation, thereby adhering to the substrate to form a dense coating. Therefore, oxidation and phase change within the coating can be effectively avoided [15,16,17,18,19]. This makes cold spraying promising for repair, surface enhancement, and additive manufacturing in a variety of industries [20].

Several scholars have investigated the preparation of MCrAlY coatings through cold spraying. Zhang et al. [21] prepared CoNiCrAlY coatings on nickel-based high-temperature alloy substrates using low-pressure cold spraying. Their study demonstrated that CoNiCrAlY coatings with low porosity and low oxygen content could be successfully fabricated using this technique. Karaoglanli et al. [13] successfully prepared CoNiCrAlY coatings with dense structures, low oxygen content, and low porosity (around 1.0%) using high-pressure cold spraying, and showed excellent oxidation resistance through detailed analysis of the oxidation behavior. Meanwhile, Richer et al. [12] compared the microstructures and oxidation behaviors of MCrAlY coatings prepared by different deposition techniques, finding that coatings produced by cold spraying exhibited denser structures and better oxidation resistance properties. Guo et al. [7,22] prepared in situ nanocrystalline MCrAlY coatings via cold spraying, demonstrating that this method promotes powder particle deformation and grain refinement, resulting in coatings with good interfacial bonding, a dense microstructure, and excellent oxidation resistance. Previous studies have confirmed that cold spraying is a viable method for preparing CoNiCrAlY coatings. Propellant gases (such as nitrogen and helium) play an important role in the preparation of cold-sprayed coatings, each offering distinct advantages. However, most current studies focus on comparing the oxidation properties of coatings prepared by different techniques, while the effects of propellant gases on coating structure and properties remain less explored. In particular, the deformation behavior of particles during deposition with different gases and their impact on coating oxidation behavior has not yet been reported.

Therefore, this study focuses on the microstructure of CoNiCrAlY coatings prepared under different process parameters (such as propellant gas) and investigates the variations in particle deformation and grain morphology during the comparative deposition process. Additionally, the study explores how these different grain morphologies affect the oxidation properties of the coatings. The results also provide valuable insights and guidance for the industrial application of cold-sprayed CoNiCrAlY coatings.

2. Experimental Details

The feedstock powder used for deposition coatings is a commercially available gas-atomized CoNiCrAlY powder (Amdry 9951, Oerlikon Metco, Mequon, WI USA). The morphology and particle size distribution information of the powder is shown in Figure 1. It can be seen that the CoNiCrAlY powder particles are almost perfectly spherical with a smooth particle surface without any satellites, indicating that the powder has good fluidity during the cold spraying process. The particle size distribution of the powder is relatively uniform with an average particle size of about 20 μm. Most of the powder size falls in the range of 10–40 μm. Figure 1c shows the cross-sectional morphology of CoNiCrAlY powder in BSE-SEM mode. It can be seen that the interior of the powder particles is dendritic in structure, and the dark gray β-NiAl phase is dispersed arbitrarily in the light gray γ-matrix phase of the whole powder. The substrate used for CoNiCrAlY coating deposition was a 5 mm thick round plate of IN625 nickel-based superalloy. Prior to cold spraying, the substrate was polished to remove its surface oxide film with low surface roughness and then cleaned with alcohol. The roughness (Ra) of the polished surface of the substrate was about 0.04 ± 0.01 μm.

In this study, CoNiCrAlY coatings with different process conditions were deposited onto the substrate using a high-pressure cold spray system (PCS1000, Saitama, Japan). Nitrogen and helium were used as carrier gases during the cold spraying process. The gas pressures were set to 5 MPa and 3 MPa, with the accelerating gas heated to 950 °C and 700 °C, respectively. Coating deposition was carried out at a speed of 500 mm/s, with the distance between the nozzle and the substrate surface maintained at 30 mm. The specific deposition parameters are provided in

Table 1. Cold spraying deposition parameters of CoNiCrAlY coatings.

Parameters	N2 Spray	He Spray
	Values	Values
Gas pressure	5 MPa	3 MPa
Gas temperature	950 °C	700 °C
Stand-off distance	30 mm	30 mm
Spray angle	90°	90°

. After deposition, the as-sprayed samples were heated to 1080 °C at a rate of 6 °C/min in a vacuum furnace and maintained at this temperature for 4 h under a pressure below 10⁻² Pa.

Isothermal oxidation experiments in static air conditions at 800, 900, and 1000 °C were carried out in a muffle furnace. To study the evolution process of the coatings in the oxidation experiments, the samples were taken out and analyzed after thermal exposure in the furnace for 20, 50, 100, 200, and 500 h, respectively. The oxidation experiment requires that the furnace temperature be stabilized at the experimental temperature, and then the samples placed in the furnace for oxidation. After reaching the predetermined oxidation time, the samples were taken out of the furnace and air-cooled to room temperature. The surface morphology, microstructure, and composition of the samples after different stages of oxidation were characterized.

The particle size distribution of the CoNiCrAlY powders was measured by a laser particle size analyzer (Mastersizer 3000, Malvern, UK). The morphology of the powders and their cross-sectional microstructure were characterized by scanning electron microscopy (SEM, GeminiSEM 300, Zeiss, Jena, Germany) and electron backscatter diffraction (EBSD).

In order to observe the microscopic morphology of the coating cross-section and understand the bonding of the coating interface, the prepared coatings were cut into 10 mm × 10 mm specimens by a wire-cutting machine and then the pieces were inlaid with resin. After grinding and polishing the sample according to the standard sample preparation process, optical microscopy (OM, DMI500M, Leica, Wetzlar, Germany), scanning electron microscopy (SEM, GeminiSEM300, Zeiss, Jena, Germany), and energy dispersive x-ray energy spectrometry (EDS, INCA-x-act, Oxford Instrument, Oxford, UK) were used to surface morphology, microstructure, and oxidation of the deposited and oxidized coatings. The TGO-scale thickness of each oxidized sample was recorded to evaluate the growth kinetics of TGO. At least six SEM images of TGO were taken for each sample, and multiple TGO thickness values were obtained on each cross-sectional image using the image analysis software Image J 1. The average TGO scale thickness of the samples was used to derive the oxidation rate constants of the coating.

The polished coatings were also knocked out of the resin, and the coating cross-sections were analyzed by electron backscatter diffraction (EBSD). EBSD images were taken at a scanning step of 0.1 μm before the recorded EBSD data were analyzed using Oxford Instruments Channel 5 software. The phase composition of the deposited coatings before and after oxidation was studied by X-ray diffraction (XRD, SmartLab III, Rigaku, Tokyo, Japan) with a scan step of 0.02°/s and a range of 20–90°.

3. Results and Discussion

3.1. Microstructures

Figure 2 shows the cross-sectional microstructure of the annealed CoNiCrAlY coating, indicating a well-bonded interface between the coating and the substrate, and a dense structure without obvious defects. Some micropores are visible at the particle–particle interfaces, which are caused by uneven deformation of the powder particles under high-speed impact during the cold spray process.

Figure 2a,b depict the CoNiCrAlY coatings prepared using nitrogen as the propellant gas, which showed a porosity of approximately 1.0%. In contrast, Figure 2c,d show the CoNiCrAlY coating prepared with helium, which is almost free of pores within the coating, resulting in higher densities than the nitrogen coating. This is mainly attributed to the higher impact velocity of the particles under helium, leading to significant heating of the particles, and triggering adiabatic shear instability (ASI). The jets and splats formed during the deformation tend to fill the voids in the coating, resulting in a dense microstructure [23].

The XRD spectra of CoNiCrAlY powder and cold-sprayed CoNiCrAlY coatings after annealing are shown in Figure 3. It can be seen that both the coating and the powder are composed of γ-matrix (Co-Ni-Cr) phase and β-NiAl phase, which proves that the cold spray deposition does not change the phase composition and no new phase is formed. It is also clear that the β-NiAl phase is uniformly distributed in the γ-matrix phase by analyzing Figure 4b,e.

The cross-sectional microstructures of the above CoNiCrAlY coating were characterized by EBSD with an area of 120 μm × 80 μm. Figure 5 shows the band contrast images, IPF, and local average misorientation (LAM) maps of CoNiCrAlY coatings in different deposition states. From the band contrast (BC) maps (Figure 5a,d), it can be seen that the CoNiCrAlY powder particles undergo strong plastic deformation during the cold spray deposition process, and the particles also have a distinct flat shape. In particular, strong distortions of the crystal structure and grain refinement occur at the particle-to-particle or particle-to-substrate boundaries (Figure 5b,e); as can be seen, these areas cannot be captured by EBSD and result in black areas in the IPF maps. It can also be seen that the degree of particle deformation is significantly higher in the helium-sprayed coating than in the nitrogen-sprayed coating, which is related to the higher particle impact velocity accelerated by the helium.

Figure 4a,d and Figure 5b,e show the EBSD inverse pole figure (IPF) maps for the as-deposited and heat-treated coatings, and the grain sizes (mean with standard deviation) within the diagram were calculated. Figure 4d demonstrates that the grain size of the powder particles is 1.36 μm. However, the internal grains of the deposited state coatings undergo significant grain refinement, with grain sizes of 0.36 μm and 0.32 μm for the two deposited states, respectively. The reason for this is dynamic recrystallization due to severe plastic deformation during deposition. It can also be seen from the IPF maps of the deposited state that the grain size within the coating becomes very heterogeneous, and the grain size in the central region of the powder particles is much larger compared to the boundary location, which is related to the greater degree of plastic deformation occurring at the interface location. After heat treatment, the grain growth can be clearly observed from the IPF maps (Figure 4a,d), and the average grain size inside the coating increases to 0.83 μm and 0.77 μm, respectively. This also shows that there is basically no unidentified area inside the coating after heat treatment, and similar to the deposited state, the grain size inside the coating is uneven, and the grains at the original interface are smaller than those in the middle of the particles. This phenomenon is mainly related to the static recrystallization and stress relief that occurs during the heat treatment [24] (the change in stress–strain during the heat treatment phase is explained in detail below).

Figure 4e,f and Figure 5e,f show the local average misorientation (LAM) for the CoNiCrAlY coatings in the as-deposited and heat treatment states. LAM maps are usually used to highlight the local variations in dislocation density and its distribution. Figure 4e,f show that the orientation difference is more pronounced in the as-deposited coating compared to the substrate. The substrate is predominantly dark blue and the coating is green to red (with a gradual increase in orientation difference from blue to red), indicating that the coating undergoes strong plastic deformation during the deposition process, thus causing high stress–strain and high dislocation density due to work hardening. By comparing the different deposition states (Figure 5e,f), it is clear that the coating prepared with helium has more red pixel dots, i.e., the coating is more severely plastic-deformed and has a higher internal stress–strain, which is due to faster particle impact velocity under helium conditions. Compared with the as-deposited state, the heat-treated coatings show a significant recovery (as shown in Figure 4e,f), proving a significant reduction in the dislocation density inside the coating. The reason is that the grain growth and recrystallization occur during the heat treatment. In addition, residual stress within the coating is significantly reduced during heat treatment, resulting in work hardening being released and the internal dislocation density being reduced [25].

3.2. Oxidation Behaviors

Figure 6 shows the XRD patterns of the coatings after oxidation at 800 °C. The X-ray diffraction results show that the oxidation products of the CoNiCrAlY coatings prepared with different propellant gases are only α-Al₂O₃. However, the peak intensity of α-Al₂O₃ gradually increases with an increase in oxidation time. Additionally, α-Al₂O₃ is not detected during the first 20 h and 50 h of oxidation at 800 °C, which is probably due to the limited amount of oxidation products that formed, which could not be detected.

Figure 7 illustrates the surface morphologies of the coatings after oxidation at 800 °C. Figure 7a,d display the surface morphologies of the N2-sprayed and He-sprayed coatings after 20 h of oxidation. As seen in the enlarged image on the upper right, needle-like oxide particles grow on the coating surface, which are primarily α-Al₂O₃, consistent with the results obtained by XRD (Figure 6). However, it is evident that the surface of the N2-sprayed coating is covered by uniformly dense needle-like oxide particles after 500 h of oxidation (Figure 7c), while the He-sprayed coating still has different clusters of particles after 500 h of oxidation, with valleys between the clusters of particles. Thus, the He-sprayed coating has less oxide content under the same oxidation conditions. The needle-like oxide particles are indicative of the formation of θ-Al₂O₃ based on morphological analysis. Previous studies have also reported that θ-Al₂O₃ is present in the early stage of oxidation, which is confirmed by the oxide morphologies. However, it is confirmed by the XRD results in Figure 6 that the oxidation products are all α-Al₂O₃, indicating that the conversion of θ-Al₂O₃ to α-Al₂O₃ has been essentially completed on the coating surface. Nevertheless, this whisker-like structure is still observed on the coating surface because the surface morphology has not completely transformed [26]. Tolpygo et al. [27] have also previously identified θ-Al₂O₃ as a transitional oxide. Although it can quickly transition to steady-state α-Al₂O₃, this phase transition does not lead to a rapid transformation of the whisker structure, which is maintained for a longer period.

In addition to the surface morphologies shown earlier, the coatings also undergo significant microstructural evolution through depth during the oxidation process. Figure 8 shows the cross-sectional morphology and elemental surface sweeps of the CS CoNiCrAlY coating after 50, 200, and 500 h of oxidation at 800 °C. It can be seen that after 50 h of high-temperature thermal exposure, a uniform and continuous thermally grown oxide (TGO) layer was formed on the surface of both coatings, and the analysis combined with the EDS surface scan (Figure 8g) showed that the oxide layer was primarily an aluminum oxide scale. With increasing time, the TGO scale of both coatings grows significantly, and the oxide layer gradually thickens. However, the microstructure remains similar, dominated by α-Al₂O₃. Meanwhile, a more detailed observation of Figure 8g shows that a small but very obvious amount of Y₂O₃ is found at the interface of TGO/coating. It is concluded that Y as an active element can promote the formation of peg-like oxide at the oxide film interface, thus improving the adhesion of the oxide film through the mechanical keying effect (or pegging effect) [28].

From the cross-sectional elemental surface sweep of the coating after oxidation (Figure 8g), it can also be seen that the formation and growth of the α-Al₂O₃ oxide layer on the surface of the coating consumes Al in the coating, and an Al depletion zone occurs inside the coating below the oxide layer, leading to the formation of a β-depletion zone. It is clearly seen from Figure 8 that three different characteristic regions of Al distribution are found in the cross-section of the oxidized coating [29]: (1) a region of high Al concentration present on the surface (oxide layer); (2) a significant decrease in Al concentration in the region below the oxide layer due to the depletion of the β-phase (Al depletion region); and (3) a distinct Al peak and valley in the lowermost region, indicating that the β-phase is randomly distributed in the γ-phase. With an increase in oxidation time, the depth of the β depletion zone becomes significantly larger as the thickness of the oxide layer gradually increases.

Figure 9 shows the XRD patterns of N₂-sprayed and He-sprayed coatings after oxidation at 900 °C. Similar to the oxidation results at 800 °C, the oxidation products of both coatings can be seen to be primarily α-Al₂O₃ throughout the oxidation process. It can also be seen that no oxidation products were detected in the coating samples at the early stage of oxidation, such as after 20 h of oxidation; both samples showed the detection of the oxide α-Al₂O₃ after 50 h of oxidation, and with an increase in oxidation time, the intensity of the α-Al₂O₃ related peaks also increased in intensity with increasing oxidation time.

Figure 10 shows the surface morphology of the coating oxidized at 900 °C for 50, 200, and 500 h. It can be observed that after 50 h of oxidation, the surface morphology of the coating changed significantly. Compared with the oxidized surface of the coating at 800 °C under the same oxidation time, the surface morphology differed significantly. From the magnified image, it is evident that there is obvious granular alumina on the surface of the coating in addition to the needle-like alumina. However, upon comparing the same coating at different oxidation times, it can be observed that after 50 h of oxidation (Figure 10a,d), the surface of the coating is dominated by needle-like alumina, and as the oxidation time increases, oxide particles of this morphology gradually disappear, and the surface of the coating is dominated by granular oxide particles. When the oxidation time reaches 500 h, the coating is almost entirely covered by granular alumina. In contrast, the evolution of the oxide morphology on the coating surface after oxidation of the helium coating is slower than that of the nitrogen coating. It can be seen that after 500 h of oxidation (Figure 10f), a small amount of whisker-like oxide particles still exist on the coating surface. The transformation from θ-Al₂O₃ to α-Al₂O₃ is also reflected in the different surface morphology change processes.

Figure 11 depicts the cross-sectional morphology of the coating after oxidation at 900 °C. After 500 h of oxidation, the oxide skin on the surface of the coating is well-preserved, and no fracture is observed. However, the oxide layer of the nitrogen coating is thicker and less uniform compared to that of the helium coating. Moreover, small “interfacial oxide pegs” appear at the TGO/coating interface, and since this oxide is at the interface, the characteristic peaks of this oxide are not observed in the XRD spectra. Nevertheless, the results of the EDS analysis indicate that these small white particles are Y₂O₃. In general, these excellent oxide pegs are believed to be beneficial for their adhesion.

Figure 12 shows the XRD spectra of all coatings after oxidation at 1000 °C for various durations. The results reveal that the oxidation products of the helium-sprayed coatings consist mainly of α-Al₂O₃ throughout the entire oxidation process. In contrast, the oxidation products of the nitrogen coating, although mostly α-Al₂O₃, differ from the previous oxidation results at 800 °C and 900 °C in that a small amount of spinel oxide is also present. As seen in Figure 12a, the diffraction peaks of spinel oxide appear only after 500 h of oxidation, but the low intensity of its characteristic peaks and the absence of other diffraction peaks suggest that the content of this oxide is minimal.

Figure 13 shows the surface morphology of the oxide skin after oxidation at 1000 °C for different lengths of time. It can be observed that the surface of the coating is still an aggregation of various oxide particles. Compared to the results of oxidation at 800 °C and 900 °C, the morphology of the oxide does not change much, but rather, the rate of the oxide morphology change is more rapid. As can be seen in Figure 13c, the surface of the nitrogen coating is completely covered with granular oxides. However, the magnified image in Figure 13f shows that there is still a certain amount of needle-like oxide on the surface of the coating. This indicates that the oxidation rate of the helium coating is lower than that of the nitrogen coating.

Figure 14 displays the cross-sectional morphology of the CS CoNiCrAlY coating after oxidation at 1000 °C for 50, 200, and 500 h. It is evident that the thickness of the TGO layer of the coating has increased significantly for the same oxidation time compared to the cross-sectional morphology at 800 °C and 900 °C. At the initial 50 h of oxidation (Figure 14a,b), the thickness of the oxide layer of the coating is already substantial, and the oxide layer of the coating continues to thicken with increasing oxidation time. There are noticeable white Y₂O₃ particles at the interface between the oxide layer and the coating, which greatly affect the adhesion of the oxide layer.

3.3. Oxidation Kinetics

The thickness of the oxide film was measured by cross-sectional SEM micrographs of the oxidized samples to assess the overall oxide growth rate. Figure 15 shows the fitted curves of TGO thickness with increasing oxidation time versus oxide film for both coatings at 800 °C, 900 °C, and 1000 °C. It can be found that the thickness of the coating oxide film satisfies the parabolic relationship with time according to Wagner’s theory [30], that is, the square of the oxide film thickness y (μm) is inversely proportional to the time t (hour) as shown in formula (1).

$$y^2=k_{p}t+C$$

(1)

where k_p is the parabolic velocity coefficient and C is a constant.

The oxidation behavior of the two coatings at different temperatures is similar, and the increase in the thickness of the oxide film is approximately parabolic with time. Figure 11a–c show the results of TGO thickness gain for both coatings during oxidation at 800 °C, 900 °C, and 1000 °C, respectively. When integrated with the results of curve fitting, the TGO thickness curves can be viewed as having two stages: (1) at the early stage of oxidation, the oxidation rate is high in the first 50 h of oxidation, and both coatings are in the fast oxidation stage; and (2) at the middle and late stages of oxidation, their growth rate decreases significantly and they are in the slow oxidation stage. This phenomenon is due to the formation of alumina as a protective oxide, which prevents further reactions of oxygen and thus reduces the overall oxidation rate.

In the case of oxidation at 800 °C, as shown in Figure 11a, the oxide film thickness of the N₂-sprayed coating after 500 h oxidation is 2.6 μm, while that of the He-sprayed coating is 2.0 μm. Thus, it is indicated that the He-sprayed coating has a lower oxidation growth rate. At 900 °C, as shown in Figure 11b, the oxide films thicken in the order of the He-sprayed coating and the N₂-sprayed coating, and their final thicknesses after 500 h are 2.13 μm and 3.03 μm, respectively. During isothermal oxidation at 1000 °C, as shown in Figure 11c, the He-sprayed coating has the smallest increase in film thickness of about 2.5 μm, while the oxide film thickness of the N₂-sprayed coating reaches approximately 4.32 μm. In summary, the oxide film thickness of the same coating increases significantly with increasing temperature, and the oxide growth rate of the helium-sprayed coating is consistently lower than that of the nitrogen-sprayed coating at different temperature conditions, as can be seen from the coatings prepared with different parameters. This suggests that helium-sprayed CoNiCrAlY coatings can provide better oxidation performance. For this scenario, the differences in microstructure and grain size of the cold-sprayed CoNiCrAlY coatings with different preparation parameters contribute to the different oxidation behavior. These will be discussed further in the following sections.

First, the microstructure of the coatings shows a more significant difference in porosity between the two coatings, and in general, the coating porosity is more significant for the mass gain, as it represents the additional surface area where oxidation occurs. P. Richer et al. [12] also experimentally verified that the coating oxide growth rate increases significantly with increasing porosity. As shown in Figure 2, the difference in oxidation behavior becomes significant as the helium-sprayed coating is extremely dense and almost porosity-free (no porosity), while the nitrogen coating has a relatively large number of pores. Secondly, the grain size of the coating also has a specific effect on the oxidation process. Considering that the oxidation of CoNiCrAlY coatings is the diffusion of Al elements to the coating surface in combination with oxygen, the diffusion rate affects the oxidation growth; whereas the previous EBSD demonstrates that helium coatings have smaller grain size, this smaller grain size and more grain boundaries, which slows down the diffusion rate of Al elements and ultimately the oxidation rate of the coating. In summary, the low porosity and smaller grain size of the helium coating both ensure better oxidation resistance.

3.4. Oxidation Mechanisms

Figure 16 shows a schematic diagram of the oxidation process of the cold-sprayed CoNiCrAlY coating. It can be clearly seen in the figure (Figure 5b,e) that the coating is a state where the β-NiAl phase is uniformly distributed in the γ phase, and the Al-rich β-phase as an aluminum reservoir is the core region where oxidation occurs.

As shown in Figure 16a, among all the elements of the coating (Co, Ni, Cr, Al), the Gibbs generation free energy of aluminum oxide generated from aluminum is the lowest and the minimum oxygen partial pressure required is also the lowest [31,32], so the Al element in the coating will quickly diffuse to the surface of the coating and combine with oxygen to form a dense layer of aluminum oxide on the surface (Figure 16b). At the same time, as shown in Figure 16b, the inner surface of the coating will drop sharply in a short period of time due to the diffusion of Al to the extent that an Al-depletion zone appears, while the Al below the Al-depletion layer will continue to diffuse to maintain the Al concentration on the inner surface of the coating. It is known that this dense, continuous TGO layer acts as a protective oxide to inhibit further oxidation of the coating, and, therefore, the Al consumption rate decreases with time. This process corresponds to the second stage of the oxidation process discussed earlier, which is the (mid-oxidation stage) slow oxidation stage. After a long period of oxidation, as shown in Figure 16c (third stage), the growth of the TGO layer leads to a reduction in Al elements on the inner surface of the coating, a significant increase in the thickness of the Al-depletion layer of the coating, and an increase in the diffusion path of Al elements in the coating, while the formation of the aluminum oxide film causes the TGO/coating interface, and, therefore, the growth rate of alumina decreases. The oxidation model diagram can also explain why the oxidation growth process of the coating follows a parabolic law (Figure 15). In addition, according to the periodic table of elements, Cr, Ni, and Co have similar atomic radii as Al elements (Cr: 1.85 Ǻ; Ni: 1.62 Ǻ; Co: 1.67 Ǻ; Al: 1.82 Ǻ), especially Cr³⁺ and Al³⁺ have the same valence bond, making Cr₂O₃ and Al₂O₃ almost miscible at high temperatures, thus leading to Ni, Cr, and Co elements to be able to diffuse outward in the grain boundaries of the Al₂O₃ film at high temperatures. The final observation of the oxidation model (Figure 16c) shows that the thickness of the Al-depletion layer in the coating is so large that it cannot maintain a suitable aluminum concentration by diffusion, resulting in an extremely low growth rate of the alumina film and almost no increase in thickness; however, elements such as Cr, Ni, and Co can continue to diffuse, and when the oxidation time is sufficient, a large amount of Cr, Ni, and Co will diffuse to the surface of the alumina film, and a mixture of NiO, Cr₂O₃, and some spinel oxides will be formed on the surface under the high oxygen partial pressure environment, such as NiCr₂O₄, NiAl₂O₄, CoAl₂O₄, CoCr₂O₄, and NiCo₂O₄, or an alternative solid solution of (Ni, Co)(Al, Cr)₂O₄, will be formed on the surface [33]. At this time, Cr, Ni, and Co oxides distinguished from alumina can be observed on the surface, and this phenomenon was verified in the oxidation study at 1000 °C. The XRD results after 500 h of oxidation (Figure 12) indicate the presence of diffraction peaks of spinel oxides, while at 800 °C and 900 °C, due to the lower temperatures, the oxides still have not reached this stage after 500 h of oxidation.

4. Conclusions

In this study, CoNiCrAlY coatings were successfully deposited on a nickel-based superalloy IN625 substrate using a high-pressure cold spray system. The microstructure and oxidation properties of the CoNiCrAlY coatings were systematically investigated. The main conclusions are as follows:

(1)

Microstructural analysis demonstrated that both CoNiCrAlY coatings, prepared with N2 and He carrier gases, exhibited dense microstructures and good interfacial bonding. However, the porosity of the helium coatings (<0.1%) was significantly lower than that of the nitrogen coatings (~0.5%), indicating that helium coatings had better densification and lower porosity.

(2)

EBSD analysis showed that both deposited coatings underwent severe plastic deformation during the deposition process, followed by stress relief within the coatings after heat treatment. The grain size of the helium-sprayed coating (0.77 μm) was slightly smaller than that of the nitrogen-sprayed coating (0.83 μm), but both coatings exhibited very fine grains.

(3)

The oxidation results revealed that the oxidation degree of both coatings increased significantly with rising oxidation temperature (800 °C, 900 °C, and 1000 °C), with oxidation depth increasing as the temperature increased. Under the same oxidation conditions, the helium coating showed better oxidation resistance than the nitrogen coating, with improvements of 23%, 29%, and 42%, respectively. Both coatings followed the Wagner parabolic law, with the thickness of the oxide film increasing parabolically with time. The oxidation rate of the coating after heat treatment was lower than that of the as-sprayed coating. The TGO thickening curve was related to the stages of the oxidation process. Factors such as the reduction in the aluminum reservoir inside the coating, the rate of elemental diffusion, and the type and amount of oxide formation on the surface influenced the oxidation process.