Hot-Corrosion Behaviour of NiCoCrW Superalloy Fabricated by Selective Laser Melting

Abstract

In this study, a NiCoCrW (K447A) superalloy was fabricated using selective laser melting (SLM). The hot-corrosion behaviour of the SLM-fabricated NiCoCrW (K447A) was investigated in Na₂SO₄ at 900 °C. Microstructural analysis and hot-corrosion testing revealed its anisotropic behaviour in different build directions. The corrosion resistance along the build direction (XOY plane) was superior to that perpendicular to the build direction (YOZ plane). The stability of grain boundaries in the XOY plane is enhanced by the fine grains and a larger carbide area fraction, effectively hindering the diffusion of S and O₂ at grain boundaries. Furthermore, the more uniform distribution of carbides in the XOY plane reduces the local stress concentration, contributing to the stability of the protective oxide film. The enhanced hot-corrosion resistance is attributed to the formation of a more continuous and dense corrosion layer on the XOY surface.

1. Introduction

Nickel-based superalloys are extensively used in aerospace, nuclear power, and industrial gas turbines [1,2] due to their excellent mechanical properties, high strength, high hardness, and resistance to high temperatures [3,4,5]. However, traditional manufacturing processes face significant challenges in forming complex components for these applications. Additive manufacturing (AM) offers an effective solution to these limitations, providing new possibilities for geometric design and multi-material development [6,7].

During the operation of engines and gas turbines, substances such as NaCl and Na₂SO₄ are typically formed on component surfaces due to fuel combustion, leading to hot corrosion [8,9]. This degradation occurs primarily due to the diffusion of elements such as oxygen and sulphur through grain boundaries or cracks. These elements undergo complex chemical reactions with the matrix, accelerating oxidation and reducing component service life [10].

The layer-by-layer deposition and repeated thermal cycling characteristic of laser AM result in significant microstructural anisotropy in different build directions [11,12]. The differential segregation of elements at interfaces and variations in crystal structure lead to distinct microstructures and reactivity [13,14], ultimately affecting the mechanical properties and hot-corrosion resistance of the material [15,16,17]. Previous studies have demonstrated the influence of the crystallographic orientation [18,19,20], crystal structure [21,22,23], and γ′ precipitate phase [24,25] on the hot-corrosion resistance of nickel-based superalloys.

While advancements in equipment technology and materials science have addressed issues such as cracking and the formation of defects in AM structures, the inherent specificity and anisotropy of the microstructure introduce randomness in hot-corrosion behaviours [2,7]. This variability hinders the development of targeted control strategies. Therefore, in-depth research on the hot-corrosion performance of AM nickel-based superalloys is crucial to enhance their practical application value.

In this study, a NiCoCrW (K447A) superalloy fabricated by SLM was subjected to hot-corrosion testing at 900 °C using sodium sulphate as the corrosive medium. This research aims to establish a correlation between the manufacturing process, microstructure, and hot-corrosion performance to inform the design of AM processes that yield nickel-based superalloys with exceptional hot-corrosion resistance.

2. Materials and Methods

The studied superalloy powder used in this study was produced by vacuum gas atomisation. As shown in Figure 1a, the powder exhibited a smooth, defect-free surface morphology with a relatively uniform particle size distribution primarily ranging from 15 to 53 μm and high sphericity with few satellite particles. The chemical composition of the powder is presented in

Table 1. Chemical composition of NiCoCrW (K447A) superalloy powders.

Elements	Co	Cr	Mo	Al	Ti	C	Ta	Hf	W	Ni
wt.%	9.70	8.41	0.70	5.46	1.18	0.16	3.13	1.46	10.50	Bal.

. Selective laser melting (SLM) was performed using an AFS-M260 system(Longyuan AFS Co., Ltd., Beijing, CN). Prior to fabrication, the oxygen content in the forming chamber was reduced to 500 ppm through suction and purging. Argon was used as a protective gas during the SLM process to prevent oxidation. The printing parameters were set as follows: laser power of 225 W, scanning speed of 700 mm/s, scanning spacing of 0.07 mm, and layer thickness of 0.04 mm. Multiple 10 mm × 10 mm × 10 mm samples were successfully fabricated using this SLM process.

The material deposition strategy and sample sampling locations are shown in Figure 2. The surfaces of the SLM-fabricated samples were polished using SiC sandpaper with grit sizes of 400, 800, and 1200 to achieve a smooth finish (Figure 2). The samples were then ultrasonically cleaned in ethanol for 30 min to remove any residual polishing debris and dried. The weight of each sample before the hot-corrosion experiment (m1) was recorded using a Sartorius BP211D electronic balance (Sartorius, Göttingen, Germany) with a precision of 0.01 mg.

Hot-corrosion tests were conducted in a muffle furnace using corundum crucibles. Na₂SO₄ was employed as the corrosive salt at a temperature of 900 °C for a duration of 100 h. The samples were initially placed on a preheated metal plate at approximately 180 °C. A saturated Na₂SO₄ solution was then applied evenly to the sample surfaces using an atomising spray gun (Manoli, Dongguan, CN). Upon the evaporation of the water, this resulted in the formation of a dense Na₂SO₄ salt film with a mass per unit area of 0.5 mg/cm².

To study the hot-corrosion kinetics, a non-continuous weighing method was employed. The samples were removed from the furnace every 10 h, air-cooled, and weighed using the electronic balance (m2). The mass gain (Δm) was calculated as Δm = m2 − m1 to maintain a consistent salt melt environment, Na₂SO₄ was reapplied to the samples every 20 h. The calculation of the mass gain per unit area followed the following formula:

$$\mathsf{\Delta }G={\displaystyle \frac{\mathsf{\Delta }m}{A}}$$

(1)

where ΔG (mg/cm²) represents the mass gain per unit area, Δm is the mass gain, and A denotes the surface area (cm²).

The microstructure and corrosion layer morphology of the samples were characterised using an FEI Quanta 650 field-emission scanning electron microscope (FEI Company, Hillsboro, OR, USA) equipped with energy-dispersive X-ray spectroscopy (EDS). To ensure the integrity and conductivity of the corrosion layer during cross-sectional observation, the samples were coated with carbon, encapsulated in epoxy resin, and then ground and polished. Phase analysis of the corrosion products was performed using a Smartlab SE X-ray diffraction (Rigaku Corporation, Tokyo, Japan). Electron backscatter diffraction (ULVAC-PHI, Kanagawa, Japan) was employed to analyse the crystal orientation, grain size, and grain boundaries of the samples. The EBSD was analysed using TSL-OIM software (v 7.2). Further microstructural analysis of the NiCoCrW (K447A) superalloy was conducted using an FEI Technai G2 F20 transmission electron microscopy (FEI Company, Hillsboro, USA).

3. Results

3.1. Microstructural Characterisation

Figure 3 shows the microstructure of the studied superalloy. The XOY plane (Figure 3a,b) is primarily composed of equiaxed grains and white particles, whereas the YOZ plane (Figure 3c,d) contains columnar grains and white particles. EDS analysis of the white particles revealed an enrichment of carbon, indicating they are carbides. The elemental composition of these carbides is provided in

Table 2. EDS results for elements (wt.%) of the composition at each point in Figure 3.

Spectrum	C	Cr	Mo	Al	Ti	Ta	Hf	W	Ni
1	13.73	8.07	1.01	3.53	2.38	6.78	7.17	14.45	42.88
2	13.12	8.84	1.27	3.76	2.66	8.42	2.30	14.88	44.75
3	10.90	8.68	0.86	4.41	1.20	4.72	5.05	9.99	54.19
4	12.87	8.01	0.62	3.75	1.93	8.21	3.99	13.05	47.57
5	12.25	7.30	0.61	3.62	2.71	11.03	5.38	14.28	42.82
6	11.63	7.47	1.41	3.63	2.84	8.96	10.07	10.88	43.11

. These microstructural observations are consistent with previous studies on NiCoCrW superalloys [3,4]. The area fraction of carbides in the samples with different building directions was calculated using Image J (v 1.8.0). The results indicate that the area fractions of carbides in the XOY and YOZ planes are approximately 9.12% and 5.16%. TEM analysis was conducted to further investigate the microstructural characteristics and the carbides.

Figure 4 shows the carbide element segregation and dislocation distribution characteristics of the studied superalloy. Analysis of selected area electron diffraction patterns and EDS revealed that the carbides are MC-type carbides (M = Hf, Ti). The carbides in the XOY plane exhibit a honeycomb distribution, while those in the YOZ plane are distributed in a band-like pattern. Carbides in nickel-based superalloys can pin dislocations, enhancing the material’s strength. Their stability at high temperatures also contributes to improved hot-corrosion resistance. Therefore, the XOY plane, with its finer grain structure and more dispersed carbides, is expected to exhibit enhanced microstructural stability.

The inverse pole figure (IPF) maps in Figure 5a,d show that the grains in the XOY plane tend towards orientation, while the grains in the YOZ plane are also close to orientation. The illustrations in Figure 5a,d demonstrate that the average grain size in the XOY plane is 3.01 μm, while in the YOZ plane, it is 4.40 μm. Consistent with the SEM and TEM observations, the grain size in the XOY plane is finer. The grain boundary maps in Figure 5b,e show that the XOY plane has a larger grain boundary area due to its finer grains. The kernel average misorientation (KAM) maps in Figure 5c,f indicate a higher local stress concentration in the YOZ plane than that in the XOY plane.

3.2. Hot-Corrosion Test Results

Figure 6 shows the surface macromorphologies of the XOY and YOZ planes after hot corrosion. The samples with different build directions exhibit different degrees of hot corrosion. After 10 h of hot corrosion, the YOZ samples showed extensive bulging on the surface, while the XOY plane remained relatively flat. After 20 and 50 h of hot corrosion, the surface roughness of the XOY plane remained relatively unchanged, whereas that of the YOZ plane specimens increased with corrosion time. After 100 h of hot corrosion, the corrosion layer on the XOY plane exhibited a better bond with the matrix, whereas the bulging parts on the YOZ plane experienced severe spalling.

The kinetic curves of the studied superalloy after corrosion for 100 h in molten Na₂SO₄ at 900 °C are shown in Figure 7. The mass gain due to hot corrosion at every stage on the YOZ plane is consistently higher than that on the XOY plane. After 100 h of corrosion, the unit area mass gains for the studied superalloy on the XOY and YOZ plane samples are 12.13 mg/cm² and 18.56 mg/cm², respectively. The average corrosion rate of the YOZ plane is 1.5 times that of the XOY plane. The hot-corrosion kinetics of the studied superalloy follows the parabolic law, and the hot-corrosion rate of the studied superalloy is represented by Equation (2) [26,27]. By comparison to other used nickel-based superalloys [18,20], both the XOY and YOZ plane samples demonstrate superior resistance to hot corrosion. The superior hot-corrosion resistance of nickel-based superalloys prepared by SLM is attributed to their better microstructural stability at high temperatures. These findings imply that the studied superalloy is well suited for applications demanding high hot-corrosion resistance.

$$k_p={\displaystyle \frac{∆G^2}{t}},$$

(2)

where t (h) is the time of hot corrosion, and k_p (mg²·cm⁻⁴/h) is the hot-corrosion rate.

3.3. Characterisation of Hot-Corrosion Products

The XRD phase analysis of corrosion products is shown in Figure 8. After hot corrosion for 100 h in Na₂SO₄ at 900 °C, the corrosion products were mainly composed of oxides. At 900 °C, in addition to the Ni element, corrosion reactions also occurred for Al, Cr, Ti, and Co within the alloy. Owing to the consistent material composition, the main products of hot corrosion in samples with different building directions featured similar phase compositions. However, the microstructures of the samples with different building directions differed, and their abilities to form dense oxide films on the surface varied, leading to different relative contents of corrosion products. As shown in Figure 6 and Figure 7, the YOZ plane shows more severe corrosion, resulting in higher diffraction peak intensities of the hot-corrosion products than those in the XOY plane. And the corrosion layer on the YOZ plane experienced severe spalling, leading to the exposure of the matrix. Consequently, the intensity of the diffraction peaks for the matrix (γ/γ′) phase in the YOZ plane is higher.

Figure 9 shows the microscopic surface morphology of the XOY and YOZ plane samples after corrosion for 100 h in Na₂SO₄ at 900 °C. The corrosion layer on the XOY plane is relatively complete (Figure 9a), with no obvious spalling, thereby providing a certain degree of protection for the studied superalloy. The high-magnification image shown in Figure 9b shows that the oxide layer consists of dense spinel and loose clusters. Combined with XRD analysis, the dense spinel is identified as (Al,Cr)₂O₃, while the loose porous products are NiO, Na₂CrO₄, and NaNiO₂. Severe spalling of the oxide layer, along with a higher prevalence of cracks and pits, is observed on the YOZ specimen surface, as shown in Figure 9c. The YOZ plane featured more cracks and pits compared to the XOY plane. Thus, the YOZ plane exhibited more severe corrosion, which is consistent with the corrosion rates depicted in Figure 7.

3.4. Cross-Sectional Morphologies of Corrosion Layer

Figure 10 shows the cross-sectional morphology and EDS results of the studied superalloy after hot corrosion for 100 h in Na₂SO₄ at 900 °C. The corrosion layer on the XOY plane is denser and more tightly bonded to the matrix, while the corrosion layer on the YOZ plane exhibits poor adhesion to the matrix. Figure 10a shows that the corrosion layer on the XOY plane is mainly composed of Al₂O₃ and Cr₂O₃. The dense Al₂O₃ layer can effectively prevent the corrosion of the matrix by Na₂SO₄. Compared to the XOY plane, the enrichment of Ni and Co within the corrosion layer on the YOZ plane is more pronounced, indicating the spalling of the Al₂O₃ oxide and the subsequent exposure of the matrix to Na₂SO₄. Additionally, the enrichment of S within the oxide layer on the XOY plane indicates that the oxide layer is more dense, preventing the diffusion of S into the matrix. Thus, S and O infiltration into the YOZ plane matrix increased the corrosion severity in the YOZ plane.

4. Discussion

4.1. Hot-Corrosion Mechanisms

The kinetic curves in Figure 7 reveal that the hot corrosion of the studied superalloy can be divided into two stages: an initial rapid corrosion stage (0–10 h) and a subsequent stage (10–100 h) with a more stable mass gain. This behaviour is attributed to the initial interaction of Na₂SO₄ with the sample surface, leading to rapid hot-corrosion reactions. As corrosion progresses, a protective oxide layer forms on the surface, impeding direct contact between the corrosive medium and the superalloy. Consequently, the corrosion mechanism transitions from a surface-controlled reaction to a diffusion-controlled reaction, resulting in a reduced corrosion rate [28,29,30].

At the test temperature of 900 °C, Na₂SO₄ decomposes according to the following reaction [21]:

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

(3)

The essence of hot corrosion is the outward diffusion of alloy elements and the inward diffusion of O and S into the matrix. Grain boundaries facilitate the diffusion of alloy elements, promoting the formation of a dense oxide film on the sample surface, which enhances resistance to hot corrosion. However, grain boundaries also serve as diffusion pathways for S and O₂, potentially leading to more severe corrosion of the matrix. The reaction between alloy elements and O₂ is shown in reaction (4) [31,32,33].

Ti and Hf are primarily present as carbides at the grain boundaries. HfC is relatively stable and does not readily participate in corrosion reactions. However, when O₂ diffuses through the grain boundaries, TiC reacts to form TiO₂, as shown in reaction (5) [24]. The standard Gibbs free energy of oxide formation at 900 °C follows the order Al₂O₃ < Cr₂O₃ < TiO₂ < CoO < NiO, as shown in Figure 11a.

$$\mathrm{x}\mathrm{M}+{\displaystyle \frac{y}{2}}{O}_2\to {\mathrm{M}}_{\mathrm{x}}{\mathrm{O}}_{\mathrm{y}}\left(\mathrm{M}=\mathrm{N}\mathrm{i},\mathrm{C}\mathrm{r},\mathrm{C}\mathrm{o},\mathrm{A}\mathrm{l},\mathrm{T}\mathrm{i}\right),$$

(4)

$${\mathrm{T}\mathrm{i}\mathrm{C}+\mathrm{O}}_2\to {\mathrm{T}\mathrm{i}\mathrm{O}}_2,$$

(5)

The reaction between S and alloy elements is shown in reaction (6). The standard Gibbs free energy of metal sulphides at 900 °C follows the order Al₂S₃ < TiS₂ < Cr₂S₃ < CrS, as shown in Figure 11b. Therefore, S reacts preferentially with Al. When Al is consumed to a certain extent, Ti reacts with S, and eventually, Cr undergoes sulphidation. As shown in Figure 12, metal sulphides can form oxides by reacting with O₂ (Reaction (7)) when the O₂ partial pressure is sufficiently high. S and O₂ have a synergistic effect in the hot-corrosion process. As shown in reaction (3), the O₂ entering the corrosion layer and the O₂ released from Na₂SO₄ together lead to an increase in the oxygen partial pressure in the matrix, promoting the oxidation of metal sulphides.

$$x\mathrm{M}+y\mathrm{S}\to \mathrm{M}x\mathrm{S}y\left(\mathrm{M}=\mathrm{T}\mathrm{i},\mathrm{C}\mathrm{r},\mathrm{A}\mathrm{l}\right),$$

(6)

$$\mathrm{M}\mathrm{x}\mathrm{S}\mathrm{y}+{\displaystyle \frac{y}{2}}{\mathrm{O}}_2\to {\mathrm{M}}_{\mathrm{x}}{\mathrm{O}}_{\mathrm{y}}+\mathrm{y}\mathrm{S}\left(\mathrm{M}=\mathrm{T}\mathrm{i},\mathrm{C}\mathrm{r},\mathrm{A}\mathrm{l}\right),$$

(7)

The Na₂O generated from the decomposition of Na₂SO₄ is an important factor contributing to the severity of hot corrosion in the superalloy. The protective Cr₂O₃ oxide film can dissolve in the presence of Na₂O to form NaNiO₂ and Na₂CrO₄, as shown in reaction (8) [34]. In summary, the diffusion of S and O₂ through grain boundaries, coupled with the dissolution of the protective oxide film by Na₂O, leads to the hot corrosion of the studied superalloy.

$$\mathrm{C}{\mathrm{r}}_2{\mathrm{O}}_3+\mathrm{N}{\mathrm{a}}_2\mathrm{O}\to \mathrm{N}{\mathrm{a}}_2\mathrm{C}\mathrm{r}{\mathrm{O}}_4\to \mathrm{C}{\mathrm{r}}_2{\mathrm{O}}_3+\mathrm{N}\mathrm{a}\mathrm{N}\mathrm{i}{\mathrm{O}}_2.$$

(8)

4.2. Microstructure Anisotropy Influence on Hot Corrosion

The hot-corrosion products on the XOY and YOZ sample surfaces are similar, consisting primarily of Al₂O₃ and Cr₂O₃, indicating that the fundamental hot-corrosion mechanisms are the same. During the SLM process, different building directions lead to differences in microstructure. The XOY plane sample has a finer grain size, a higher area fraction of carbides, and a more uniform carbide distribution (Figure 4 and Figure 5). Finer and more uniform grains typically reduce the stress concentration, thereby lowering the risk of corrosion cracking [18,19,23]. Grain boundaries are pinned by carbides, and the stability of grain boundaries is increased by a higher carbide area fraction, thereby reducing the stress concentration. The matrix deforms more uniformly during hot corrosion due to the uniform distribution of carbides, thereby reducing local stress, enhancing the adhesion of the oxide film, and reducing the spalling behaviour of the oxide film, thus improving the resistance to hot corrosion. Thus, the hot-corrosion performance differs significantly between the two build directions due to their distinct microstructures.

The hot-corrosion schematic of the studied superalloy in Na₂SO₄ at 900 °C is shown in Figure 13. S and O₂ diffuse through the grain boundaries into the matrix, where they selectively react with alloying elements. TiC and HfC precipitate at the grain boundaries, and these carbides exhibit high stability. Moreover, Hf and Ti have higher oxidation resistance than Al, Ni, and Co. As shown in Figure 4, the area fraction of carbides on the XOY plane is larger, effectively slowing down the hot-corrosion rate. As hot-corrosion progresses, a dense oxide film (Al₂O₃ and Cr₂O₃) forms on the sample surface. Due to the low solubility of Na₂SO₄ in this oxide film, it acts as a protective barrier to the matrix, reducing the rate of hot corrosion. The finer grain size and more uniform distribution of carbides on the XOY plane result in a lower local stress concentration than that in the YOZ plane, contributing to the structural stability of the oxide film. The grain refinement increases the uniformity of the microstructure, reducing local stress and decreasing defects at the grain boundaries, which in turn reduces the diffusion paths for corrosive media. A higher carbide area fraction pins the grain boundaries, preventing slippage and promoting the formation of a dense oxide film, thereby enhancing the resistance to hot corrosion. The uniform distribution of carbides allows the matrix to deform more uniformly, reducing local stress and improving the adhesion of the oxide film, thus reducing spalling behaviour. These factors collectively act to significantly reduce the rate of hot corrosion and enhance the hot-corrosion resistance. The dissolution of Cr₂O₃ can lead to the formation of holes or cracks in the corrosion layer (reaction (8)). S and O₂ can then diffuse through these defects to the substrate surface, where they react with the matrix, leading to further hot corrosion. Therefore, the matrix is primarily protected by the Al₂O₃ layer. As shown in Figure 5f, the YOZ plane exhibits greater local stress, leading to poor adhesion between the corrosion layer and the matrix. After the severe spalling of the corrosion layer, the matrix surface is re-exposed and reacts with S and O₂ (reaction (4)–(7)), resulting in more severe corrosion on the YOZ plane.

5. Conclusions

In this study, a NiCoCrW (K447A) nickel-based superalloy was fabricated by selective laser melting (SLM), and its microstructure and hot-corrosion behaviour in Na₂SO₄ at 900 °C for 100 h were investigated. The major conclusions are summarised as follows:

• The hot corrosion of the SLM-fabricated NiCoCrW (K447A) superalloy exhibits anisotropy in different build directions. The XOY plane (parallel to the build direction) exhibits superior hot-corrosion resistance compared to that in the YOZ plane (perpendicular to the build direction).
• The microstructure of the NiCoCrW (K447A) superalloy shows significant anisotropy. The XOY plane has finer grains, a more uniform distribution of carbides, and a lower local stress concentration.
• The excellent hot-corrosion resistance of the XOY plane is attributed to the larger area fraction and more uniform distribution of carbides, which contribute to the stability of the oxide film. The oxide film formed on the XOY surface is more stable, providing better protection for the underlying matrix.

Thus, this study provides a new approach for designing hot-corrosion-resistant materials fabricated by SLM. By adjusting the deposition strategy to tailor the microstructure, it is possible to enhance the material’s oxidation resistance and hot-corrosion resistance.