Oxidation Property of a Fourth-Generation Powder Metallurgy FGH4108 Nickel-Based Superalloy

Abstract

Isothermal oxidation kinetics of a fourth-generation powder metallurgy FGH4108 nickel-based superalloy is investigated at 800 °C to 1100 °C by X-ray diffraction (XRD), scanning electron microscopy (SEM), and energy dispersive X-ray spectroscopy (EDS). At 800 °C and 900 °C, the oxidation kinetic curves of the FGH4108 superalloy follow parabolic law. At 1000 °C, the oxidation kinetic curve follows cubic law. At 1100 °C, the oxidation kinetic curve has two distinct parts: the first part follows a parabolic law, and the second one obeys a linear law. Cross-sectional morphologies and elemental distributions show that the oxide film consists of two parts at 800 °C: the outer layer is a continuous dense protective Cr₂O₃ oxide film, and the inner layer is a discontinuous Al₂O₃ oxide layer. At 900–1100 °C, the oxides consist of three layers: the outermost is the oxides of Cr₂O₃ and TiO₂, the middle is a continuous oxide of Al₂O₃, and the innermost is dotted oxides of TiO_2. The thickness of the inner TiO₂ oxide layer increases with the increase of oxidation temperature. On this basis, the oxidation behavior of the FGH4108 superalloy at high temperatures is confirmed to be controlled by the diffusion of Cr, Al, Ti, and O. From the aspect of oxidation resistance, the long-term service temperature of the FGH4108 superalloy should not exceed 1000 °C.

1. Introduction

The term superalloy refers to Ni, Fe, or Co-based materials that can work continuously for a long time under high temperatures and complex stresses and are widely used in the preparation of hot end components of aero engines and gas turbines [1,2,3,4,5]. As the aerospace industry develops, the operating temperatures of aero engines are becoming higher and higher to achieve a higher thrust-to-weight ratio and higher fuel utilization efficiency [6,7,8]. With the aim of improving the properties of the alloy, the type and content of strengthening elements in the matrix are increasing, which leads to deformation resistance and severe segregation in the process of forging and casting [9,10,11]. However, the powder superalloy has good machinability because component segregation occurs only within the particle size of the metal powder [12,13]. The powder superalloy has a uniform fine microstructure and exhibits excellent strength, fatigue properties and surface stability at high temperatures [14,15,16]. Therefore, powder metallurgy nickel-based superalloys have received increasing attention as aerospace engine materials with excellent mechanical properties and corrosion resistance for manufacturing high-performance turbine disks [17,18].

It is well known that oxidation resistance is one of the most important aspects of nickel-based superalloys since poor oxidation resistance hinders their possible applications at high temperatures [12]. The oxidation behaviors of advanced nickel-based superalloys at temperatures higher than 700 °C have been studied extensively [19,20,21,22]. Brenneman et al. [19] studied the oxidation behavior of GTD111 superalloy at 900 °C in air and concluded that the oxidation resistance was provided by the continuous dense Cr₂O₃ oxide band, rather than the discontinuous Al₂O₃ oxide layer. Liu et al. [23] studied the oxidation behavior of single-crystal nickel-based superalloy DD32 in air at 900 °C and 1000 °C and found that two oxidation steps appeared in the oxidation kinetics, which were controlled by NiO growth and Al₂O₃ growth, respectively. Berthod et al. [20] studied the oxidation behavior of six (Co + Ni)-based Cr-rich alloys with different Co/Ni ratios at 1250 °C and concluded that the higher the Ni content, the better the oxidation resistance of the alloys. Moffat et al. [21] studied the oxidation behavior of a carbide network in low or high carbon Co-based superalloys at 1000 °C. They concluded that the carbide in low carbon alloys is more soluble than that in high carbon alloys, thus reducing the continuity of the carbide network and ultimately leading to an increase in the thickness of the oxide layer. Xiao et al. [24] studied the cyclic oxidation behavior of Pt-modified γ’ nanocrystalline coating at 1150 °C and confirmed that the nanocrystalline structure promoted the formation of α-Al₂O₃ oxide coating with lower Pt and Al contents, which improved the anti-spalling performance. As a fourth-generation powder metallurgy nickel-based superalloy, FGH4108 is highly alloyed and reinforced by γ’ precipitation with a mass fraction of about 60%. Therefore, it can work under harsh conditions of temperature and load. However, its oxidation behavior needs to be studied in detail.

In this paper, the high-temperature oxidation behavior and kinetics of FGH4108 superalloy at different temperatures are investigated by isothermal oxidation tests. The oxidation mechanism is analyzed to provide a general understanding for further improvement and better utilization of the FGH4108 superalloy.

2. Materials and Methods

The composition of the FGH4108 superalloy is shown in

Table 1. Chemical compositions of FGH4108 superalloy (wt%).

Co	Cr	Al	Ti	C	B	Zr	Mo + W	Nb + Ta + Hf	Ni
18.0	12.0	3.0	3.0	0.06	0.04	0.05	6.0	7.0	Balance

. FGH4108 alloy was prepared by vacuum induction-melting gas atomization (VIGA) + hot isostatic pressing (HIP) + hot extrusion + forging + heat treatment. The relative density of the alloy was determined by the Archimedean drainage method to be 100%. The microstructure and crystallography of the alloy were investigated through optical microscopy (OM). Figure 1 shows the typical microstructure of the alloy. The grain size grade was estimated to be ASTM E112 class 6–7 using OM image processing software, ImageJ 1.52t.

The specimens were wire cut into 15 × 10 × 5 mm³ sized oxidation test specimens, and their surfaces were sandpapered to 1000 #, ultrasonically cleaned, and then dried and set aside. Before oxidation, the specimens were weighed by an electronic balance with an accuracy of 0.1 mg. The high temperature oxidation resistance was tested by the weight gain at 800 °C, 900 °C, 1000 °C, and 1100 °C in ambient conditions. The oxidation times were set as 1 h, 3 h, 5 h, 10 h, 25 h, 50 h, 75 h, and 100 h. After oxidation, the specimen was air-cooled to room temperature and weighed by the electronic balance again. The surface morphologies of the oxidation products were detected by scanning electron microscopy (SEM), and the phase constitutions were analyzed by X-ray diffraction (XRD). The XRD data was processed by the Jade 6.5 software, and the Inorganics database of the software was used. The cross section of the oxidized specimen was observed by SEM to explore the morphology of the oxidation products. The samples for the observation cross section were mechanically polished with 60 #, 150 #, 800 #, to 2000 # SiC sandpaper in turn. Energy dispersive X-ray spectroscopy (EDS) was applied to study the elemental distribution of the oxidation product films. Without a particular explanation, the results in the figures and tables in this paper all represent the average data of three samples under specific experimental conditions.

3. Results and Discussion

3.1. Oxidation Kinetics

Figure 2 shows the kinetic curves of the FGH4108 alloy after 100 h of oxidation at 800 °C, 900 °C, 1000 °C, and 1100 °C. It can be seen that the weight gain increases as the temperature rises and the time increases. After 100 h oxidation, the oxidation weight gains are 3.69 g/m², 4.41 g/m², 11.82 g/m², and 94.97 g/m² at 800 °C, 900 °C, 1000 °C, and 1100 °C, respectively.

It is known that the mass gain of nickel-based superalloy during oxidation can be described by the Wanger formula [2,25,26]

$${\left(△m/S\right)}^n=kt,$$

(1)

where $△m$ is the oxidation weight gain, S is the surface area of the specimen, k is a constant related to the oxidation rate, t is the oxidation time, and n is the oxidation rate index. Fitting equations for each temperature are shown in

Table 2. Fitting equations for each temperature.

Temperature (°C)	Time (h)	Kinetic Expressions	Correlation Coefficient R2
800	0–100	${\left(△m/S\right)}^{1.95}=0.12t$	0.924
900	0–100	${\left(△m/S\right)}^{1.85}=0.16t$	0.949
1000	0–100	${\left(△m/S\right)}^{2.85}=11.47t$	0.966
1100	0–25	${\left(△m/S\right)}^{2.16}=193.07t$	0.978
	25–100	${\left(△m/S\right)}^{1.13}=\left(1.21t+52.74\right)$	0.990

and Figure 1. It can be seen that the oxidation rate of the FGH4108 alloy is mainly temperature dependent. With the increase in temperature, the n decreases, while the oxidation rate constant k increases by one order of magnitude. At 800–900 °C, FGH4108 alloy has good oxidation resistance, and the oxidation kinetic curves follow the parabolic law. At 1000 °C, oxidation gain increases and nearly follows cubic law. At 1100 °C, oxidation gain increases significantly, and the alloy has poor oxidation resistance. Besides, the oxidation kinetic curve can be divided into two parts; the pre-oxidation period follows the parabolic law but gradually changes to the linear law with increased oxidation time. It can be seen from Table 2 that with the increase of oxidation temperature, the value of k gradually increases, indicating that the degree of oxidation increases, especially when the temperature reaches 1100 °C; so, the parabolic rule is lost.

3.2. Surface Morphology and Phase Composition of the Oxide Layer

Figure 3 displays the surface morphologies after oxidation for 100 h. It can be seen from Figure 3a that the alloy is oxidized slightly at 800 °C, and only a small amount of oxides is observed at the surface. Figure 3b exhibits that after oxidation at 900 °C for 100 h, the oxide film starts to bulge, but no crack is found. Figure 3c indicates that after 1000 °C, the oxide film is relatively loose and poorly bonded to the substrate due to lateral growth or wrinkling. Besides, a small amount of spalling occurs during cooling, revealing surface holes and grain boundary grooves. Figure 3d shows that significant flaking occurs after oxidation at 1100 °C and the oxidation resistance greatly reduces.

Figure 4 displays XRD analyses after oxidation for 5 h, 50 h, and 100 h. It can be seen from Figure 4a that only Cr₂O₃ can be identified after oxidation at 800 °C for 5 h, and a small amount of Cr₂O₃ emerges for 50 h while both TiO₂ and Cr₂O₃ appear for 100 h. Figure 4b exhibits that only Cr₂O₃ can be identified after oxidation at 900 °C for 5 h, which is the same as that at 800 °C. After oxidation for 50 h, a small amount of Cr₂O₃ also emerges, and TiO₂ can be identified. After oxidation for 100 h, a small amount of NiCr₂O₄ appears besides Cr₂O₃ and TiO₂. Figure 4c indicates that TiO₂ can be detected evidently after oxidation at 1000 °C for 5 h. When the oxidation time extends to 50 h, TiO₂ appears in addition to Cr₂O₃. After oxidation for 100 h, Al₂O₃ and NiCr₂O₄ exist, as well as Cr₂O₃ and TiO₂. Figure 4d shows that many oxidation products (TiO₂, Cr₂O₃, Al₂O₃, NiO, and NiCr₂O₄) emerge at 1100 °C even for 5 h. As the oxidation time extends, the amounts of all oxides increase while their species remain unchanged.

3.3. Cross-Sectional Analysis of the Oxide Layer

Figure 5 demonstrates the cross-sectional morphology and elemental distribution of the oxide layer after oxidation at 800 °C for 100 h. It can be seen that the oxide layer can be divided into two layers. The outermost layer is enriched in Cr due to a continuous dense protective Cr₂O₃ oxide film. The inner layer forms a black oxide of Al near the substrate as a discontinuous Al₂O₃ oxide film.

Figure 6 shows the cross-sectional morphology and elemental distribution of the oxide layer after oxidation at 900 °C for 100 h. It is seen that the oxide layer is composed of three parts. The outermost layer forms a continuous layer of dense Cr₂O₃ oxide with TiO₂ present, i.e., a mixed layer of Cr₂O₃ and TiO₂ oxide film. The middle layer is a continuous layer of Al₂O₃ oxide. The innermost layer is a dotted layer of uniformly distributed TiO₂.

Figure 7 illustrates the cross-sectional morphology and elemental distribution of the oxide layer after oxidation at 1000 °C for 100 h. It is evident that the oxide layer also includes three parts. The outermost layer is enriched with Cr, Ti, and Ni owing to a mixed layer of loose Cr₂O₃, TiO₂, and spinel oxide NiCr₂O₄ film. Many holes form in the oxide film, and the oxides appear to flake off. The middle layer is a continuous layer of Al₂O₃ oxide. The innermost layer is a dotted, uniformly distributed layer of TiO₂.

Figure 8 shows the cross-sectional morphology and elemental distribution of the oxide layer after oxidation at 1100 °C for 100 h. It is seen that the oxide layer still includes three parts. The outermost layer is enriched with Cr and Ti and is loose and porous. The Cr₂O₃ oxide film continues to form and flakes off. The middle layer is a continuous Al₂O₃ oxide with varying thicknesses at different locations. The innermost layer is a layer of TiO₂ oxide uniformly distributed in the shape of spots.

3.4. Oxidation Mechanism

3.4.1. Oxidation Kinetics

According to the oxidation kinetics, together with surface morphology and phase composition, the oxidation process can be separated into four stages (see Figure 9).

At Stage I, oxygen molecules in the air are physically adsorbed on the surface of the alloy due to van der Waals forces, followed by the decomposition of the oxygen molecules and the chemisorption of free electrons in the alloy [14]. Afterwards, the alloying elements on the surface react with O, and the oxide starts to nucleate. Due to the easy diffusion of alloying elements through grain boundaries, grain boundaries provide a favorable location for the nucleation and growth of the oxide. Notably, Ti reacts more readily with O than Cr to form oxides thermodynamically, but Cr₂O₃ oxide appears earlier than TiO₂ (see Figure 4). Actually, the Cr content in the alloy is much higher than that of Ti. As a result, the oxide detected first is Cr₂O₃. Subsequently, a mixed layer of Cr₂O₃ and TiO₂ oxide films comes into being. The oxide film produced at this stage is not dense.

At Stage II, Ti and Cr in the matrix continue to diffuse towards the outer oxide film. Simultaneously, O diffuses inwards, and the oxidation reactions continue to occur, forming a continuous dense oxide film. Due to the Cr depletion in this region, internal oxidation of Al in the alloy begins to occur.

At Stage III, the oxide film grows further inwards, which increases the thickness of the film and makes the Ti and Cr outward diffusion more difficult. At this moment, Al oxidation proceeds, and a continuous layer of Al₂O₃ oxide in the inner layer generates. At the same time, Ni reacts with O to produce an unstable NiO oxide. Afterwards, NiO reacts with Cr₂O₃ to produce NiCr₂O₄. This hinders the occurrence of internal oxidation due to the more protective NiCr₂O₄. Accordingly, a dense oxide film forms. As the reaction progresses, holes are created in the oxide layer and at the interface due to the Kirkendall effect. As the holes are created, the oxide film becomes less dense, and O continues to diffuse through the holes into the inner layer of the substrate. Consequently, the depletion of Cr and Al occurs near the substrate due to the creation of a continuous Cr₂O₃ outer oxide layer and Al₂O₃ inner oxide. Besides, Ti reacts with O in the innermost layer to create a dotted TiO₂ oxide. Therefore, this oxide is not protective since it is distributed discontinuously and thus cannot block the diffusion of O. At this stage, the oxide layer can be clarified. The outermost layer forms a continuous dense mixed layer of Cr₂O₃ and TiO₂ oxide with NiCr₂O₄ spinel oxide. The middle layer forms a continuous layer of Al₂O₃ internal oxide. The innermost layer forms a discontinuously dotted layer of TiO₂ oxide.

At Stage IV, the oxide flakes off and regrows. When the dense oxide film breaks down due to a large difference in expansion coefficient, the substrate is exposed to the air and continues to oxidize [27,28,29]. In this case, a new oxide film appears and hinders further oxidation until the oxide flakes off again. On this basis, oxidation is an iterative process of growth-flaking-regrowth-flaking of the oxide film.

3.4.2. Selective Oxidation

Figure 4 confirms that at a temperature range of 800 °C to 1100 °C, the oxides include Cr₂O₃, TiO₂, Al₂O₃, and NiCr₂O₄, which aligns with the reported results [30,31,32]. For the oxidation of metals [13,33], the Ellingham diagram is widely used to calculate the free energy, $\Delta G^0$, and describe the oxide species thermodynamically [14,34,35,36,37]. As a function of temperature, $\Delta G_{{\mathrm{Cr}}_2{\mathrm{O}}_3}^0$, $\Delta G_{{\mathrm{TiO}}_2}^0$, $\Delta G_{\mathrm{NiO}}^0$, and $\Delta G_{{\mathrm{Al}}_2{\mathrm{O}}_3}^0$ are expressed by [13,14,36,37] as

$$\frac{4}{3}\mathrm{Cr}+{\mathrm{O}}_2(\mathrm{g})=\frac{2}{3}{\mathrm{Cr}}_2{\mathrm{O}}_3(\mathrm{s}),\Delta G_T^{\theta }=-753.12+0.1826T\left(\mathrm{kJ}\text{·}{\mathrm{mol}}^{-1}\right)$$

(2)

$$\frac{4}{3}\mathrm{Al}+{\mathrm{O}}_2(\mathrm{g})=\frac{2}{3}{\mathrm{Al}}_2{\mathrm{O}}_3(\mathrm{s}),\Delta G_T^{\theta }=-1116.29+0.2088T\left(\mathrm{kJ}\text{·}{\mathrm{mol}}^{-1}\right)$$

(3)

$$\mathrm{Ti}+{\mathrm{O}}_2(\mathrm{g})={\mathrm{TiO}}_2(\mathrm{s}),\Delta G_T^{\theta }=-944.75+0.1854T\left(\mathrm{kJ}\text{·}{\mathrm{mol}}^{-1}\right)$$

(4)

$$2\mathrm{Ni}+{\mathrm{O}}_2(\mathrm{g})=2\mathrm{NiO}(\mathrm{s}),\Delta G_T^{\theta }=-497.16+0.1887T\left(\mathrm{kJ}\text{·}{\mathrm{mol}}^{-1}\right)$$

(5)

in which T denotes the temperature (K). NiCr₂O₄ is formed from NiO and Cr₂O₃ according to the following reaction.

$$\mathrm{NiO}(\mathrm{s})+{\mathrm{Cr}}_2{\mathrm{O}}_3(\mathrm{s})={\mathrm{NiCr}}_2{\mathrm{O}}_4(\mathrm{s}),\Delta G_T^{\theta }=-1378.26+0.3677T\left(\mathrm{kJ}\text{·}{\mathrm{mol}}^{-1}\right)$$

(6)

In terms of Equations (2)–(6), the free energies of the above five oxides at 800 °C to 1100 °C are calculated and shown in Figure 10. It is clear that all calculated values of the free energies are negative, indicating that all these alloying elements are likely to undergo oxidation reactions to form the corresponding oxides. Moreover, NiCr₂O₄ is formed from NiO and Cr₂O₃ spontaneously.

It is notable from Figure 10 that Al has the highest affinity to O among the alloying elements in FGH4108, and Al₂O₃ is the stablest oxide. However, Figure 4 demonstrates that Cr₂O₃ appears first. This is because the Ellingham diagram only illustrates the magnitude of the affinity of the different alloying elements for oxygen, but the type of oxides that emerge is also influenced by the elemental content factor. Although Cr₂O₃ is not as stable as Al₂O₃, the content of Cr in the alloy (12.24 wt%) is much higher than that of Al (2.96 wt%). As a result, Cr₂O₃ appears firstly by selective oxidation.

In addition, Al₂O₃ is a typical n-structured oxide, and the n-structured oxide film contains anionic vacancies and excess electrons [38]. This oxide relies on the inward diffusion of O²⁻ to the Al and the bonding reaction at the oxide film–substrate interface [39]. Hence, the Al₂O₃ oxide layer in the inner layer of the oxide film is observed by the cross-sectional analysis. Furthermore, the selective oxidation of the alloy is controlled by the diffusion of the elements, and Al diffuses much more slowly in the oxide film than Cr and Ti. This causes the presence of the Al₂O₃ oxide layer in the inner layer. Consequently, Cr and Ti diffuse to the outer layer to react first with O to form oxides, whereas Al oxidation occurs mainly in the inner layer.

3.4.3. Spalling of Oxides

As mentioned in the previous section, the oxidation behavior of FGH4108 superalloy is controlled by the diffusion of Cr, Al, Ti, and O. As shown in Figure 2, the oxidation kinetic curves present a parabolic rule at 800–1000 °C, indicating that the growth rate of the oxide film is inversely proportional to the thickness of the film. At this time, a thick protective film is formed on the metal surface, and the growth rate of the film is determined by the diffusion rate [14,25]. At 1100 °C, the front part of the oxidation kinetics curve is parabolic, and the second part is linear. Figure 11 shows the ratio of oxidation weight gain to time at 800 °C to 1100 °C. At 800–1000 °C, the ratio of oxidation weight gain to time gradually decreases with the increase of time and eventually trends towards 0 g·m⁻²·s⁻¹. This indicates that a dense oxide film is formed on the surface of the alloy, and with the gradual thickening of the oxide film, the diffusion rate of ions decreases, thus reducing the weight gain rate of oxides. At 1100 °C, although the ratio of oxidation weight gain to time also decreases with the increase of time, when the time exceeds 30 h, the ratio of oxidation weight gain to time trends towards 0.5 g·m⁻²·s⁻¹. This indicates that, in this case, the oxide film does not play an adequate role in protecting the alloy, i.e., it does not form a sufficiently dense oxide film. Therefore, oxidation weight gain increases linearly with time. From the aspect of oxidation resistance, it can be concluded that the long-term service temperature of the FGH4108 superalloy should not exceed 1000 °C.

As shown in Figure 3, obvious oxide spalling occurs in the sample oxidized at 1100 °C. This oxidative spalling results in a linear increase of oxidation weight gain. The falling oxide particles will cause solid particle erosion of the engine parts, making the engine efficiency decrease rapidly and even endanger the operation safety of the engine. The spalling of the oxide film is closely related to the volume ratio between the oxide and the alloy consumed to form the oxide (Pilling-Bedworth Ratio, PBR), and the value of PBR is an important criterion for judging the integrity of oxide film, and is expressed in [40] as

$$PBR=V_{\mathrm{O}}/V_{\mathrm{A}},$$

(7)

in which V_O denotes the volume of the oxide, and V_A denotes the volume of the alloy consumed to form the oxide. When the PBR of oxide is less than one, the volume of the oxide film is too small to cover the entire alloy surface, and tensile stress exists in the oxide film [34]. In this case, the oxide film is easy to crack, so the oxide film does not have a protective action. When the PBR of the oxide is greater than three, the oxide is subject to large compressive stress. Therefore, the PBR value of oxides with protective properties is usually in the range of one to two [33]. However, in addition to PBR, the bond strength of the interface should be combined to analyze whether the oxide film falls off easily. The strain (ε) caused by the volume difference between oxide and alloy can be expressed as [33],

$$\epsilon =\sqrt[3]{PBR}-1,$$

(8)

According to Hooke’s law, stress (σ_O) can be obtained by combining strain:

$${\sigma }_{\mathrm{O}}=E_{\mathrm{O}}\epsilon ,$$

(9)

in which E_O denotes the Young’s modulus of oxides. According to Figure 4 and Figure 8, XRD and EDS analysis confirmed that the oxide film of the alloy at 1100 °C is mainly composed of NiCr₂O₄, Cr₂O₃, and Al₂O₃. The PBR and Young’s modulus of NiCr₂O₄, Cr2O₃, and Al₂O₃ are listed in

Table 3. Material parameters of oxides.

Oxide	PBR	Reference	EO (Pa)	Reference	σO (GPa)
Al2O3	1.29	[40]	320	[41]	28.35
Cr2O3	1.95	[41]	260	[42]	64.83
NiCr2O4	2.05	[27]	233	[33]	62.99

. Put the values in Table 3 into Equations (8) and (9), and the growth stresses are obtained (see Table 3). It can be seen that the growth stress of the Al₂O₃ film is smaller than that of Cr₂O₃ and NiCr₂O₄. As a result, only de-cohesion occurs on the Al₂O₃ scale, while cracking and spalling occur on the Cr₂O₃ and NiCr₂O₄ scales. The spalling of the Cr₂O₃ and NiCr₂O₄ scale induces the linearly increasing oxidation weight gain at 1100 °C.

3.5. Comparison of Antioxidant Properties

Figure 12 illustrates the oxidation rate for 100 h at different temperatures. Figure 12 illustrates the oxidation rates for 100 h of FGH4095 (first-generation powder superalloy) [23], FGH4096 (second generation) [22] and FGH4108 at different temperatures. It can be seen that, compared with FGH4095 and FGH4096, FGH4108 has excellent oxidation resistance at no more than 1000 °C. According to HB 5258, “Test Method for Determination of Oxidation Resistance of Steels and Superalloys” (China Standard), FGH4108 belongs to the full oxidation resistance grade at 1000 °C or less. Due to the high Al content of FGH4108, a dense Al₂O₃ layer can be formed during the oxidation process (Figure 4a–c), which prevents the further diffusion of metal ions and oxygen and inhibits the occurrence of internal oxidation. In particular, the higher Hf content in FGH4108 further promotes the rapid formation of an α-Al₂O₃ oxide film and improves the oxidation stripping resistance of the alloy [43]. Therefore, it has a better antioxidant capacity at less than 1000 °C. When the temperature reaches 1100 °C, compared with FGH4095 and FGH4096, FGH4108 showed poor antioxidant activity but still belonged to the antioxidant class (Figure 12). The main reason is that the Al₂O₃ layer on the surface is no longer dense at high temperatures (Figure 9), which facilitates the diffusion of oxygen and leads to serious internal oxidation. Furthermore, since the internal oxidation mainly generates TiO₂, the higher Ti content of FGH4108 further promotes internal oxidation. The repeated spalling of the oxide film in the oxidation process ultimately leads to higher oxidation weight gain.

4. Conclusions

The oxidation kinetic curves of the alloy at 800 °C and 900 °C conform to the parabolic law. The oxidation kinetic curve at 1000 °C conforms to the cubic law. The oxidation kinetic curve at 1100 °C consists of the parabola and linear laws.

At 800 °C, the oxide film consists of two parts: the outer layer is a continuous dense protective Cr₂O₃ oxide film, and the inner layer is a discontinuous Al₂O₃ oxide layer. At 900–1100 °C, the oxide film consists of three layers: the outermost layer is a mixed layer of Cr₂O₃ and TiO₂ oxide film, the middle layer is a continuous layer of Al₂O₃ inner oxide, and the innermost layer is a dotted TiO₂.

The oxidation behavior of the FGH4108 superalloy is confirmed to be controlled by the diffusion of Cr, Al, Ti, and O through a selective oxidation process. The spalling of the Cr₂O₃ and NiCr₂O₄ scale induces the linearly increasing oxidation weight gain at 1100 °C. From the aspect of oxidation resistance, the long-term service temperature of the FGH4108 superalloy should not exceed 1000 °C.