Optimizing the High-Temperature Oxidation Resistance of Nb-Si-Based Alloys by Adding Different Ti/Mo/Hf Elements

Abstract

As a candidate material for turbine blades in aerospace engines, Nb-Si-based alloys have attracted significant research attention due to their high melting point and low density. However, their poor high-temperature oxidation resistance limits practical applications. Different alloying elements, including Ti, Mo, and Hf, were added to Nb-Si-based alloys to study the microstructural evolution of alloys. Additionally, the oxidation behavior and the oxidation kinetics of different alloys, as well as the morphology and microstructure of oxide scale and interior alloys at 1523 K from 1 h to 20 h were analyzed systematically. The current findings indicated that the Mo element is more conducive to promoting the formation of high-temperature precipitates of β-Nb₅Si₃ than the Ti and Hf elements. Inversely, the Ti element tends to cause the transition from high-temperature-phase β-Nb₅Si₃ to low-temperature-phase α-Nb₅Si₃, while the Hf element improves the appearance of the γ-Nb₅Si₃ phase but inhibits the other phases and refines the primary Nbss effectively. Noteworthily, compared with the oxidation weight gain of different alloys, Nb-16Si-20Ti-5Mo-3Hf-2Al-2Cr alloy has excellent high-temperature oxidation resistance, in which the oxidation products are TiNb₂O₇, Nb₂O₅, SiO₂, TiO₂, and HfO₂. It can be determined that in the oxidation process, the Ti element will preferentially form an oxide film of TiO₂, thereby wrapping around the matrix phases, protecting the matrix, and improving the antioxidant capacity, while the Hf element can form an infinite solid solution with the matrix and consume the small number of oxygen atoms entering the matrix, so as to achieve the effect of improving the oxidation resistance.

1. Introduction

As a turbine blade material, the following four requirements must be met [1]: high corrosion resistance, excellent overall performance (high creep resistance and strength, and a certain toughness), a low expansion coefficient, and good processing performance. As a candidate for new high-temperature structural materials, the melting point of silicon-based intermetallic compounds must exceed 2250 K [2,3]. Nb-Si intermetallic compounds are the most favorable competitor for turbine engine materials with comprehensive properties, especially Nb₅Si₃ intermetallic compounds [4], which have a melting point of 2757 K. The Nb-Si alloy has a lower density (7.16 g/cm³) and a higher melting point (about 1073 K) than the nickel-based superalloy intermetallic compound.

However, the high-temperature oxidation resistance of Nb-Si alloys is poor [5], which has caused a bottleneck restricting the development of Nb-Si alloys. However, its oxidation resistance can be effectively improved through alloying [1]. The common alloying elements of Nb-Si alloys are Ti [6], B [7], Hf [8], Al [9], Cr [10], Mo [11], Sn [12], Ta [13], etc. Among them, Ti and Hf can significantly reduce the diffusion rate of O in alloys. However, the content of Ti should not exceed 24% [14] to avoid further oxidation of the Nb-Si alloy. A high content of Cr can promote the formation of Laves (Cr₂Nb) [15], which significantly improves the oxidation resistance of Nb-Si. The metastable phase (Nb₃Si) will then be decomposed into Nbss and Nb₅Si₃ by the addition of Al [16], enhancing the stability of the phase (Nb₅Si₃) in the Nb-Si alloy. Al₂O₃ has a dense oxide film preventing O from permeating the Nb-Si alloy. Mechanical properties are taken into account while maintaining high-temperature oxidation resistance. Therefore, Cr and Al additions are generally kept at a low content [17,18,19]. Si can effectively enhance the oxidation resistance of the Nb-Si alloy by forming a dense SiO₂ layer, which inhibits oxygen diffusion. Even in Si-containing alloys, the optimized distribution and activity of Si further strengthen this protection [20,21,22]. It has been reported that Mo can reduce the pesting of Nb-Si alloys [23,24]. Sn elements can significantly reduce Nb-Si alloy’s pest-related oxidation sensitivity.

Research on the high-temperature oxidation resistance of multicomponent Nb-Si alloys is still in the preliminary stages, both at home and abroad. This article studies the oxidation behavior of Nb-Si-Ti-Mo-Hf-Al-Cr at 1523 K and analyzes the phase composition, microstructure, and composition of the oxide film after different oxidation times. A preliminary study on the high-temperature oxidation resistance of Nb-Si alloys has been conducted to study the mechanism of oxidation of Nb-Si-Ti-Mo-Hf-Al-Cr alloys.

2. Methods

2.1. Material Preparation

Six ingots with different chemical compositions (the chemical composition is expressed as atomic percentage, if not explicitly stated in this article), as shown in

Table 1. Nominal chemical composition (in at.%) of the experimental alloys, including Si, Ti, Mo, Hf, Al, Cr, and balance Nb.

No.	Si	Ti	Mo	Hf	Al	Cr	Nb
A1	16	18	10	5	2	2	Bal
A2	16	20	15	1	2	2	Bal
A3	16	22	5	3	2	2	Bal
A4	18	20	5	3	2	2	Bal
A5	16	20	5	5	2	2	Bal
A6	16	20	5	3	2	2	Bal

, were prepared by using raw materials of high purity Nb (>99.9 wt.%), Si (>99.99 wt.%), Ti (99.4 wt.%), Hf (>99.99 wt.%), Mo (>99.4 wt.%), Al (>99.9 wt.%) and Cr (>99.4 wt.%). Each ingot was arc-melted in a water-cooled copper crucible under argon atmosphere protection at 2473–2673 K, with a vacuum level lower than 5 × 10⁻³ Pa, and remelted 5 times for homogeneity. The weight and size of the samples were about 50 g and Φ20 mm × 30 mm; these were then cut into 6 × 6 × 3 mm³ dimensions for performing the microstructural analysis and oxidation tests. Unit area normalization was used to eliminate edge effects. The compositions used in this research were designed based on prior studies [6,8,17]. To balance the stability of the phase and the antioxidant properties, the ratio of Nb:Si deviated slightly from 3:1.

2.2. Measurement and Analysis Methods

The chemical compositions of the phases were determined by a combination of X-ray diffraction (XRD) (Empyrean, Panalytical, NL) and X-ray photoelectron spectroscopy (XPS, ESCALAB 250Xi, ThermoFisher, Waltham, MA, USA). The microstructure of the samples was analyzed using scanning electron microscopy (SEM, Quanta 200FEG, FEI, USA) with energy dispersive spectroscopy (EDS, Quanta 200FEG, FEI, USA). The primary dendrite size and the secondary dendrite size were measured and calculated using Image-Pro 11 software. The oxidation tests were performed in a high-temperature heat treatment furnace at 1523 K. The six surfaces of the specimen were polished until smooth with SiC papers and washed with alcohol. Then, each specimen was placed in a corundum crucible heated at 1523 K for 2 h to remove any moisture. Oxidized samples were extracted from the furnace at a particular time at 1523 K to measure the weight gain. Then, the mass difference before and after oxidation was calculated using a precision analytical balance with an accuracy of 0.0001 g to obtain the oxidation weight per unit area. The dendrite sizes of the samples were measured and counted using Image-Pro software.

3. Results and Discussion

3.1. Microstructural Evolution Induced by Different Chemical Compositions in Nb-Si-Ti-Mo-Hf-Al-Cr Alloys

The XRD pattern shows that there are two main phases in terms of Nbss and Nb₅Si₃ phases in the alloys numbered A1 to A6, as well as the three isomers (α, β, γ) of Nb₅Si₃ that appeared [25], as demonstrated in Figure 1. These are the α-phase (low-T tetragonal Cr₃B₃ type), β-phase (high-T tetragonal W₅Si₃ type), and γ-phase (hexagonal Mn₅Si₃ type, stabilized by impurities). The samples were not annealed, but the cooling rates after arc melting were controlled (~100 k/s), ensuring that the β-Nb₅Si₃ phase could be preserved. The oxidation resistance of Nb₅Si₃ varies with different structures; the high-temperature stability of the β-Nb₅Si₃ phase is stronger, which is due to its higher structural stability and generally higher Si content. High Si content is conducive to promoting the formation of a protective SiO₂ layer, thereby slowing oxygen diffusion and matrix oxidation. The second phase is γ-Nb₅Si₃, which can generate a denser oxide film at 1373 K; the main components are SiO₂ and TiO₂, which can effectively inhibit oxygen diffusion and reduce internal oxidation. The α-phase oxide film is porous and is mainly composed of Nb oxides (such as Nb₂O₅); it offers poor protection and is prone to severe internal oxidation. When alloying elements such as Ti, Hf, and Mo are added to Nb-Si-based alloys, the kinetics can be accelerated by lowering eutectoid reaction temperatures. In this case, only the stable Nb₅Si₃ phase can still occur in these alloys, which proves that alloying elements can effectively promote the eutectoid reaction of Nb₃Si→Nbss + Nb₅Si₃. This is a significant finding for eliminating metastable phases [26]. Additionally, by comparing the 6 groups of samples, it can be seen that more phase fractions of β-Nb₅Si₃ can form in alloys with the highest Mo content (A1 and A2 alloys), indicating that the Mo element is more conducive to the formation of the high-temperature precipitate β-Nb₅Si₃ than Ti and Hf elements in Nb-Si alloy [27]. Assuming that the Si content remains unchanged in the alloys, the increase of Ti element addition can reduce the phase fraction of β-Nb₅Si₃ but promote the content of α-Nb₅Si₃, implying that the Ti element can transform the high-temperature phase β-Nb₅Si₃ to the low-temperature phase α-Nb₅Si₃. Moreover, in a comparison of A4, A5, and A6 alloys, it can be seen that the increase in Hf content will promote the formation of γ-Nb₅Si₃ instead of other phases [28].

In order to further understand the microstructural evolution of the different alloys, SEM analysis was performed as shown in Figure 2, and the sizes of the phases and dendrites were measured and calculated using Image-Pro software. In addition, EDS tests were conducted to determine the composition of the precipitated phase in alloys. The results are shown in

Table 2. Energy spectrum analysis of micrographs in Figure 2; at.%.

Energy Spectrum Point	Nb	Si	Ti	Hf	Mo	Al	Cr
A	39.7	19.1	19.0	3.0	15.4	1.4	2.4
B	38.2	39.3	11.9	5.2	4.7	0.5	0.3
C	45.5	2.8	17.2	2.6	27.5	2.0	2.5
D	29.1	36.3	24.4	3.3	4.3	1.4	1.4
E	45.2	3.8	17.0	4.4	9.7	1.5	18.5
F	21.4	38.0	25.1	7.8	1.8	1.4	4.4
G	61.4	5.2	13.9	2.0	14.4	1.6	1.5
H	42.0	39.4	10.6	3.3	2.5	1.8	0.4
I	55.4	6.1	18.3	3.4	12.4	2.5	2.1
J	37.8	38.8	11.0	5.2	4.8	1.7	0.7
O	58.5	4.7	18.2	2.6	11.7	2.0	2.3
P	38.8	36.2	14.1	3.9	5.8	0.8	0.5

, where the data represent averages of 5 points per phase (±0.5 at.% error). It can be seen that the size of the Nbss phase in A1 alloy is relatively large; the mean primary dendrite length is 83.9 μm, and the mean of the secondary dendrite size is 23.5 μm. Meanwhile, the distribution of Nbss/Nb₅Si₃ eutectic in the alloys is relatively scattered [29] and has a fishbone shape, as shown in Figure 2a,b. The Nbss phase of the A2 alloy still has dendritic morphology and the dendrites are able to grow more fully, with 168.1 μm as the mean primary dendrite size and 32.4 μm as the mean secondary dendrite size, as shown in Figure 2c,d. In addition, part of the Nbss phase presents dispersed small particles (34.2 μm), forming a flower-like eutectic structure as shown in Figure 2d. However, the microstructure of the A3 alloy is different from the others (Figure 2e,f). Specifically, the dendrite phase of the Nbss disappears and transforms to a granular morphology phase with a size of about 21.6 μm in the A3 alloys; this is evenly distributed to form a flower-type eutectic structure with Nb₅Si₃. The interlamellar spacing of these eutectics ranges from 1.0 μm to 1.6 μm. The A4 samples are hypereutectic alloys, and the primary phases are mainly formed from bulk or strip Nb₅Si₃. Some eutectic structures present a network shape (as shown in Figure 2g,h), but others are chaotic eutectic structures. In contrast, the A5 samples are hypo eutectic alloys, and the Nbss has a granular morphology, while the lamellar eutectic composition is an extended structure distributed in the Nbss matrix [30], as shown in Figure 2i,j. With regard to the hypoeutectic alloy of A6 samples in Figure 2k,l, the white phases are the Nbss and the black are the Nb₅Si₃. It can be seen that the Nbss has a granular morphology, but most of the Nb₅Si₃ is in the form of long strips, and the eutectic material has a network structure.

The elemental surface distribution analysis of the A3 alloy was carried out by spectral analysis with an electron probe, as depicted in Figure 3. As shown in the first picture in Figure 3, the white color is Nb SS, and the dark gray color is Nb₅Si₃, and the grid form is in the eutectic phase. It is easy to see that the Mo, Al, and Cr elements are mainly dissolved in the Nbss, while the Ti and Hf elements are primarily dissolved in the Nb₅Si₃. The distribution of each alloying element is relatively uniform in each phase, but the Cr element is more prominent in the eutectic structure, and Mo has a distinct transition layer at the interface of Nbss/Nb₅Si₃. Although Ti and Nb are infinitely mutually soluble, Ti more easily reacts with Si to form a (Nb, Ti)₅Si₃ compound in the eutectic structure [31].

To study the specific distribution of alloying elements in the microstructure in detail, a surface scanning analysis of the A6 alloy was conducted, as shown in Figure 4. The alloy elements are uniformly distributed in the alloys, without significant segregation from the surface scan of the alloy. Mo and Cr elements are mainly dissolved in the Nbss. Ti, Hf, and Al elements are primarily distributed in the Nb₅Si₃. In this case, Al occupies the Si sites, forming Nb₅(Si, Al)₃ solid solutions [31]. The solid solubility of Ti and Mo in Nbss is higher than that of Hf because the radius of Mo atoms is close to that of Nb atoms [32], which causes the lattice distortion of Nbss to be smaller. At the same time, Hf exhibits the opposite relationship, which causes the significant lattice distortion of Nbss.

3.2. Oxidation Behaviour and Oxidation Kinetics of Alloys at 1523 K for Different Time Periods

The 6 × 6 × 3 mm³ samples were cut from the alloy ingots and subjected to a static oxidation test at 1523 K. The oxidation weight gain per unit area of the A1 to A6 alloys was obtained after 5 h of oxidation, as shown in Figure 5. It can be seen that the oxidation gain of the A6 alloy was the smallest, meaning that this alloy had the most significant high-temperature oxidation resistance.

In order to further explore the oxidation mechanism, the A6 alloy was oxidized for 1 to 10 h (time interval is 1 h), and the weight gain per unit area of the samples was calculated, as shown in Figure 6. The samples were oxidized simultaneously at 1523 K, with one specimen being quenched in water hourly. The furnace temperature fluctuation was controlled within ±5 °K. From 1 h to 10 h, the samples’ macroscopic morphology shows that the adhesion between the oxide layer and the substrate was good. There was no crack on the surface of the oxide layer after oxidation for 1 h. However, the adhesion became poorer with an increase in oxidation time. Tiny cracks appeared between the oxide layer and the substrate after oxidation for 5 h. The cracks became larger after oxidation for 10 h.

Table 3. Oxidation thickening and the resulting oxidation weight gain of A6 alloy after oxidation for 1 h to 10 h.

Oxidation Time	1 h	2 h	3 h	4 h	5 h	6 h	7 h	8 h	9 h	10 h
thickness (um)	23	77.9	116.5	129.5	144.5	188.5	243	323	429	438
Oxidation weight gain (mg/mm2)	0.21	0.26	0.42	0.46	0.75	0.95	1.00	1.16	1.19	1.38
Square of oxidation weight gain (mg2/mm4)	0.04	0.07	0.18	0.21	0.57	0.90	1.00	1.36	1.41	1.90

shows the average thickness and oxidation weight gain per unit area of A6 alloy at different oxidation times.

According to the Wagner theoretical model [33], the diffusion rate of the internal positive and negative ions through the oxide film determines the oxidation rate of the metal during high-temperature oxidation. The chemical position is the main driving force of ion diffusion. When the alloy undergoes high-temperature oxidation, its oxidation kinetic curve follows the laws of parabola. The model assumes that when t = 0 and Δm = 0, then:

$$\left(\Delta m\right)=k_{p}t$$

(1)

We differentiate Equation (1) as follows:

$${\displaystyle \frac{d\Delta m}{dt}}={\displaystyle \frac{k_p}{2\Delta m}}$$

(2)

where k_p is the alloy oxidation parabolic constant, Δm is the alloy oxidation weight gain, and t is time. Equation (2) is the oxidation weight gain model of the Wagner theoretical model. However, this model does not take that into account during the actual experiment. The high temperature of the oxidizing gas flow will cause uneven heat in the region, resulting in uneven temperatures in the oxidized alloy. The Wagner model assumes that the oxidation temperature at any time is constant, at a fixed value. Therefore, the oxidation kinetic curve, when calculated according to the model, shows a significant error compared to the actual situation. Based on this situation, this paper uses the method of Daniel Monceau et al. [34] to correct the oxidation calculation of the alloy when using the Wagner model. When t = t_i and Δm = Δm_i, Equations (1) and (2) become Equations (3) and (4), as follows:

$$\left(\Delta m\right)-\left(\Delta m_i\right)=k_p\left(t-t_i\right)$$

(3)

$${\displaystyle \frac{d\Delta m}{dt}}={\displaystyle \frac{k_p}{2\Delta m_i}}$$

(4)

The oxidation weight gain of the A6 alloy is taken as the ordinate. The oxidation time is the abscissa used to draw the oxidation kinetic curve of the A6 alloy, as shown in Figure 7a. Furthermore, the relationship between the square of oxidation weight gain (Δm)² and the oxidation time t is obtained [35]. As shown in Figure 7b, it can be seen that (Δm)² is substantially linear, with an oxidation time of t. Thus, the oxidation curve of the A6 alloy follows a parabolic law [36].

It can be seen from Equation (4) that the k_p value of the alloy can be determined by calculating the oxidation weight gain at different times. The optimized alloy’s 1523 K oxidation parabolic constant was calculated to be two-stage parabolic kinetics, such as the transient state (0–3 h, kₚ = 0.12 mg²/(mm⁴·h)) and the steady state (3–10 h, kₚ = 0.24 mg²/(mm⁴·h)). The value of k_p of the A6 alloys aligned with the literature values of kₚ = 0.18–0.30 mg²/(mm⁴·h) for Nb-Si-B alloys [33], which confirmed that the A6 alloy had superior oxidation resistance.

3.3. Morphologies and Microstructure of Oxide Scale and Interior Alloys

Figure 8 shows the morphology of the oxide film of the A5 alloy and A6 alloy after 5 h. It can be seen from Figure 8 that a few cracks grew by a rod-shaped oxide, with a growth direction perpendicular to the oxidized surface. The oxides were in close contact with each other in the A6 alloy. This structure prevented O from quickly entering the interior alloy during oxidation. Thus, the structure effectively prevented the alloy from oxidation, which is one of the reasons why the oxidation weight of A6 gain was most significant after oxidation for 5 h. The oxide film of A5 alloy grows in a sheet form, and the contact between the oxides is not tight. This sheet structure makes it easy for it to form microcracks. Then, O can diffuse into the interior alloy through these microcracks [37], producing severe oxidation.

After the alloy had oxidized, the oxide film that formed on the surface was peeled off and then ground into a fine powder. X-ray diffraction analysis was used to analyze the phase composition of the oxide film. The X-ray diffraction pattern of the A6 alloy oxidation from 5 h to 20 h is shown in Figure 9.

The oxidation products of the A6 alloy after oxidation were mainly TiNb₂O₇, Nb₂O₅, SiO₂, TiO₂, and HfO₂. The phase corresponding to the most substantial peak in the XRD pattern was TiNb₂O₇. However, the volume fraction of Nb₂O₅ and TiO₂ was not large, indicating that the Nb and Ti elements react with O simultaneously during oxidation. Because the activity of Ti is higher than the Nb, during the oxidation process, the growth rate of TiO₂ exceeds the growth rate of Nb₂O₅, and the faster-growing TiO₂ will gradually wrap around the slower-growing Nb₂O₅. The XRD results in Figure 9 show that the peak of TiO₂ in the oxide film was 458.5eV (Ti 2p₃/₂), confirming the formation of the TiO₂ passivation layer, which is superior to Nb₂O₅ formation and improves oxidation resistance by wrapping around the matrix phase. Then, a solid phase reaction occurs [38]: Nb₂O₅ + TiO₂→TiNb₂O₇. With the prolongation of oxidation time, the content of TiNb₂O₇ increases, and the content of Nb₂O₅ and TiO₂ also increases. This also proves that Nb and Ti simultaneously react with O at the beginning of oxidation. The content of HfO₂ also increases with increasing oxidation time, because the Hf is sufficiently diffused to combine with O as the oxidation time lengthens [39].

When the A6 alloy is oxidized, its structure (from the outside to the inside) is the oxidized outer layer (oxide film) to the transition layer of the interior metal, through which the phase composition of the oxidized outer layer is analyzed by X-ray diffraction. The transition layer and the interior metal microstructure are shown in Figure 10. It can be seen from Figure 10 that as the oxidation time increases, the thickness of the transition layer gradually decreases until it disappears after oxidation for 20 h. This is because when the oxidation time is short, the alloying elements cannot diffuse to the outside [40]. Most of the O is consumed when O passes through the oxide film. The remaining small part of O is consumed by the active element Hf in the alloy, so the interior alloy remains unoxidized, as shown in Figure 10b. When the oxidation time is extended, the alloying elements of the internal alloy have this longer diffusion time to combine with an increased amount of O entering the internal alloy, which causes severe oxidation.

Through the internal alloy figures given for the alloy after different oxidation times, it is known from [41] that a small amount of white granular oxide begins to appear in Nbss. Most granular oxide appears at the interface of Nbss/Nb₅Si₃ in the initial oxidation stage. As the oxidation time is prolonged, the content of white oxide increases, showing two shapes, namely, blocks and strips. The bulk oxide is in Nbss, and the long strip oxide appears at the interface of the Nbss/Nb₅Si₃ phase, while the black needle-like oxide appears in Nbss, which is messy in its distribution.

To determine the composition of each phase, EDS analysis was performed on the internal alloy after oxidation for 20 h, as shown in Figure 10d and

Table 4. Energy spectrum analysis of the oxidation of A6 alloy for 20 h in Figure 10d; at.%.

Spectrum Point	O	Nb	Si	Ti	Hf	Mo	Al	Cr
A	2.8	38.7	37.1	11.8	3.4	4.2	1.5	0.5
B	11.0	59.2	4.7	7.3	1.0	12.0	1.7	3.3
C	14.2	29.9	23.3	22.0	2.6	4.4	2.6	1.1
D	19.2	21.3	27.5	12.5	15.5	1.9	1.5	0.5

. The gray smooth tissue (Point A) in the tissue is the Nb₅Si₃ phase. It can be seen that the Nb₅Si₃ phase has a high antioxidant capacity, while the Nbss (Point B) is easily attacked by O. The alloying elements that are dissolved in the Nbss will react with O to form a variety of oxides. The content of Ti in the black acicular oxide (Point C) is significantly higher than that in Nbss, indicating that it is TiO₂. The white bulk oxide (Point D) is HfO₂.

To study the oxidation mechanism of A6 alloy in detail, it is necessary to know the type and electronegativity of the oxides of each alloying element. The physical parameters of each alloying element are shown in

Table 5. Physical parameters of the alloying elements.

Element	Nb	Si	Ti	Hf	Mo	Al	Cr
Highest valence oxide	Nb2O5	SiO2	TiO2	HfO2	MoO3	Al2O3	Cr2O3
Electronegativity	1.41	1.91	1.38	1.16	1.47	1.61	1.65

. It can be seen from Table 4 that the electronegativity of elements in the A6 alloy (Nb-16Si-20Ti-5Mo-3Hf-2Al-2Cr) has the following order from small to large: Hf, Ti, Nb, Mo, Al, Cr, and Si [42]. It is known from its electronegativity that Hf is the most active and has a significant affinity with O. Therefore, when O enters the internal alloy, Hf first combines with O to form HfO₂. Second, Ti reacts with O. At this point, the A6 alloy experiences selective oxidation. By comparing the content of Hf in Nbss before and after oxidation, it can be seen that the content of Hf after oxidation is significantly reduced because Hf combines with O to form granular HfO₂ under low O partial pressure. The formation of bulk HfO₂ mainly occurs because metal bonds bond the metal atoms in Nbss, meaning that the diffusion of Hf in Nbss is relatively easy [41]. When O is present at the interface of Nbss/Nb₅Si₃, O preferentially reacts with Hf, resulting in a concentration gradient of Hf in the local microstructure. Then, the Hf in Nbss diffuses from the interior of Nbss along the concentration gradient to the interface, forming oxides of Hf. The Hf in the Nb₅Si₃ is more robust, due to the formation of the Si-Hf-Si covalent bond [43], meaning that the Hf in Nb₅Si₃ is difficult to diffuse to the Nbss/Nb₅Si₃ interface. The electronegativity of the Ti element is lower than Hf, meaning that TiO₂ quickly appears at the interface of the two phases.

Figure 10c,f shows the linear scanning of the A6 alloy at 1523 K for 7 h and 20 h along the vertical oxide film/internal interface, indicating the gradients of O, Ti, and Hf. Working from the dotted line region (transition layer) in Figure 10b, the content of O in the oxide film is the highest, and the O content in the transition layer is low. However, the O content in the internal alloy is the lowest, and it is maintained at the same level. This shows that when O passes through the transition layer, the O is mainly consumed, and the amount of O entering the matrix is extremely small. Ti, Hf, and Al have multiple peaks in the transition layer, and the peak value corresponds to that of O. This indicates that Ti, Hf, and Al preferentially react with O to form oxides in the transition layer, consuming the O molecules from the outside. Then, various oxides also prevent O from continuing to penetrate the internal alloy, which is why the O content of the internal alloy is very low [44]. The Hf in the internal alloy also shows a peak, indicating that a small amount of O entering the internal alloy is first consumed by Hf, resulting in no severe oxidation phenomenon.

After oxidation for 20 h, the transition layer has disappeared. High Ti, Hf, Si, and Al peaks are seen near the interface between the oxide film and the internal alloy. This indicates that Ti, Hf, Si, and Al migrate to the vicinity of the interface and undergo oxidation. The content of Mo is still significantly lower than that of the oxide film, but the content of Mo in the internal alloy is increased. This is similar to the change in Nb, which occurs because Mo and Nb are infinitely miscible with each other [45]. When O enters the internal alloy, Mo does not preferentially react with O. It is protected by active elements when the alloy is oxidized. Conversely, Mo oxide (MoO₃) [46,47] evaporates easily at high temperatures, which is also one of the reasons why the content of Mo in the internal alloy is higher than that in the transition layer.

4. Conclusions

(1)

We mainly identified Nbss and Nb₅Si₃ and a small amount of metastable phase (Nb₃Si) in the Nb-Si-Ti-Mo-Hf-Al-Cr alloy. These alloying elements (Ti, Hf, Mo) can eliminate the metastable phase. The Mo element has the function of stabilizing β-Nb₅Si₃, while the presence of Hf can refine the primary Nbss in the alloy, and the Ti element can promote the formation of α-Nb₅Si₃.

(2)

Nb-Si-Ti-Mo-Hf-Al-Cr alloy oxidizes at 1523 K at different times. The oxidation products are TiNb₂O₇, Nb₂O₅, SiO₂, TiO₂, and HfO₂, respectively. During the oxidation process, a solid phase reaction occurs: Nb₂O₅ + TiO₂→TiNb₂O₇. The formation of TiNb₂O₇ can effectively inhibit oxygen diffusion, while Hf consumes oxygen via a solid solution.

(3)

The oxidation resistance of Nb-Si-Ti-Mo-Hf-Al-Cr alloys is primarily affected by solid solution elements. When the alloy is oxidized, the Nb₅Si₃ shows extreme oxidation resistance, but O cracks quickly corrode the Nbss and are generated due to the deformation stress between the Nbss and Nb₅Si₃, which exacerbates the oxidation of the internal alloy. Hf can rapidly migrate to the Nbss/Nb₅Si₃ interface to form HfO₂, which consumes O, preventing O from diffusing into the internal alloy and effectively improving the oxidation resistance of the alloy. Optimizing the Ti/Hf/Mo ratios synergistically can effectively improve the oxidation resistance and mechanical properties of alloys.