Effects of Gd₂O₃ Content on the Infrared Emissivity and Ablation Resistance of HfB₂/SiC/TaSi₂ Coating at 4400 kW/m²

Abstract

To improve the infrared emissivity and the ablation resistance of HfB₂/SiC/TaSi₂ coatings for serving in heat flux of 4400 kW/m², HfB₂/SiC/TaSi₂ coatings with different contents of high-emissivity Gd₂O₃ were prepared using atmospheric plasma spraying. The highest emissivity in 3–5 μm can reach up to 0.92 at 1273 K when the Gd₂O₃ content is at 10 vol.%. The increase in the emissivity is attributed to the additional electronic transitions induced by oxygen vacancies, which are generated by substituting Hf⁴⁺ with Gd³⁺. Due to the high emissivity, the surface temperature of the coating modified with 10 vol.% Gd₂O₃ was decreased by ~100 K. Meanwhile, the mass and the liner ablation rate are confirmed to be 4.28 × 10⁻⁷ kg/s and 2.15 × 10⁻⁷ m/s, respectively, which are decreased by 80% and 31% compared to the undoped HfB₂/SiC/TaSi₂ coating. During ablation, HfB₂/SiC/TaSi₂/Gd₂O₃ coating was oxidized to HfO₂, Gd₂Ta₂O₇, HfSiO₄, and GdTaO₄. A stable Hf–Ta–Gd–Si–O multiphase glass was formed on the surface of the coating, which could restrain oxygen penetration. However, the excessive amount of Gd₂O₃ is detrimental to the ablation performance due to its consumption of the SiO₂ glass layer. These findings indicate that the addition of an appropriate amount of Gd₂O₃ could improve the anti-ablation performance of the modified coating.

1. Introduction

Carbon/carbon (C/C) composites have been used as structural and thermal components owing to their excellent mechanical properties and thermal resistance. However, the poor oxidation resistance of C/C limits its application in hypersonic vehicles [1,2]. The preparation of ultra-high-temperature ceramics (UHTCs) on the surface of C/C composites can block direct contact between C/C and oxygen at higher temperatures [3].

HfB₂ is a representative UHTC with a low thermal expansion coefficient, high thermal conductivity, low catalytic coefficient, and high-temperature phase stability [4,5]. It has been reported that HfB₂ has good oxidation resistance below 1473 K by forming liquid B₂O₃, which can hinder oxygen penetration and seal cracks and holes. However, when the temperature exceeds 1473 K, its protective ability is destroyed owing to the rapid volatilization of B₂O₃ [6]. It has been reported that adding 20–30 vol.% SiC to HfB₂ can result in excellent antioxidant ablation performance below 1873 K [7]. Nevertheless, the formation of a SiC-depleted region remains a problem for the HfB₂–SiC system. The additional introduction of TaSi₂ improves the density and oxidation resistance of UHTCs as an additive phase because Ta₂O₅ reduces the volatilization rate of the glass SiO₂ phase. Opila et al. [8] conducted a detailed analysis of the effect of TaSi₂ addition to borides and showed that the addition of Ta can enhance the glass-phase viscosity and reduce HfO₂ oxygen vacancies. However, it was also pointed out that the addition of TaSi₂ did not enhance the antioxidant capacity of the coating at 2073 K and higher. Zhang et al. [9] have prepared HfB₂–SiC–TaSi₂ coating by a spark plasma sintering process and studied its oxidation behavior below 2073 K; the results showed that the generated Hf–Ta–B–Si–O composite glass layer could effectively reduce the coating-like diffusion rate and enhance the oxygen barrier ability of the coating.

However, TaSi₂ did not improve oxidation protection capability above 2073 K because of Ta₂O₅ volatilization and melting of the metal oxides. This also renders the HfB₂/SiC/TaSi₂ system unable to meet the demands for high-temperature use. Therefore, it is urgent to seek a more stable additive phase for the HfB₂-based composite system in ultra-high-temperature environments (≥2073 K).

In addition to improving the high-temperature properties of UHTCs, adding a radiation component to UHTCs is an effective way to indirectly improve high-temperature properties through cooling of the surface temperature during hypersonic flight via reradiation of the absorbed heat [10]. Preliminary work has indicated that rare-earth iron or rare-earth oxide doping can improve the surface infrared emissivity of coatings, which, in turn, enhances the radiation exchange capacity; this coupling can promote the antioxidant properties of the coatings [11,12,13,14,15,16,17,18]. Tan et al. [19] have studied the oxidation protection of a rare-earth-modified ZrB₂–SiC coating using an oxyacetylene torch and found that rare-earth-modified coatings (Sm₂O₃ or Tm₂O₃ modified ZrB₂–SiC) displayed high thermal stability and offered additional oxidation protection. Liu et al. [13] have prepared REO–HfO₂ (REO = rare earth oxide = Tb₄O₇, Gd₂O₃, or Sm₂O₃) coatings and pure HfO₂ coatings using atmospheric plasma spraying (APS) and studied the infrared radiation performance of these coatings. The results showed that the REO–HfO₂ coatings displayed higher infrared emissivity than pure HfO₂ coatings and exhibited excellent thermal resistance at 1873 K without undergoing a phase change and exfoliation. Tan et al. [14] have studied the hemispherical emissivity and ablation resistance of a Sm-doped ZrB₂/SiC coating and found that 5 mol% Sm-doped ZrB₂/SiC coating had the highest emissivity with dense oxide scale formation. Previous studies have shown that doped REOs could exhibit better oxygen-blocking capability and form a dense oxide scale upon oxidation [20,21,22,23,24]. Feng et al. [20] and Li et al. [21] have prepared La₂O₃-modified HfC–SiC coatings and ZrB₂ coatings by supersonic atmospheric plasma spraying (SAPS), and the results showed that the incorporation of La₂O₃ could improve the ablation resistance. Wang et al. [22,23] have prepared Y₂O₃-modified MoSi₂ coating by SAPS, and the results showed that adding Y₂O₃ at an appropriate concentration could increase the viscosity of SiO₂ glass. Qian et al. [24] have studied the oxygen-blocking capability of La₂O₃-modified HfB₂–SiC coating and found that 10 vol.% of La₂O₃ additive could improve the oxygen barrier properties.

Gd³⁺, as a lanthanide rare earth element in the middle position, has an oxide Gd₂O₃ with a melting point of ~2603 K, which is higher than that of Ta₂O₅ (2058 K). Meanwhile, Gd₂O₃ can react with the coating oxidation product HfO₂ to form Gd₂Hf₂O₇, thus influencing the structural stability of the oxide layer. Gd₂Hf₂O₇ has a cubic pyrochlore structure (Fd3m) with a maximum pyrochlore–fluorite transition temperature of 2600–2800 K, and no Gd₂Hf₂O₇ structural changes are observed below this temperature [17]. Due to the high melting point of Gd₂O₃ and high infrared emissivity and its ability to influence thermal stability, Gd₂O₃ can be added to HfB₂/SiC/TaSi₂ coatings to improve the high-temperature oxidation resistance properties (>2073 K) of the coatings. However, there are few studies on Gd₂O₃-modified coatings, and the effect of Gd₂O₃ at different contents on the ablation resistance of HfB₂/SiC/TaSi₂ has not been reported.

In this study, Gd₂O₃-modified HfB₂/SiC/TaSi₂ coatings were prepared by APS. The microstructure and component distribution of the coatings were examined to study the effect of varying contents of Gd₂O₃ on the characteristics of the coatings. In addition, the infrared radiative performance and ablation behavior of the coatings were systematically investigated to evaluate the effect of emissivity on the anti-ablation resistance properties of the Gd₂O₃-modified HfB₂/SiC/TaSi₂ coatings. Our findings will provide candidate materials for UHTCS, aiming at its potential ultra-high-temperature application in the aerospace field.

2. Experimental Procedure

2.1. Preparation

HfB₂ particles (size range 1–3 μm, purity ≥ 99.9%, Beijing Zhongjinyan New Material Technology Co., Ltd., Beijing, China), SiC particles (size range 1–3 μm, purity > 99.9%, Beijing Zhongjinyan New Material Technology Co., Ltd.), TaSi₂ particles (size range 1–3 μm, purity ≥ 99.9%, Beijing Zhongjinyan New Material Technology Co., Ltd.) and Gd₂O₃ particles (500 μm, purity ≥ 99.9%, Beijing HuaweiRuike Chemical Co., Ltd., Beijing, China) were used as raw materials for the preparation of HfB₂, SiC, TaSi₂, and Gd₂O₃ agglomerate powders, respectively. Five HfB₂/SiC/TaSi₂/Gd₂O₃ mixtures with different ratios were designed, as listed in

Table 1. As-prepared coating composition.

Sample	HfB2 (vol.%)	SiC (vol.%)	TaSi2 (vol.%)	Gd2O3 (vol.%)
Gd0	77.8	11.1	11.1	0
Gd5	73.8	10.6	10.6	5
Gd10	70	10	10	10
Gd15	66.2	9.4	9.4	15
Gd20	62.4	8.8	8.8	20

. Among the five mixtures, the volume ratio of TaSi₂ and SiC remained unchanged, and the volume ratio of Hf to Si in each group was fixed at 7:3, whereas the addition of Gd₂O₃ was varied at 0, 5, 10, 15, and 20 vol.%.

The raw particles were ball milled (JM-15L, Changsha Tianchuang Powder Co., Ltd., Changsha, China) with deionized water (Beijing Wanxin Chemical Industry Trade Center, Beijing, China) and PVA (0.4 wt%, Forsman Technology (Beijing) Co., Ltd., Beijing, China) for 5 h to produce a uniform slurry. The slurry was then immediately injected into a spray drying apparatus (Wuxi Dongsheng LGZ-8 spray drier, Wuxi, China) to agglomerate the powders. Before spraying, the agglomerated powders were treated using induction plasma spheroidization (IPS) to improve their density and sphericity. The IPS-treated particles with a particle size of 10–50 µm were sieved and used as feedstock powders for APS. The IPS treatment of the powders was carried out using a PL-35 model system (TEKNA Plasma System Inc., Sherbrooke, Canada). The other parameters of the IPS treatment of the HfB₂/SiC/TaSi₂/Gd₂O₃ agglomerate powders are listed in

Table 2. Parameters of IPS procedure.

Parameter	Set Value
Ar flow rate (m3/s)	4.7 × 10−4
H2 flow rate (m3/s)	4.7 × 10−5
Processing chamber pressure (kPa)	1.0 × 102
Power (kW)	30

.

The coatings were prepared with US-made Praxair GTS-5500 equipment using APS, and the plasma torch used was SG 100. Powders were plasma sprayed onto C–C substrates (ø 0.025 m × 0.010 m) with an average coating thickness of roughly 200–250 µm. The parameters used for the deposition of the coatings are listed in

Table 3. Parameters of the APS process.

Parameter	Set Value
Current (A)	9.0 × 102
Main gas: Ar (m3/s)	1.5 × 10−3
Auxiliary gas: He (m3/s)	8.3 × 10−4
Carrier gas: Ar (m3/s)	2.0 × 10−4
Spraying distance (m)	7.5 × 10−1

.

For convenience, a naming convention based on the actual Gd₂O₃ concentration was used. For example, the Gd5 coating contains 5 vol.% Gd₂O₃ in the HfB₂/SiC/TaSi₂ coating.

2.2. Ablation Testing

The oxyacetylene flame (OAF) was selected to evaluate the anti-ablation property of the HfB₂/SiC/TaSi₂/Gd₂O₃ coatings. To achieve a long-term ablation of 180 s, the Mini-spray jet (Gastechnik GmbH & Co. KG, Neustadt/Wied, Germany) equipped with cooling air was used as the oxyacetylene spray gun in the OAF equipment. The local heat flux of the flame was measured using a Gordon gauge (AnYi Instrumentation Ltd., Shanghai, China), and the surface temperature of the specimens was obtained using an MR1S dual colorimetric infrared thermometer (Raytek Marathon, Santa Cruz, CA, USA).

During ablation, the distance between the gun and the specimen was 0.045 m. The pressure and flow rate of oxygen and acetylene were 0.5 MPa, 2.5 × 10⁻⁴ m³/s, and 0.05 MPa, 1.7 × 10⁻⁴ m³/s, respectively. In this case, the peak heat flux of the flame was 4400 kW/m². The ablation property of the coatings was characterized using the mass and linear ablation rates, which were calculated by the following formulae:

$$R_m=\frac{\Delta m}{\Delta t}$$

(1)

$$R_l=\frac{\Delta l}{\Delta t}$$

(2)

where $\Delta m$ and $\Delta l$ are the change in the mass and thickness of the sample, respectively, and $\Delta t$ is the ablation time. The ablation rates of the samples were based on the average values calculated for five formulation samples.

2.3. Characterization

The composition of HfB₂/SiC/TaSi₂/Gd₂O₃ coatings before and after oxidation testing was analyzed by X-ray diffraction (XRD, X’ Pert PRO), with Cu Kα radiation, a step-size of 0.04°/s, and a 1 s counting rate. The morphology of the coatings was analyzed by scanning electron microscopy (SEM, JSM-6460), and energy dispersive spectroscopy (EDS) was used for elemental analysis. The porosity of the sprayed coatings before ablation was estimated by image analysis using Image J software and the average from 10 pictures was taken for each sample. An infrared spectrometer (IR-2, Shanghai Chengbo Optoelectronic Technology Co., Ltd., Shanghai, China) was employed to measure the infrared emissivity of the coatings at room temperature, whereas the infrared emissivity at 1273 K was tested using a Nicolet 6700 infrared radiation instrument at the China National Supervision and Test Center for Infrared and Industry Galvanothermy Product Quality (Wuhan, China).

3. Results and Discussion

3.1. Phase Composition and Morphology of as-Sprayed Coatings

Figure 1 shows the XRD patterns and probable phases of the as-prepared coatings before ablation. As shown in Figure 1, the phase composition of the as-prepared coatings was mainly HfB₂ (PDF#75-1049, hexagonal, and spatial group structure is P6/mmm (191)), and diffraction of SiC and TaSi₂ was not observed. The reasons for the disappearance of the SiC and TaSi₂ diffraction peaks may be as follows: (i) HfB₂ and SiC form a eutectic phase, HfSi₂ (PDF#72-1201, orthorhombic, and spatial group structure is Cmcm (63)), during the process [25]. (ii) The amount of TaSi₂ is relatively small. (iii) There is a low penetration depth of X-rays in SiC. In addition, the Gd₂O₃ (PDF#86-2477, cubic, and spatial group structure is la-3(206)) phase was found in the Gd5, Gd10, Gd15, and Gd20 coatings, and the intensity of the Gd₂O₃ diffraction peak gradually increased with increasing Gd₂O₃ content. It can be seen in Figure 1 that the Gd₂O₃ not only reacted with other phases, but also did not melt with the other phases. The HfO₂ (PDF#21-0904, orthorhombic, and spatial group structure is Pbcm (57)) diffraction peak was observed in all five coatings, indicating that the powder at the edge of the jet during the spraying process was oxidized by O₂ in air owing to the absence of a protective atmosphere. In addition, the agglomerated powders in the center of the jet inevitably react with oxygen when sprayed onto the C–C substrate. HfO₂ was also produced while the powder on the substrate cooled down.

Figure 2 shows the surface morphology of the as-sprayed coatings. All five coatings show good powder-spreading characteristics with few surface undulations and flat fusion condensation spreading layers. This planar laminar feature can increase the contact area with the deposited powder, reduce the interlayer defects caused by differences in shape orientation, and facilitate a dense coating. In particular, the spreading performance of the coating improved with increasing Gd₂O₃ content. This indicates that the introduction of Gd₂O₃ with a lower melting point can improve the deposition effect of the powder and increase the density of the coating.

Figure 3a–e show the cross-sectional morphology of the as-sprayed coatings. It can be seen that the thickness of all the coatings was 200–300 μm, and these coatings were perfectly integrated with the C–C substrate. It can be seen that the Gd0 coating has a large coarse area. With an increase in Gd₂O₃ addition, the loose area decreases and the cross-sectional spalling crater characteristics gradually weaken.

Table 4. Porosity of as-prepared coatings.

Sample	Porosity (%)
Gd0	19.91
Gd5	18.55
Gd10	16.73
Gd15	12.32
Gd20	11.28

lists the porosity of the five coatings. It can be seen that the coating porosity decreases from 19.91% to 11.28%, indicating that the introduction of Gd₂O₃ improves the powder deposition efficiency and helps to enhance the coating density. With increased density, resistance to oxidation improves. Figure 3f shows the elemental distribution of the Gd10 coating, and it can be seen that the Hf, Ta, Si, and Gd in the coating are uniformly distributed, indicating that there was no elemental bias inside the coating.

3.2. Infrared Radiation Property of as-Prepared Coatings

Figure 4 shows the infrared emissivity of the Gd₂O₃-modified coatings at room and high temperature in the 3–5 µm band. At room temperature, the emissivity increased from ε = 0.69 for the Gd0 coating to ε = 0.72, 0.73, 0.74, and 0.71 for the Gd5, Gd10, Gd15, and Gd20 coatings, respectively. At temperature T = 1273 K, the emissivity increased from ε = 0.85 for Gd0 coating to ε = 0.88, 0.92, 0.91, and 0.89 for the Gd5, Gd10, Gd15, and Gd20 coatings, respectively. The emissivity improved with an increase in Gd₂O₃ content, and the highest emissivity was reached for the Gd15 coating at room temperature and the Gd10 coating at high temperatures. Moreover, the high-temperature emissivity was much higher than that at room temperature for all five coatings. This is attributed to the increased lattice vibration and electron–phonon coupling at high temperatures, which improves infrared lattice absorption and multi-phonon absorption.

The above variation in emissivity can be attributed to two reasons: Gd₂O₃ has a high emissivity [26] and Gd³⁺ ions with larger radii and lower atomic weights enter the HfO₂ lattice, occupying the Hf⁴⁺ ion positions. Some oxygen ions in the lattices spill over to form oxygen vacancies, which then attract and trap electrons to form color centers [27]:

$$\begin{gathered} {Gd}_2{O}_3\stackrel{{HfO}_2}{\to }2{Gd}_{Hf}^{\prime }+3O_O+V_O^{..} \\ {O}_O\to V_O^{..}+2e^{-}+\raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$2$}\right.O_2 \end{gathered}$$

where ${Gd}_{Hf}^{\prime }$ is a Gd³⁺ ion occupying the position of Hf, $O_O$ represents O²⁻ ion in its normal position, $V_O^{..}$ is an oxygen vacancy, and $e^{-}$ is an electron.

The oxygen vacancies and color centers are the same as those in this discussion. Oxygen vacancies can cause localization of the electronic state in the material bandgap, promoting additional electronic transitions. According to [28], neutral oxygen vacancies significantly increase the emissivity of ZrO₂, owing to a mid-gap state caused by oxygen vacancies in ZrO₂. In this study, the oxygen vacancies produced by the substitution of Hf⁴⁺ with Gd³⁺ have energy levels inside the HfO₂ band gap. Thus, both the Gd³⁺ emission band and the oxygen vacancies produce localized energy levels inside the HfO₂ band gap, such that exciting electrons into these localized energy levels promotes increased infrared emissivity. In addition, Gd₂O₃ could react with oxidation products (e.g., HfO₂) to form a new phase, Gd₂Hf₂O₇. Because the cation radii ratio (r_Gd³⁺/r_Hf⁴⁺) is lower than 1.46 and there is an excess amount of anions (O²⁻), Gd₂Hf₂O₇ has a defective fluorite-type structure [13,29], which is expected to improve material emissivity.

When compared with the Gd0 coating, the maximum difference of the high-temperature emissivity of Gd₂O₃-modified coating in the 3–5 µm band is 0.07, which is an improvement of 8.2%, indicating that the introduction of Gd₂O₃ has better infrared radiation performance, and Gd₂O₃ mainly acts in the high-temperature oxidation phase of the coating.

Therefore, the improvement in infrared emissivity with increasing Gd₂O₃ content from 0 to 10 vol.% can be attributed to the production of Gd-doped HfO₂ and Gd₂Hf₂O₇ phase because they both have high emissivity. When doped with 10 vol.% Gd₂O₃, the Gd-doped HfO₂ reaches a saturated doping state. The decrease in infrared emissivity with an increase in Gd₂O₃ content from 10 to 20 vol.% is due to the reduction of Gd-saturated HfO₂, which has a higher emissivity than Gd₂Hf₂O₇.

3.3. Ablation Resistance of Coatings

3.3.1. Macro-Oxidation Behavior of the Coatings

Figure 5 shows the local heat flux distribution and surface temperature curves of the coatings during 180 s of ablation. As shown in Figure 5a, the peak heat flux was 4400 kW/m² within a radial distance of 0.002 m, and the heat flux values decreased radially outwards. Illustrations (i) and (ii) in Figure 5b show flame photographs during the free jet and ablation tests, respectively. Because the ablation gun in this study had a cooling device, the free jet flame length was 91 mm, which is shorter than the flame length without a cooling device [30].

Figure 5b presents the surface temperature curves of the five coatings during heating. The surface temperature increased rapidly for the Gd0 and Gd20 coatings and, subsequently, it reached a steady state value equal to ~2173 K for the Gd0 coating and ~2073 K for the Gd20 coating. For the Gd5, Gd10, and Gd15 coatings, the surface temperature increased continuously during the ablation process, indicating that their oxidation products change continuously in this ablation environment. During ablation, identical heating and thermal convections occur in the ablation center; thus, the surface temperature is dominated by heat radiation and conduction of the coating materials. Compared with Gd0 coating, the Gd₂O₃-modified coatings exhibited a surface temperature ~100 K lower; this is consistent with Gd₂O₃-doped coatings having a lower emissivity than Gd0 coating.

Macrographs of all the coatings before and after ablation are shown in Figure 6. The ablation centers of the coatings that suffered the most heat flux were chosen to evaluate their ablation resistance and detailed microstructure. Before ablation, the surfaces of all the coatings appeared dark grey; therefore, Figure 6a was used to represent all the coatings. In Figure 6b–f, the ablation edges of the five coatings changed to white. The ablation centers of Gd0, Gd5, and Gd10 were white, whereas those of the other two coatings were yellow, which is related to the Gd₂O₃ content. Additionally, edge warping appeared in the Gd0 and Gd5 coatings, and spalling areas and bubbles appeared in the Gd20 coating. It can be speculated that the Gd10 and Gd15 coatings exhibit better ablation resistance.

Figure 7 shows the mass and linear ablation rates under the same ablation conditions. The Gd0 coating exhibits the highest mass and linear ablation rates of 2.14 × 10⁻⁶ kg/s and 3.13 × 10⁻⁷ m/s, respectively. With an increase in the Gd₂O₃ content, the mass and linear ablation rate of the composites decreased. When the addition of Gd₂O₃ was 10 vol.%, the mass and linear ablation rates were 4.28 × 10⁻⁷ kg/s and 2.15 × 10⁻⁷ m/s, 80% and 31% lower than those of the Gd0 coating, respectively. These results suggest that the oxidative ablation resistance of the surface-modified composites was effectively improved by adding Gd₂O₃. There are three possible reasons for this. First, the modified coatings have a denser microstructure, as seen in Figure 2 and Figure 3. Second, the emissivity of the modified coatings is much higher than that of the coatings without modification, as shown in Figure 4. Third, the ablative oxidation products of the modified coatings are more stable, producing a better oxygen barrier.

3.3.2. Microstructure Analysis of the Ablation Layer

Figure 8 shows the surface morphology and related EDS results of the ablated Gd0 and Gd5 coatings. After ablation for 180 s, a porous structure formed on the surface of the ablated Gd0 coating (Figure 8a,b), and these open pores have a size larger than 10 μm. In addition, the main phase is composed of Hf–Ta–O in the Gd0 coating in Figure 8c and no SiO₂ phase is detected. At a surface temperature of ~2173 K in the Gd0 coating, the SiO₂ glass phase volatilizes violently, which causes serious structural damage to the surface, resulting in a large number of open pores. Therefore, the oxide penetrates the internal coating and causes a high mass ablation rate of the Gd0 coating (see Figure 7).

With the incorporation of Gd₂O₃, the compactness of the oxide layer increases significantly, as indicated by the surface morphology in Figure 8d–l. For the Gd5 coating, it can be seen that there is a grayish-black liquid droplet that covers the surface of the cluster particle tissue and occluded pores, and other defects can be seen on the surface, as shown in Figure 8d,e. In Figure 8f, it can be inferred that the grayish-black liquid droplet is a borosilicate glass phase, which is mainly composed of SiO₂ and B₂O₃, and the presence of this self-healing glass phase maintains the sealing property of the oxide layer.

For the Gd10, Gd15, and Gd20 coatings, it was found that the characteristics of glass-phase coverage are gradually weakened compared with those of Gd5, and the surfaces are mainly characterized by particle clusters (see Figure 8g,l). There are large micropores (<1 μm) observed on these particle clusters under high magnification for the ablated Gd10, Gd15, and Gd20 coatings (see Figure 8h,j and Figure 8l, respectively). The size of the micropores decreased with increasing Gd₂O₃ content. The pores suggest the possible formation of a liquid phase during oxidation. The liquid phase eventually becomes volatilized as the local temperature exceeds the evaporation temperature. This liquid phase is likely B₂O₃, SiO₂, or a combination of both. Owing to the strong sealing of the coating resulting in a shallow depth of oxygen erosion, there is no internal oxide layer of sufficient depth to provide SiO₂ transport for the surface layer. When the oxidation temperature is higher than 2073 K or the oxygen partial pressure is low, SiC undergoes active oxidation to generate gaseous SiO, intensifying the generation of surface-layer pores.

Because of the similar surface morphology of the Gd10, Gd15, and Gd20 coatings, the Gd10 coating was used for detailed morphology analysis. Figure 9 shows a high-magnification image and the related EDS results for the ablated Gd10 coating shown in Figure 8g,h. Lighter and darker regions can be observed. In Figure 9b,c, it can be seen that spot1 (lighter region) is mainly composed of Hf, Ta, and Gd oxides, whereas spot2 (darker region) is mainly composed of Hf and Gd oxides and a small amount of Si. Combined with the fact that the surface of Gd10 still shows excellent oxidation resistance, it can be inferred that in addition to the sealing effect of the glass phase, the introduced Gd₂O₃ is also a key factor determining the oxidation resistance of the coating under the test conditions.

The cross-sectional morphology of the ablated coatings is shown in Figure 10. Compared with the cross-sectional morphology before ablation in Figure 3, the cross-sectional morphology of the five coatings changed significantly after the ablation test, and the Gd₂O₃-modified coatings show similar layered structure. Figure 11 shows the EDS point-to-point line scan obtained across the Gd10 coating, starting from the ablated surface (x = 0 μm) to the interior coating (x = 150 μm). The O content of the outer layer is significantly higher than that of the inner layer, indicating that the outer layer is an oxide layer where oxidation reactions occur, whereas the inner layer is non-oxidized. Therefore, the ablated coating can be separated into a loose porous oxidized layer and a dense non-oxidized layer, and there is a clear film between the two layers.

The Gd0 coating was completely oxidized and there were many loose areas and large holes in the cross-section. This may be due to internal B₂O₃ and SiO escaping to the surface layer as well as the depletion of SiC. During the escape, the volatile gas damages the particle structure, loosening the porous oxide layer. Because the surface temperature (~2173 K) of the Gd0 coating is above the glass-phase volatilization temperature (Figure 5), the glass-phase self-healing function cannot be performed, and, because of the large porosity of the Gd0 coating, oxygen quickly penetrates the inner coating during the ablation process, and the coating loses its oxygen barrier function.

With Gd₂O₃-modified coating, a layered structure can be observed. In the Gd5 coating, large pores similar to those in the Gd0 coating are also observed, and the formation of these pores is likely related to the depletion of SiC. However, the oxide layer of the Gd5 coating is significantly thinner than that of the Gd0 coating. This can be attributed to the fact that the addition of Gd₂O₃ increases the density of the Gd5 coating (Table 4), which, in turn, hinders the diffusion of oxygen into the interior of the coating.

In the Gd10 coating, the oxidized layer appears relatively dense, and the thickness of the oxidized layer is lower than that of the Gd5 coating. The microstructure of the Gd15 coating is similar to that of the Gd10 coating, which has a denser structure. In the Gd20 coating, the oxidized layer contains many pores and the thickness of the oxidized layer increases compared to that of the Gd10 coating, further revealing that excessive Gd₂O₃ content aggravates the oxidation of the coating.

The thicknesses of the oxidized and non-oxidized layers of all the ablated coatings were measured and are summarized in

Table 5. Thickness of oxidized and non-oxidized layers for different coatings.

Coating	Oxidized (μm)	Non-Oxidized (μm)
Gd0	250 ± 10	0
Gd5	80 ± 10	170 ± 10
Gd10	16 ± 10	234 ± 10
Gd15	25 ± 10	225 ± 10
Gd20	52 ± 10	198 ± 10

. It can be seen that the Gd10 and Gd15 coatings exhibit the best oxidation resistance among all the coatings, having impact macrostructure and the largest thickness of non-oxidized coatings after 180 s ablation. However, the Gd0 coating was completely oxidized during the ablation test. Therefore, doping with Gd₂O₃ was beneficial for oxidation resistance.

To further compare the differences between the oxidized and non-oxidized layers of the coatings, the Gd10 coating with the best oxidation resistance was selected for observation at higher magnification, as shown in Figure 12. The oxide layer on the upper side of the red dashed line in Figure 12a is loose and contains holes, whereas the non-oxide layer below the red line retains the sprayed state characteristics and is dense. Figure 12b shows a magnified cross-section of the oxide layer, which is composed of minor oxide particles and dark, glassy phases embedded in the interstices. The EDS analysis in Figure 12b shows that region 1 contains Hf, Ta, Gd, Si, and O, indicating that the particle organization is mainly a compound of Hf–Ta–Gd–Si–O, mainly Hf, Si, and O with a small amount of Gd, indicating that the glass phase is SiO₂ and HfSiO₄, consistent with the results obtained by ablating the surface. Figure 12c shows a magnified image of the non-oxidized layer wherein the interstitial-phase distribution of bright and dark bands in the non-oxidized layer, with no loose particles, indicates a good ablation protection effect of Gd10 coating under the oxyacetylene combustion flow test of 4400 kW/m².

Figure 13 shows the XRD patterns and probable phases of the ablated coatings. The phase compositions of the Gd0 coatings are primarily cubic c-HfO₂ (PDF#53-0560) and Hf₆Ta₂O₁₇ (PDF#44-0998, orthorhombic). HfO₂ undergoes a transition from monoclinic to cubic at 1650 °C accompanied by volume shrinkage and a transition from cubic to monoclinic, causing volume expansion during cooling. HfSiO₄ (PDF#77-1759, tetragonal, and spatial group structure is 141/amd (141)), GdTaO₄ (PDF#24-0441, monoclinal, and spatial group structure is 12/A), and m-HfO₂ (PDF#78-0049, monoclinic, and spatial group structure is P21/c(14)) were detected in the Gd5, Gd10, and Gd15 coatings. With an increased content of Gd₂O₃, Gd₂Hf₂O₇ (PDF#24-0425, cubic, and spatial group structure is Fm3m (225)) was detected in the Gd10, Gd15, and Gd20 coatings. Gd₂Hf₂O₇ has excellent high-temperature thermal stability of the oxide layer, which can reduce stress within the coating and hinder crack sprouting [31]. Considered together with the edge warping of the ablated Gd0 and Gd5 coatings, it can be assumed that the introduced Gd₂O₃ plays a role in stabilizing the HfO₂ phase, suppressing volume change in the coating owing to phase change and improving the thermal suitability of the coating and substrate.

3.4. Ablation Mechanism

During the ablation process, HfB₂, SiC, TaSi₂, and Gd₂O₃ may react as follows [32,33,34,35]:

$$\mathrm{H}\mathrm{f}{\mathrm{B}}_2\left(\mathrm{s}\right)+{5/2\mathrm{O}}_2\left(\mathrm{g}\right)=\mathrm{H}\mathrm{f}{\mathrm{O}}_2\left(\mathrm{s}\right)+{\mathrm{B}}_2{\mathrm{O}}_3\left(\mathrm{l}\right)$$

(3)

$${\mathrm{TaSi}}_2\left(\mathrm{s}\right)+{13/4\mathrm{O}}_2\left(\mathrm{g}\right)=1/2{\mathrm{Ta}}_2{\mathrm{O}}_5\left(\mathrm{s}\right)+2\mathrm{Si}{\mathrm{O}}_2\left(\mathrm{l}\right)$$

(4)

$${\mathrm{B}}_2{\mathrm{O}}_3\left(\mathrm{l}\right)={\mathrm{B}}_2{\mathrm{O}}_3\left(\mathrm{g}\right)$$

(5)

Ta₂O₅(s) = Ta₂O₅(l)

(6)

$$\mathrm{S}\mathrm{i}\mathrm{C}\left(\mathrm{s}\right)+2{\mathrm{O}}_2\left(\mathrm{g}\right)=\mathrm{S}\mathrm{i}{\mathrm{O}}_2\left(\mathrm{l}\right)+\mathrm{C}{\mathrm{O}}_2\left(\mathrm{g}\right)$$

(7)

$$2\mathrm{S}\mathrm{i}\mathrm{C}\left(\mathrm{s}\right)+3{\mathrm{O}}_2\left(\mathrm{g}\right)=2\mathrm{S}\mathrm{i}{\mathrm{O}}_2\left(\mathrm{l}\right)+2\mathrm{C}\mathrm{O}\left(\mathrm{g}\right)$$

(8)

$$\mathrm{S}\mathrm{i}\mathrm{C}\left(\mathrm{s}\right)+{\mathrm{O}}_2\left(\mathrm{g}\right)=\mathrm{S}\mathrm{i}\mathrm{O}\left(\mathrm{g}\right)+\mathrm{C}\mathrm{O}\left(\mathrm{g}\right)$$

(9)

$$\mathrm{S}{\mathrm{i}\mathrm{O}}_2\left(\mathrm{l}\right)=\mathrm{S}{\mathrm{i}\mathrm{O}}_2\left(\mathrm{g}\right)$$

(10)

$$\mathrm{Hf}{\mathrm{O}}_2\left(\mathrm{s}\right)+\mathrm{Si}{\mathrm{O}}_2\left(\mathrm{l}\right)=\mathrm{HfSi}{\mathrm{O}}_4\left(\mathrm{s}\right)$$

(11)

$$\mathrm{G}{\mathrm{d}}_2{\mathrm{O}}_3\left(\mathrm{s}\right)+2\mathrm{Hf}{\mathrm{O}}_2\left(\mathrm{s}\right)=\mathrm{G}{\mathrm{d}}_2\mathrm{H}{\mathrm{f}}_2{\mathrm{O}}_7\left(\mathrm{s}\right)$$

(12)

$$\mathrm{G}{\mathrm{d}}_2{\mathrm{O}}_3\left(\mathrm{s}\right)+{\mathrm{Ta}}_2{\mathrm{O}}_5\left(\mathrm{s}\right)=2\mathrm{GdTa}{\mathrm{O}}_4\left(\mathrm{s}\right)$$

(13)

$${\mathrm{Ta}}_2{\mathrm{O}}_5\left(\mathrm{s}\right)+6\mathrm{Hf}{\mathrm{O}}_2\left(\mathrm{s}\right)={\mathrm{Hf}}_6{\mathrm{Ta}}_2{\mathrm{O}}_{17}\left(\mathrm{s}\right)$$

(14)

To analyze the oxidation mechanism of the coatings, the Gibbs free energies (ΔG) for reactions R3-R4, R7-R9, and R11 were calculated, as shown in Figure 14. The related thermodynamic data were obtained from FactSage 8.2 software. From Figure 14, in the temperature range of 500~2200 K, the ΔG of all reactions is negative, indicating all the reactions can occur spontaneously. From the thermodynamic calculation results, it indicates that TaSi₂ of the outer layer is preferentially oxidized to form a porous Ta₂O₅ layer owing to its lower Gibbs free energy, and the oxidation trend of HfB₂ is slightly weaker than TaSi₂. Owing to the lower ΔG of reaction 7 than those of reaction 3-4, the generated silica glass with a good self-healing ability can fill defects of the loose structure and porous Ta₂O₅ and HfO₂ forming a relatively dense outer layer.

In the Gd0 coating, the HfB₂/SiC/TaSi₂ coating is oxidized to HfO₂, Ta₂O₅, and SiO₂ according to reactions (3), (4), and (7). HfO₂ reacts with Ta₂O₅ to generate a new stable phase (Hf₆Ta₂O₁₇) (reaction (14)). When the ablation temperature (2173 K) is higher than the melting points of Ta₂O₅ (~2058 K) and SiO₂ (1873–1973 K), a portion of SiO₂ and Ta₂O₅ melt and evaporate. Evaporation of the gaseous products (CO, CO₂, B₂O₃, SiO₂, and Ta₂O₅) leaves small pores in the coating. The Ta–Si–O glass layer can seal some of these voids. Because Hf₆Ta₂O₁₇ has low thermal conductivity (2.89 W/m K) and a high melting point (~2723 K), Hf₆Ta₂O₁₇ particles play a pinning role in the Ta–Si–O glassy layer, inhibiting the formation and propagation of cracks [36]. Unfortunately, the porous structure of the original Gd0 coating (see Table 2) offers channels for oxygen diffusion, thereby degrading the ablation performance (see Figure 5 and Figure 8).

For the HfB₂/SiC/TaSi₂/Gd₂O₃ coating, the introduction of Gd₂O₃ results in a lower surface temperature under the same test conditions due to the increased emissivity, which directly reduces the oxidation rate of the coating and the volatilization rate of the glass phase. Moreover, the melting point of Gd₂O₃ (~2573 K) is significantly higher than those of Ta₂O₅ (~2058 K) and SiO₂ (1873–1973 K), avoiding the negative impact of lowering the oxide eutectic temperature on the thermal stability of the oxide layer. When the surface temperature is 2073 K, Gd₂O₃ replaces Ta₂O₅ and SiO₂ to play a role in stabilizing the HfO₂ phase transition owing to its high melting point, thus promoting a dense oxide layer. Gd₂O₃ reacts with Ta₂O₅ to generate GdTaO₄ (reaction 12), and HfO₂ reacts with SiO₂ to generate a new stable phase (HfSiO₄) (reaction 11). In the Gd5 coating, HfSiO₄, GdTaO₄, and HfO₂ are the main phases on the ablated surface and have a pinning effect on the glass layer because of their high melting points. Therefore, the Gd5 coating has a thinner oxidation layer than the Gd0 coating (Figure 10b); however, the oxidation layer exhibits a porous structure owing to the presence of a small amount of Gd₂O₃.

With an increasing amount of Gd₂O₃, a new phase of Gd₂Hf₂O₇ (melting point ~2573 K) is generated. The Gd₂Hf₂O₇ phase can suppress the volume expansion accompanying the HfO₂ cubic/monoclinic phase transition during the cooling process and reduce stress concentration in the coating. However, the decomposition temperature of Gd₂Hf₂O₇, with a cubic pyrochlore structure, is ~2373 K [17], so it can maintain phase stability when T ≥ 2073 K. In the Gd10 and Gd15 coatings, HfSiO₄, GdTaO₄, HfO₂, and Gd₂Hf₂O₇ are the main phases. The dissolution of higher-melting-point oxides forms a stable Hf–Ta–Gd–Si–O multiphase glass on the surface of the coating (see Figure 9c). High-melting-point particles, such as HfO₂, GdTaO₄, HfSiO_4, and Gd₂Hf₂O₇, are dispersed in the Hf–Ta–Gd–Si–O glassy layer as “hard particles” when the glass layer is softened during oxidation at high temperatures. As a result, the Gd10 and Gd15 coatings exhibit better resistance to oxidation than the Gd5 coating (See Figure 10c,d).

However, excessive Gd₂O₃ can cause excessive consumption of SiO₂ in the Gd20 coating, resulting in the formation of more Gd₂Hf₂O₇, making the generated glass very viscous and imparting its self-healing ability. Therefore, doping with a high content of Gd₂O₃ (>15 vol.%) negatively affects the oxidation-blocking capability of the HfB₂/SiC/TaSi₂ coating.

4. Conclusions

Gd₂O₃-modified HfB₂/SiC/TaSi₂ coatings were successfully prepared by APS. Through the oxyacetylene flame test, the ablation resistance of the modified coatings was evaluated at 4400 kW/m² for 180 s. The effects of varying Gd₂O₃ contents on the microstructure, infrared radiative property, and ablation resistance of the HfB₂/SiC/TaSi₂ were investigated. With the growth of Gd₂O₃ content increased from 0 vol.% to 20 vol.%, the coating porosity decreased from 19.91% to 11.28%, indicating that the addition of Gd₂O₃ improved the coating density. Moreover, because of the additional electronic transitions promoted by oxygen vacancies in HfO₂, the addition of Gd₂O₃ improved the infrared emissivity of the HfB₂/SiC/TaSi₂ coatings. The Gd10 coating possessed the highest emissivity in the 3–5 μm band at 1273 K, up to 0.92. Due to the highest emissivity, the Gd10 coating possessed the lowest surface temperature of 2073 K. Thus, the Gd10 coating had mass and linear ablation rates of 4.28 × 10⁻⁷ kg/s and 2.15 × 10⁻⁷ m/s, 80% and 31% lower than those of the Gd0 coating, respectively. Further, the Gd10 coating showed excellent oxygen barrier performance due to the stable Hf–Ta–Gd–Si–O glass formed on the surface of the Gd10 coating, which inhibited oxide penetration into the internal coating. Nevertheless, an excessive amount of Gd₂O₃ is detrimental to the ablation performance owing to its excessive consumption of SiO₂. This study demonstrates that the addition of an appropriate amount of Gd₂O₃ is an effective way to improve the ablation resistance of HfB₂/SiC/TaSi₂ coatings at a heat flux of 4400 kW/m².