Preparation of ZrB2-MoSi2 high oxygen resistant coating using nonequilibrium state powders by self-propagating high-temperature synthesis

To achieve high oxygen blocking structure of the ZrB2-MoSi2 coating applied on carbon structural material, ZrB2-MoSi2 coating was prepared by spark plasma sintering (SPS) method utilizing ZrB2-MoSi2 composite powders synthesized by self-propagating high-temperature synthesis (SHS) technique as raw materials. The oxygen blocking mechanism of the ZrB2-MoSi2 coatings at 1973 K was investigated. Compared with commercial powders, the coatings prepared by SHS powders exhibited superior density and inferior oxidation activity, which significantly heightened the structural oxygen blocking ability of the coatings in the active oxidation stage, thus characterizing higher oxidation protection efficiency. The rise of MoSi2 content facilitated the dispersion of transition metal oxide nanocrystals (5–20 nm) in the SiO2 glass layer and conduced to the increasing viscosity, thus strengthening the inerting impact of the compound glass layer in the inert oxidation stage. Nevertheless, the ZrB2-40 vol%MoSi2 coating sample prepared by SHS powders presented the lowest oxygen permeability of 0.3% and carbon loss rate of 0.29×10−6 g·cm−2·s−1. Owing to the gradient oxygen partial pressure inside the coatings, the Si-depleted layer was developed under the compound glass layer, which brought about acute oxygen erosion.


Introduction 
Carbon structural materials are widely exploited in aerospace high-temperature components, as indispensable structural materials in the national defense and civil fields, due to the low density, high specific strength, temperature [8][9][10]. However, ZrB 2 is challenging to achieve densification due to its strong covalent bond and low diffusion coefficient [11,12], which would also become porous during the oxidation process [13]. To suppress such defects, MoSi 2 , the most commonly used sintering additive, was generally used to increase the compactness of ZrB 2 [14,15]. Moreover, the excellent film-forming properties of MoSi 2 can effectively reduce the oxidation consumption of ZrB 2 , thus making the ZrB 2 -MoSi 2 coating present great potential [16][17][18]. Recently, pressureless sintering methods are adopted to prepare the ZrB 2 -MoSi 2 coatings, exhibiting great potential. Jiang et al. [18] successfully prepared ZrB 2 -MoSi 2 -SiC-Si coating on graphite matrix by slurry dipping and vapor silicon infiltration methods. After oxidation at 1873 K for 150 h, the mass loss of the coating sample was only 0.21%. Yao et al. [19] used two-step pack cementation method to prepare SiC/SiC-MoSi 2 -ZrB 2 coating on C/C composites, and the weight loss after oxidation at 1773 K for 510 h was 0.5%. Zhang et al. [20] prepared ZrB 2 -MoSi 2 /SiC coating on C/C composites by combining the pack cementation method and the slurry painting method. After oxidation at 1273 and 1773 K for 30 and 10 h, the weight loss rates of the samples were 5.3% and 3.0%, respectively. Nevertheless, pressureless sintering generally requires ultra-high sintering temperature (above 2200 K) to achieve densification, which will lead to the generation of coarse grains, thereby weakening the oxidation resistance of the coatings.
Considering the oxidation protection of the coatings is a dynamic process, the structural oxygen blocking ability of the coating itself and the inert oxygen blocking ability of the self-generated compound glass layer constitute the oxidation protection of the coating for carbon matrix. In our previous work, we have found that the dispersion of Zr-oxides in the Zr-Si-B-O compound glass layer would strengthen the oxidation protection effect of the coating [21]. Nonetheless, for the ZrB 2 -MoSi 2 coatings, the by-products B 2 O 3 and MoO 3 have intense volatility, and abundant gas volatilization will inevitably destroy the stability of the compound glass layer. Thus, it is necessary to improve the dynamic oxygen blocking ability of the ZrB 2 -MoSi 2 coating by constructing high oxygen resistant coating structure of the coatings. However, little work has been reported about the construction of ZrB 2 -MoSi 2 high oxygen blocking coating structure.
To improve the structural oxygen blocking ability of the coating, it is necessary to develop densification sintering technique at a lower temperature [22]. Spark plasma sintering (SPS) is well known as the low-temperature pressured rapid sintering technology, which can remarkably inhibit grain growth as well as bringing high density under pressure conditions [23][24][25]. In our previous work, HfB 2 -SiC-TaSi 2 coating [22] and ZrB 2 -SiC coating [13] have been prepared by SPS technology, both of which reveal excellent hightemperature oxidation resistance.
Utilizing high activity powders to improve the sintering performance is another way to enhance the structural oxygen blocking ability of the coatings. Recently, self-propagating high-temperature synthesis (SHS) method has been widely exploited in the preparation of composites [26][27][28][29][30], which has superior heating and cooling rates, being conducive to generate high activity due to the concentration of defects and non-equilibrium phases in the products. Iatsyuk et al. [28] used SHS method to prepare ZrB 2 -SiC ceramics with fine grains and high density. The porosity was only 1.5%, and the Vickers hardness was 24.3-25.0 GPa. Licheri et al. [31] prepared 40 vol%ZrB 2 -12 vol%ZrC-SiC composites by SHS method, with a relative density of 99.5%, the Vickers hardness of 16.9±0.2 GPa, and the fracture toughness of 5.9±0.5 MPa·m 1/2 .
Although the combination of SPS and SHS methods to prepare ZrB 2 -MoSi 2 high oxygen resistant coating exhibits much potential, little work has been reported. Furthermore, the inhibition mechanism of the ZrB 2 -MoSi 2 coating for penetration of oxygen is still not clear, which can be used to fundamentally regulate the coating structure to obtain better oxygen blocking effect. Therefore, in this paper, ZrB 2 -MoSi 2 composite powders were first prepared by SHS technique, using which ZrB 2 -MoSi 2 coatings were prepared on graphite matrix by SPS technique. To avoid the interference of matrix oxidation consumption and reflect the authentic law of protection of the ZrB 2 -MoSi 2 coatings, graphite samples, coating standard samples, and coating samples were prepared simultaneously. Furthermore, the oxygen blocking ability of the ZrB 2 -MoSi 2 coatings prepared by SHS and commercial powders (CP) was quantitatively compared at 1973 K. The microstructure and forming mechanism of the self-generated compound glass layer were also further analyzed.

1 Synthesis of ZrB 2 -MoSi 2 powders by SHS
In this experiment, Zr (48 μm, 99.9% purity), B (1-3 μm, 95% purity), Mo (1-3 μm, 99.95% purity), and Si (1-3 μm, 99.9% purity) powders were used to prepare ZrB 2 -MoSi 2 composite powders by SHS technique, which were manufactured by Beijing Licheng Innovation Metal Materials Technology Co., Ltd., China. In order to explore the influence of MoSi 2 content on oxygen blocking of the coatings, three components were designed, which were 60 vol%ZrB 2 -40 vol%MoSi 2 (ZM40), 70 vol%ZrB 2 -30 vol%MoSi 2 (ZM30), and 80 vol%ZrB 2 -20 vol%MoSi 2 (ZM20). To guarantee the accurate phase contents of ZrB 2 -MoSi 2 composite powders, we designed the content of the original powders according to the volume ratio of ZrB 2 and MoSi 2 . Primarily, the weighed powders were put into ball mill with the ball-to-powder weight ratio of 5:1, which were milled at 300 r/min for 5 h. The mixed powders were dried in an oven at 333 K for 12 h. Subsequently, the dried powders were cold-pressed into a cylinder of 16 mm × 20 mm at a pressure of 200 MPa for 90 s. Afterward, the pressed sample was placed on the sample table of self-propagating reactor, whose upper surface was closely attached with molybdenum wire. Moreover, the reactor was vacuumized before introducing Ar gas as protective gas. After the protective gas was filled, the voltage regulating system applied current to the molybdenum wire, and then gradually preheated the sample until the self-propagating reaction occurred. When the combustion began, the voltage was turned off immediately until the self-propagation reaction of the sample was completed. Lastly, the composite powders obtained after combustion were crushed and grounded for 1 h, and then placed in a 333 K oven to dry for 12 h.

2 Sample preparation
To compare the difference in oxidation resistance between CP and SHS powders, ceramic samples (φ 12 mm) and coating samples (φ 15 mm) of ZM20-ZM40 were prepared by SPS furnace (Labox-125, Sinter Land Inc., Japan). The sintering temperature, pressure, heating rate, and holding time were 1773 K, 30 MPa, 180 K/min, and 5 min, respectively. The size of graphite matrix was Ф12 mm × 2.5 mm, and the density was 1.6 g/cm 3 . ZrB 2 (1-3 μm, 99.9% purity) and MoSi 2 (1-3 μm, 99.9% purity) of CP were also manufactured by Beijing Licheng Innovation Metal Materials Technology Co., Ltd., China. After SPS sintering, the graphite paper on the surface of the sample was polished clean with sandpaper to avoid data inaccuracy caused by the loss of graphite paper during the oxidation process.

3 Oxidation tests and microstructure characterization
The isothermal temperature oxidation test was carried out in a resistance furnace (SGMM4/19AE, Luoyang, China) at 1973 K. During the oxidation process, the analytical balance with sensitivity of 0.1 mg was utilized to weigh the crucibles containing the samples, and the oxidation kinetics, permeability, and protection efficiency of coatings and coating standard samples were further explored through the change of sample weight. Moreover, X-ray diffraction analysis (XRD, Bruker D8 ADVANCE XRD, BRUKER AXS, Germany) was used to detect the phase of the surface, an scanning electron microscope (SEM, FEI Quanta 250, FEI, USA) was exploited to analyze the microstructure, and its corresponding energy dispersive spectrometer (EDS, FEI Quanta 250, USA) scrutinized the composition and distribution of the phase, electron probe X-ray microanalysis (EPMA-8050G, SHIMADZU, Japan) was used for accurate qualitative analysis of the element distribution, and the transmission electron microscope (TEM, Tecnai G2 F20, FEI, USA) was utilized to evaluate the phase structure of the powders synthesized by the SHS technique. The cross-section backscatter SEM and cross-section EPMA samples were polished by SiC sandpapers with various specifications (100, 400, 800, and 1200 mesh), then ultrasonically cleaned with ethanol and placed in a 333 K oven to dry for 2 h.

4 Evaluation of oxygen blocking ability
To quantitatively evaluate the antioxidant capacity of the ZrB 2 -MoSi 2 system, based on our previous work [22], Eqs. (1)-(6) characterized the dynamic protection process of the coatings in various oxidation stages.
where a  represents the active factor, M  is the change of sample weight, T is the oxidation time, and C is a constant.
where  represents the protective efficiency; α s is the structure factor; i  is the inerting factor; p, j m , s, j m , and c, j m are the weight gain per unit area of pure graphite, coating standard sample, and coating sample at j times of weighing, respectively; and t is the weighing interval (t = 10 min).
where g V is the carbon loss rate; P is the oxygen permeability; c, j V , s, j V , and p, j V are the weight change rate of coating sample, coating standard sample, and graphite matrix in j times of weighing, respectively. Figure 1 shows the XRD patterns of ZrB 2 -MoSi 2 composite powders with various components synthesized by SHS technique. The synthesized powders consist only of ZrB 2 (#89-3930) and MoSi 2 (#41-0612), and no residual elemental phase is found, which is similar to the self-purify characteristics of SHS process in high-temperature combustion [32]. Figure 2 indicates the TEM and EDS micrographs of pure coating standard samples sintered by SPS technique. It can be found by EDS spectrum that there are only ZrB 2 and MoSi 2 phases in the samples, which also proves that there is no phase transformation in the coating after SPS sintering. Furthermore, according to Fig. 2(a), the average grain size of the coating standard samples prepared by SHS powders after SPS sintering is merely 10-20 nm, further demonstrating that the SHS and SPS techniques can decrease the grain size and fine grain strengthening [33]. It can also be found from Fig. 2(b) that there is no noticeable phase interface between the various phases of the samples after sintering, which realizes superior metallurgical bonding.

1 Microstructure of ZrB 2 -MoSi 2 powders and coating
In order to further characterize the bonding performance of coating and substrate before oxidation, the cross-section backscatter SEM micrographs of CP and SHS coating samples were compared in Fig. 3. It can be seen that CP-ZM20 component ( Fig. 3(a)) has severe cracks after sintering, and noticeable loosening is found in its enlarged local diagram ( Fig. 3(g)), which is related to the sintering characteristics of ZrB 2 [16]. Although SHS-ZM20 component also has a small number of microcracks ( Fig. 3(b)), the phase structure after sintering is more compact than that of commercial powders. Additionally, with the increase of MoSi 2 content, superior metallurgical bonding is achieved between the coatings and the substrates, the various phases are evenly distributed, and there is no apparent aggregation phenomenon, especially the SHS-ZM40 component ( Fig. 3(f)). This phenomenon is closely related to the sintering performance, which further  verifies that the SHS synthesized powders have superior sintering activity.

2 Oxidation behavior comparison of SHS and CP coatings at 1973 K
In view of the interference caused by the oxidation loss of graphite matrix, the weight change of coating samples cannot reflect the actual law of protection. Therefore, the oxidation kinetics of SHS and CP coating standard samples at 1973 K were studied. Figures 4(a) the oxidation resistance of the coating standard samples in the initial oxidation process. The smaller the a  is, the more excellent the oxidation resistance of the coating itself is. Figure 5(a) displays the weight gain rate curves of coating standard samples oxidized at 1973 K for 100 min. It can be seen from Fig. 5(a) that the weight of the samples first increases rapidly and then tends to be stable. In the active oxidation stage, the ZrB 2 -MoSi 2 composite coating surface has not yet formed steady glass layer, which predominantly relies on the coating structure itself to block oxygen. With the prolongation of the oxidation time, the glass layer on the coating surface gradually forms and tends to be steady, and the oxidation enters the inert oxidation stage. At this time, the oxygen blocking is fulfilled by the joint action of the coating itself and the compound glass layer, in which the oxidation rate of the coating would be significantly reduced. In order to scrutinize the impact of the reduction of oxidation activity on the protection ability, based on Eq. (2), we further investigate the oxidation protection efficiency of the coating with various components. As shown in Fig. 5(b), the protective efficiency of all the samples decreases first and then increases. This is mainly because, in the active oxidation stage, the suppression of oxygen by the coating is mainly achieved by its own structure. In contrast, in the inert oxidation stage, with the extension of oxidation time, the inhibition of oxygen by the coatings is significantly enhanced due to the formation of self-healing glass layer, thus expanding the oxidation protection efficiency of the coatings.
In addition, the inflection point of the oxidation protection efficiency curve of the CP coating appeared in about 10 min, while that of the SHS coating prepared with SHS powders appeared in 20-30 min. The prolongation of the turning point time is chiefly due to the significant decrease of the oxidation activity of the coatings prepared by SHS powders, which delays the formation of self-healing glass layer on the coating surface. However, although the inflection point of oxidation protection efficiency of SHS coating appears late, the lowest value of oxidation protection efficiency at the inflection point is far higher than that of the CP coating. Figure 5(c) shows the density of coating samples with various components and the difference between the actual density and the theoretical density. As shown in Fig. 5(c), the density of SHS coating samples is higher than that of CP coating samples, and with the increase of MoSi 2 content, the difference between the actual density and the theoretical density gradually decreases, from 1.54 to 1.09 g/cm 3 (ZM20), from 1.46 to 1.03 g/cm 3 (ZM30), and from 1.10 to 0.79 g/cm 3 (ZM40), respectively. Therefore, the improvement of coating density can effectively reduce the oxidation activity of coatings. Although it will delay the inflection point time, it can effectively improve the oxidation protection efficiency of the coating. Because of the close relationship between the enhancement of coating density and its structural oxygen blocking ability, based on Eq. (3), as shown in Fig. 5(d), we calculate the structure factor (α s ) to evaluate the oxygen blocking of the coating itself. The smaller the structure factor, the more solid the oxygen blocking of the coating structure. As can be seen from Fig. 5(d), the structure factors of all components of SHS coatings were lower than that of CP coatings, which decreased by 59.82% (ZM20), 20.44% (ZM30), and 44.95% (ZM40), respectively, indicating that the coating prepared by SHS powders could effectively heighten the structural oxygen blocking ability of the coatings.
However, the stability of glass layer is also highly crucial as another oxygen blocking core in the inert oxidation stage. In order to avoid the distortion of the oxidation mechanism caused by internal matrix oxidation, the inert oxygen blocking ability of the glass layer is evaluated by coating standard samples. According to Eq. (4), the smaller the inerting factor, the greater the oxygen blocking of the glass layer. As seen from Fig. 6(a), with the increase of MoSi 2 content, both the inerting factors of coating standard samples prepared by CP and SHS powders diminish continuously, and the inerting factors of CP coating standard samples are significantly taller than those of SHS standard samples, revealing that the samples prepared by SHS powders have superior inerting ability.
With the intention of analyzing the effect of the surface glass layer on the inert oxygen blocking, Figs. 6(b)-6(g) show the cross-section backscatter SEM and EDS micrographs of various CP and SHS coating standard samples after oxidation at 1973 K for 100 min. Compared with the SHS components, CP coating standard samples exposed severe oxidation erosion, especially the CP-ZM20 component ( Fig. 6(b)). There are large transverse cracks and evident loosening phenomena inside the samples. By comparing the distribution of O elements in the three components, the thickness of oxygen diffusion depth of CP-ZM20, CP-ZM30, and CP-ZM40 components reduces in turn, which are 345.6, 186.1, and 133.8 μm, respectively   Fig. 6(h)). However, for SHS components, there are no apparent crack, pore, or other defect after oxidation. The thicknesses of the oxygen diffusion depth enlarge with the increase of MoSi 2 content, which are 38.68, 46.69, and 61.32 μm, respectively. It is easy to spot that with the increase of MoSi 2 content, the diffusion depth of O element in CP samples decreases. This is because the antioxidant capacity of the coating depends to a large extent on the coating structure after the sintering. In Fig. 5(c), we have found that the density of sintered CP powders is lower than that of SHS powders, and the loose structure of low-density samples leads to the increase of oxygen diffusion depth and even obvious cracks ( Fig. 6(b)). The increase of MoSi 2 content can expand the density of the samples, and form the thicker SiO 2 glass layer, which greatly limits the diffusion of oxygen in the interior. Nevertheless, for the SHS samples, the dense structure causes the oxygen blocking mechanism to shift from the structure to the component dominated. Because the components with high ZrB 2 content will produce more gas by-products due to oxidation, leading to the formation of more oxidation pores on the surface [12]. Meanwhile, the lower content of MoSi 2 leads to the lack of sufficient SiO 2 to pad these oxidation pores, resulting in a thinner glass layer. ZM40 component can rapidly shape steady SiO 2 glass layer at the same time owing to sufficient silicide and relatively low boride. When the oxygen diffusion is not severe, the increase of glass layer thickness enhances the stability of the samples. It shrinks the possibility of oxygen diffusion into the interior, and thus attains better oxygen blocking effect, which also clarifies why the SHS-ZM40 component in Fig. 6(a) exhibits the lowest inerting factor. Therefore, it is further confirmed that the SHS powders have lower oxidation activity than the CP powders at the ultra-high temperature of 1973 K.
In order to reveal the specific effect of the glass layer on the oxygen blocking process, Fig. 7 shows the XRD patterns of the SHS coating samples with various components after oxidation at 1973 K for 100 min. After oxidation, the ZrB 2 and MoSi 2 phases on the surface of the samples completely disappear, with only ZrO 2 , SiO 2 , MoB (#51-0940), ZrSiO 4 (#33-1485) phases present, and no obvious Mo 5 Si 3 phase appears. ZrO 2 is dominated by monoclinic structure (#72-1669) and a small amount of tetragonal structure (#79-1768), which is consistent with Ref. [34]. Moreover, it can be easily observed from Fig. 7 that the peak of SiO 2 amorphous state becomes more and more evident with the increase of MoSi 2 content. As we all known, MoO 3 , as a by-product of MoSi 2 , begins to volatilize at 923 K [35], while B 2 O 3 , due to its low melting point (723 K), also has intense volatility, especially at the ultra-high temperature of 1973 K, which clarifies why B 2 O 3 and MoO 3 phases are not detected in XRD. The possible reactions in the oxidation process are shown in Eqs. (R1)-(R9) [35,36].
On the other hand, the film-forming quality of the compound glass layer on the coating surface is associated with the internal structure of the coatings and plays a crucial role in the high oxygen blocking system. Therefore, Fig. 8 describes the surface SEM micrographs of SHS and CP coating samples after oxidation at 1973 K for 100 min. Combined with XRD and EDS, it can be discovered that the black and gray phases are SiO 2 , the white phases are ZrO 2 , ZrSiO 4 , and the bright white phases are MoB, which are dispersed in the glass layer on the surface. Based on its partial enlarged diagram, it can be detected that the oxides are not uniformly distributed, but appear local agglomeration phenomenon. The reason is that plenty of ZrSiO 4 and ZrO 2 are precipitated from the high content of ZrB 2 , the low content of silicide leads to the insufficient supply of SiO 2 glass on the surface, and the oxide particles cannot be homogeneously dispersed, resulting in the oxide agglomeration. Furthermore, the content of the gas by-product B 2 O 3 will increase accordingly, and a large amount of pores appear in the coating, which provides diffusion channels for the infiltration of oxygen [37]. The CP-ZM20 (Figs. 8(a) and 8(b)) component has large cracks on the surface due to its higher oxidation activity, which also explains why the component in Fig. 5(b) shows the lowest oxidation protection efficiency.
However, the surface of the ZM40 coating sample (Figs. 8(i)-8(l)) tends to be flat, without obvious holes and cracks. With the increase of MoSi 2 content, the oxygen diffusion rate of oxidation product SiO 2 is relatively low [35]. At the ultra-high temperature of 1973 K, the pores caused by gas by-products are continuously filled with SiO 2 . Therefore, there is no obvious aggregation of Zr-oxides and Mo-oxides on the surface. Although there are a few unhealed microcracks on the surface of the CP-ZM40 (Figs. 8(i) and 8(j)) component, compared with the component with low content of MoSi 2 , the flatness of the surface can still characterize the outstanding stability of ZM40 components.
As shown in Figs. 8(e)-8(h), the white phases on the surface of ZM30 coating sample are appreciably reduced, and bulky holes are created simultaneously. Among them, the surface of CP-ZM30 formed significantly more white oxides than SHS-ZM30 coating sample, which is related to the coating structure after oxidation. It has been found in Fig. 6(c) that CP-ZM30 component has more severe oxygen erosion during the oxidation, giving rise to the breeding of a large amount of transition metal oxides. In addition, in the enlarged micrographs, it can also be found that there are lots of ZrSiO 4 and ZrO 2 in the pores, which indicates that the Zr-oxides may not only precipitate on the surface, but may be dispersed in the SiO 2 glass layer, thereby forming the compound glass layer with excellent oxygen blocking.

3 Oxygen blocking mechanism of coatings
To deeply analyze the structural changes of the coatings after oxidation, Fig. 9 shows the cross-section EPMA micrographs of SHS coating samples after oxidation at 1973 K for 100 min. Based on Fig. 9, the interior of the coating is still dominated by ZrB 2 , and there are no penetrating cracks or large holes, indicating that the three coating samples of ZM20-ZM40 display excellent effects of blocking oxygen penetration at 1973 K. However, an interesting phenomenon is likewise discovered in the spectra of Si element, that is, all coating samples have the absence of Si elements under the glass layer, which is called Si-depleted layer. The Si-depleted layers of ZM20-ZM40 component coating samples are 165.7, 95.03, and 64.09 μm, respectively, which is the same as that of coating standard samples. In fact, Fahrenholtz [38] has proved that the appearance of the depletion layer is not only related to pressure and temperature, but also associated with the structure after sintering. Moreover, the SiO 2 glass is in molten state at 1973 K, the gas by-products of ZrB 2 and MoSi 2 will accelerate the flow inside the coating, the convection of liquid glass inside the coating will lead to the delamination in the oxidation region [24,39], and the low oxygen partial pressure in the coating will lead to the violent volatilization of silicon ( Fig. 10(a)), resulting in the appearance of Si-depleted layers. In the interior of the coating samples, the O element penetration of the SHS-ZM20 component is supreme, which further verifies that the oxidation pores formed by high content of boride will turn into the channel of oxygen diffusion, resulting in generous oxygen penetration. With the www.springer.com/journal/40145 increase of MoSi 2 content, the degree of oxygen penetration into the interior decreases in turn, establishing the relatively straightforward layer structure, which is consistent with the change of glass layer thickness in Fig. 6. It further confirms that the lower the inerting factor of the coating samples, the bigger the thickness of the glass layer, and the better the inert protection effect of the glass layer.
For the delamination phenomenon in EPMA micrographs (Fig. 9), the volatile phase generated during the oxidation process plays a crucial role. Consequently, according to the thermodynamic model constructed by Fahrenholtz [40] and the ZrB 2 -SiC volatility diagram calculated by Kubota et al. [24], Fig. 10(a) plots the volatility diagram of volatile phases (MoO 3 , B 2 O 3 , SiO 2 , and SiO) that may exist in the oxidation process of coating samples at 1973 K; other secondary volatile phases (ZrO 2 , MoB, etc.) have minor influence, and hence they are not taken into account [34], in which the thermodynamic data is provided by HSC chemistry software. Moreover, it can be seen from Fig. 10(a) that the vapor pressure of B 2 O 3 (g) in air is the tallest, 2420 Pa. And the vapor pressure of MoO 3 (g) is 822 Pa, which is inferior to B 2 O 3 (g). The equilibrium vapor pressure of SiO 2 (g) and SiO(g) is merely 1.57×10 -2 and 2.21×10 -3 Pa, respectively. Thus, when the oxidation enters the active oxidation stage, the oxygen partial pressure (p O 2 ) inside the coating is approximately equal to the oxygen partial pressure in the air (p O 2 = 0.2 atm). No matter the coating surface or the interface between the coating and the substrate, the oxidation reaction will take place in varying degrees. At this time, the oxidation products MoO 3 (g) and B 2 O 3 (g) in Eqs. (R1) and (R2) volatilize first. With the oxidation entering the inert oxidation stage, the well-sealed compound glass layer has been shaped on the surface, and both the volatilization of internal gas and the erosion of external oxygen have been efficaciously impeded, which makes the pressure inside the coating samples have observable gradient alteration, and the oxygen partial pressure from the coating surface to the inner matrix similarly demonstrates the phenomenon of gradual decrease.
Nonetheless, the boundary between diverse layers is the initial region of oxidation, and the oxygen partial pressure at equilibrium represents the necessary condition for oxidation, which determines the oxidation reaction sequence [23], and further reflects the change of coating structure. Consequently, the equilibrium oxygen partial pressure of Eqs. (R1)-(R6) at 1973 K was calculated by HSC chemistry, as shown in Fig. 10(b). The equilibrium oxygen partial pressure of Eq. (R5) is minimum at 1973 K, which is only 5.01× 10 -16 atm, indicating that MoSi 2 preferentially oxidizes with B 2 O 3 and O 2 to form MoB phase at low oxygen partial pressure [34], and Eq. (R6) is barely inferior to Eq. (R5). Furthermore, the simultaneous oxidation of ZrB 2 and MoSi 2 also preferentially forms stable MoB phase. While the Mo 5 Si 3 phase in R3 can also be shaped at insufficient oxygen partial pressure, when it reaches the surface of the compound glass layer, the oxygen partial pressure is similar to air (p O 2 = 0.2 atm), and Mo 5 Si 3 phase will continue to react with O 2 to generate MoO 3 volatile phase (Eq. (R4)). What is more, for the components with high boride content, due to the high oxidation activity of ZrB 2 and the lack of SiO 2 glass formation, extensive gas by-products are generated. In the volatilization process, the gas by-products convey the Si of the Si-depleted layer to the coating surface, bringing about the Si-depletion layer. With the rise of MoSi 2 content, the density of the coating enhances and the amount of gas by-products generated decrease, which also explains why the thickness of the Si-depleted layer is inversely proportional to MoSi 2 content in EPMA (Fig. 9), and only MoB phase is detected in XRD, but no other Mo-oxides are unearthed. Figure 11 presents the high-resolution TEM micrographs of the surface glass layer of the SHS-ZM40 coating sample after oxidation at 1973 K for 100 min. Figure 11(b) is the partial enlarged diagram of Fig. 11(a). Substantial nanocrystalline particles are dispersed in the compound glass layer with a diameter of about 5-20 nm. TEM micrographs can further confirm in Figs. 11(c) and 11(d) that the oxides scattered in the glass layer are primarily ZrO 2 and ZrSiO 4 . Although it is detected in Fig. 8(i) that white Zr-oxides precipitate on the surface, more Zr-oxides are embedded in the glass layer to establish steady compound glass layer.
However, in addition to the formation of stable compound glass layer on the surface, the focus of the high oxygen blocking system is the internal structure of the coating, which is the core determining whether the high oxygen blocking can be fulfilled. Figure 12 is the structure diagram of the ZrB 2 -MoSi 2 coating samples after oxidation at 1973 K for 100 min. Figure 11(b) has proved that the compound glass layer is composed of ZrO 2 and ZrSiO 4 nanocrystalline particles dispersed in the molten SiO 2 glass, while ZrO 2 has merely one  coordination number in the silicate glass, which is located in the cube [ZrO 8 ]. Due to the large ion radius, Zr 4+ is mainly outside the structural network in silicate glass, and the higher melting points of ZrO 2 (2973 K) and ZrSiO 4 (2773 K) have good stability. However, the content of Zr-oxides is not as high as possible. SiO 2 glass exists in the form of [SiO 4 ], which is connected through "bridging oxygen" to form the three-dimensional structure. Excessive ZrO 2 content in the glass will cause [SiO 4 ] to slightly deform and tend to be compacted, and thus the molecular volume is reduced, the thermal expansion coefficient is expanded, and the phenomenon of oxide accumulation occurs.
It can be seen from Fig. 10 that although B 2 O 3 , the gas by-product of ZrB 2 oxidation, has intense volatility at 1973 K, MoSi 2 can react with it to produce MoB, as shown in Eq. (R3), thereby reducing the pressure difference triggered by the formation of oxidation pores in the coating, and thus reducing the possibility of oxygen penetrating the substrate. It can also be attained from the EPMA micrographs in Fig. 9 that the Si-depleted layer decreases with the enhancement of MoSi 2 content, and the internal oxidation of the coating declines accordingly. Furthermore, the decrease of B 2 O 3 volatilization will correspondingly multiply the viscosity of the glass layer, hence avoiding the rapid migration of borosilicate glass due to low viscosity [34], and augmenting the stability of the compound glass layer.
However, to evaluate the oxidation resistance performance, only the structure and inerting factors are not enough. Fundamentally, the quantitative evaluation of the loss degree of the internal carbon matrix can truly manifest the oxygen blocking effect of the system. Therefore, this study calculates the carbon loss rate and oxygen permeability of the SHS and CP coating samples according to Eqs. (5) and (6), as shown in Fig. 13. The carbon loss rate of ZM20-ZM40 coatings prepared by SHS powders is 1.22×10 -6 , 0.87×10 -6 , and 0.29×10 -6 g·cm -2 ·s -1 , and the oxygen permeability is 1.09%, 0.55%, and 0.30%, respectively. It can be further found from Fig. 13 that the oxygen permeability and carbon loss rate of the CP coating samples are higher than the corresponding SHS coating samples. Furthermore, compared with SHS-ZM20 and SHS-ZM30, the carbon loss rate of SHS-ZM40 decreased by 76.2% and 66.7%, and the oxygen permeability decreased by 72.5% and 45.5%, respectively. This proves that the 60 vol%ZrB 2 -40 vol%MoSi 2 composite powders synthesized by SHS technique, after preparation of the coating, have decent resistance at 1973 K due to their lower oxidation activity and excellent structural stability.

Conclusions
SHS and SPS techniques were utilized to enhance oxygen blocking of ZrB 2 -MoSi 2 composite coatings under 1973 K aerobic environment. The ZrB 2 -40 vol%MoSi 2 coating sample prepared using SHS powders exhibits the best oxygen blocking ability at 1973 K. Compared with the CP coatings, the densities of the SHS coatings were significantly enhanced, whose density difference with the theoretical density was reduced from 1.54 to 1.09 g/cm 3 (ZM20), from 1.46 to 1.03 g/cm 3 (ZM30), and from 1.10 to 0.79 g/cm 3 (ZM40), respectively. Due to the increase of density, the oxidation activity of the SHS coating was effectively reduced, and the structural oxygen resistance of the coating in the active oxidation stage was enhanced. The oxygen partial pressure gradient inside the coating led to the appearance of the Si-depleted layer, which critically boosted the complexity of oxygen infiltration. The rise of MoSi 2 content promoted the dispersion of transition metal oxides in SiO 2 , which conduced to shape the homogenous compound glass layer, and enlarged the viscosity of the glass layer. Compared with SHS-ZM20 and SHS-ZM30 samples, the carbon loss rate of the SHS-ZM40 component was diminished by 76.2% and 66.7%, and the oxygen permeability was decreased by 72.5% and 45.5%, respectively.