Synthesis of In Situ ZrB2-SiC-ZrC Coating on ZrC-SiC Substrate by Reactive Plasma Spraying

In situ synthesis feasibility of ZrB2-SiC-ZrC composite coatings on ZrC ceramics by reactive plasma spraying (RPS) was investigated. To help to understand the phase evolution during plasma spraying process, reaction behavior in the ZrH2-Si-B4C system was explored carefully by differential scanning calorimetry. The results indicated that the phase transformation sequence in the ZrH2-Si-B4C system could be described as ZrH1.66, Zr3O, ZrC, ZrB2, Zr2Si, ZrSi, and SiC. The prior formation of ZrC was due to high diffusion rate of C atoms from B4C. ZrB2 was produced above 1100 °C. As the temperature increased, SiC were finally formed by the reaction of ZrC with ZrSi and B4C. The RPS composite coatings mainly consisted of ZrB2, SiC, and ZrC phases, except for a small fraction of ZrO2 phase. The microstructural characterization exhibited more dense melted splats, which appears to increase gradually with the increase in spraying currents and distances. The coatings had typical lamellar structure and adhered to the substrate well. The microhardness values were higher than 1000 HV1, but there were few variations with varying spraying currents and distances.


Introduction
Zirconium carbide (ZrC), which is known as a refractory and chemically stable compound, presents with a high melting point, high strength, high thermal conductivity, and chemical stability, and is a promising candidate for being a structural component used in next-generation rocket engines and hypersonic spacecrafts [1][2][3]. Unfortunately, poor high-temperature properties, such as thermal oxidation and ablation, prevent its use in many applications, including the aforementioned high-temperature structural applications. However, in most cases, the oxidation damage originates from the surface of the materials. Therefore, a coating is an effective solution to enhance the performance of ZrC substrate materials at extreme environments by surface modification. ZrB 2 , with addition of 20-30 vol% SiC, exhibits good properties at high temperatures [4][5][6], which make it good candidate among coating materials applied on C/C composites [7,8]. Luckily, ZrB 2 -SiC [9,10] has a similar coefficient of thermal expansion (CTE) as ZrC [11], and has good chemical compatibility with ZrC [12]. These properties show that ZrB 2 -SiC ceramics are a feasible option for a protective coating of ZrC.
Among many fabrication methods, RPS combines atmospheric plasma spraying (APS) with self-propagation high-temperature synthesis (SHS), in which the desired phases for the coating could be formed in situ by the SHS reaction during the plasma spraying process [13]. It is suitable for the deposition of high temperature and refractory coatings [14][15][16].
The main objective of the present research was the feasibility of in situ synthesis ZrB 2 -SiC-ZrC composite coatings by RPS without transitional adhesive layer on ZrC ceramics. The coating was obtained by using an SHS reaction in ZrH 2 -Si-B 4 C mixed powders during the plasma jet. Presently, reports on reactive plasma sprayed coatings for ZrC ceramics 2 of 12 substrate are relatively rare. Therefore, it is necessary to carry out basic work on the reaction behavior in the ZrH 2 -Si-B 4 C system and identify the fabrication characterizations of the coating. It is expected that the results obtained in this study can provide some valuable information for promoting the application of ZrC materials under ultra-hightemperature environments.

Experimental Methods
The raw powders used to prepare the ZrB 2 -SiC-ZrC coating were commercial powders of ZrH 2 (>99% in purity, ≤38 µm), Si (>99% in purity, ≤5 µm), and B 4 C (>99% in purity, ≤3.5 µm). The mole ratio of ZrH 2 -Si-B 4 C powders was designed according to Equation (1) [17,18] to achieve ZrB 2 -based ceramics containing 20 vol% SiC. xZrH 2 + yB 4 C + (3y − x)Si → 2yZrB 2 + (x − 2y)ZrC + (3y − x)SiC + xH 2 (1) The powder mixture was wet-mixed at 45 r/min for 36 h in a roller ball mill using ZrO 2 balls with the additions of deionized water, Gum Arabic and tri-ammonium citrate, where Gum Arabic is a binder for agglomerating powders, and tri-ammonium citrate is a defoaming agent used to remove the bubbles in the mixture. The obtained slurry was spraydried and then sieved to produce the agglomerated particles with 50-100 µm. The substrate samples of dimension 10 × 20 × 5 mm 3 were cut from the ZrC-SiC composite billet. ZrC-SiC composites were prepared by hot pressed sintering and the details can be found in our previous work [19]. Before spraying, the substrate samples were grit-blasted by alumina sand and ultra-sonic-cleaned with ethanol. SEM morphology of the substrate before and after grit-blasting is shown in Figure 1. The ZrB 2 -SiC-ZrC coatings on the surface of the ZrC-SiC composites were prepared by the spraying apparatus, as schematically shown in Figure 2. The composite, agglomerated powders were deposited onto the substrate by atmospheric plasma spraying using a Unicoat Spraying System (F4 spraying gun, Sulzer Metco, Switzerland) with Ar and H 2 as plasma gases, in which the flow rates of the gases were 35 L/min and 12 L/min, respectively. The flow rate of the carrier powder gas Ar was 2.5 L/min. Spray distances from the plasma gun nozzle to the substrate were selected as 80, 100, and 140 mm. These spraying parameters were based on previous work [20], and the detailed information is listed in Table 1. Each spraying experiment was performed about 3~5 times. Since ZrB 2 -SiC-ZrC coatings were first in situ synthesized by RPS on ZrC-SiC ceramic substrate, it is necessary to investigate the reaction process of the ZrH 2 -Si-B 4 C system. It should be noted that the reaction in the DSC apparatus was different from that in plasma spraying because of the disparity in the processing conditions, particularly for the heating rate and heat loss rate, but it was very similar, to some extent, as far as the reaction process. Therefore, the studies on DSC are valuable in understanding the reaction behavior and phase evolution during the plasma spraying process [21]. Thermal analysis of the composite agglomerated powders was carried out by an STA 449F3 Jupiter ® -type thermal analyzer, and were conducted in a flowing argon gas using a heating rate of 10 • C/min. The phase constituents of DSC products and as-prepared coating were identified by Xray diffraction (XRD, Philips X'-Pert PRO). The morphologies of DSC products, and the surfaces and cross-sections of the coatings, were characterized using a scanning electron microscope (SEM) (Model FEI Sirion-200, Holland) equipped with an energy-dispersive X-ray spectrometer (EDS). Furthermore, transmission electron microscopy (JEM-2100) was also employed to examine the surface microstructure of the coating. The hardness of the coatings was measured with a Vickers hardness tester (HV-1000A) at a load of 1000 g with a dwell time of 15 s (5 indents for each sample).

Results and Discussion
During the plasma spraying process, the size distribution and morphology of composite agglomerated powders must be suitable, such as the uniform sphericity a proper cohesion strength, so as to ensure complete reaction in the plasma spraying p cess and to ensure uniform coatings [6,22]. The morphology of agglomerated partic used for RPS is shown in Figure 3. Spray-dried particles were spherical and near spheric with sizes of 50-100 μm. The interior of the particles exhibited a relatively dense structu implying good contact among the reactants. Meanwhile, every composite agglomera particle can be recognized as an independent, SHS-reactive unit, which is beneficial to ignition of the SHS reaction during the RPS process [14]. Thus, the characteristics of composite powders endowed with better flowability of spray-dried particles and favor complete SHS reaction during the RPS process.

Results and Discussion
During the plasma spraying process, the size distribution and morphology of the composite agglomerated powders must be suitable, such as the uniform sphericity and proper cohesion strength, so as to ensure complete reaction in the plasma spraying process and to ensure uniform coatings [6,22]. The morphology of agglomerated particles used for RPS is shown in Figure 3. Spray-dried particles were spherical and near spherical, with sizes of 50-100 μm. The interior of the particles exhibited a relatively dense structure, implying good contact among the reactants. Meanwhile, every composite agglomerated particle can be recognized as an independent, SHS-reactive unit, which is beneficial to the ignition of the SHS reaction during the RPS process [14]. Thus, the characteristics of the composite powders endowed with better flowability of spray-dried particles and favored complete SHS reaction during the RPS process.

Results and Discussion
During the plasma spraying process, the size distribution and morphology of the composite agglomerated powders must be suitable, such as the uniform sphericity and proper cohesion strength, so as to ensure complete reaction in the plasma spraying process and to ensure uniform coatings [6,22]. The morphology of agglomerated particles used for RPS is shown in Figure 3. Spray-dried particles were spherical and near spherical, with sizes of 50-100 µm. The interior of the particles exhibited a relatively dense structure, implying good contact among the reactants. Meanwhile, every composite agglomerated particle can be recognized as an independent, SHS-reactive unit, which is beneficial to the ignition of the SHS reaction during the RPS process [14]. Thus, the characteristics of the composite powders endowed with better flowability of spray-dried particles and favored complete SHS reaction during the RPS process.   Figure 5a,b, it can be found that only ZrH2 and Si phases were detected below 400 °C; however, a large quantity of Zr3O and a trace of ZrH1.66 appeared, besides unreacted starting Si phase, when quenched at 700 °C , indicating that ZrH2 mainly decomposed between 400 °C and 700 °C. Additionally, the appearance of ZrH1.66 peaks implies that the dehydrogenation of ZrH2 was a gradual process with the sequence of ZrH2 and ZrHx, which is consistent with the result reported in Reference [17]. In this case, in the formation of the Zr3O phase, it can be presumed that ZrH2 powders or dehydrogenation products ZrHx and Zr reacted with elemental O, which was derived from Gum Arabic and tri-ammonium citrate in the powder system. However, the provided amount of O in the system was not sufficient and the dehydrogenation of ZrH2 released hydrogen, which means that the system is in a reducing atmosphere, rich in H2, for a short time, thus favoring the formation of Zr3O instead of stable ZrO2. A strong exothermic peak appeared at about 819 °C. The corresponding XRD results heated to 800 °C in Figure 5c show that lots of Zr3O and unreacted Si were only detected in the final product. Therefore, it can be predicted that the conversion reaction of ZrHx or Zr to Zr3O predominated at about 800 °C followed by continuous dehydrogenation of ZrHx. After a strong exothermic peak of 819 °C, an endothermic peak with the minimum at about 972 °C was observed in the DSC curve. The XRD result of the product heated to 960 °C in Figure 5d show that a trace of ZrC was found in addition to Zr3O and Si phases. Some researchers have concluded that B4C is primarily composed of B11C icosahedra, rich in boron, linked by C-B-C intericosahedral chains [23,24]. At high temperatures, carbon atoms are more easily dissociated compared with boron atoms; namely, the diffusivity of carbon from B4C crystal also is higher than that of boron [25]. In addition, because Zr is a more reactive metal than Si, any free Zr will preferentially react with C species to form Zr compounds and leave unreacted Si. Therefore, it can be concluded that C dissociated from B4C rapidly diffused into Zr3O, and reduced Zr3O to form ZrC, forming B-rich boron carbide at about 960 °C. In the temperature range of 1000 °C-1200 °C, several exothermic peaks were very close, which may be caused by continuous occurrence of several exothermic reactions. The corresponding quenched product, as shown in Figure 5e, suggests that the phase type at 1100 °C was similar to that at 960 °C, but the content of Zr3O significantly decreased and ZrC increased instead, illustrating that the transformation from Zr3O to ZrC was continuing. With quenching temperature rising above 1200 °C, two new phases of ZrB2 and Zr2Si were detected with the disappearance of Si phase, as shown in Figure 5f. Wu et al. [26] suggests that B and C atoms from B4C diffuse faster than Zr and Si. Additionally, Si has a lower reactivity than Zr, so ZrB2 and ZrC preferentially formed by reaction between Zr3O and  Figure 5a,b, it can be found that only ZrH 2 and Si phases were detected below 400 • C; however, a large quantity of Zr 3 O and a trace of ZrH 1.66 appeared, besides unreacted starting Si phase, when quenched at 700 • C, indicating that ZrH 2 mainly decomposed between 400 • C and 700 • C. Additionally, the appearance of ZrH 1.66 peaks implies that the dehydrogenation of ZrH 2 was a gradual process with the sequence of ZrH 2 and ZrH x , which is consistent with the result reported in Reference [17]. In this case, in the formation of the Zr 3 O phase, it can be presumed that ZrH 2 powders or dehydrogenation products ZrH x and Zr reacted with elemental O, which was derived from Gum Arabic and tri-ammonium citrate in the powder system. However, the provided amount of O in the system was not sufficient and the dehydrogenation of ZrH 2 released hydrogen, which means that the system is in a reducing atmosphere, rich in H 2 , for a short time, thus favoring the formation of Zr 3 O instead of stable ZrO 2 . A strong exothermic peak appeared at about 819 • C. The corresponding XRD results heated to 800 • C in Figure 5c show that lots of Zr 3 O and unreacted Si were only detected in the final product. Therefore, it can be predicted that the conversion reaction of ZrH x or Zr to Zr 3 O predominated at about 800 • C followed by continuous dehydrogenation of ZrH x . After a strong exothermic peak of 819 • C, an endothermic peak with the minimum at about 972 • C was observed in the DSC curve. The XRD result of the product heated to 960 • C in Figure 5d show that a trace of ZrC was found in addition to Zr 3 O and Si phases. Some researchers have concluded that B 4 C is primarily composed of B 11 C icosahedra, rich in boron, linked by C-B-C intericosahedral chains [23,24]. At high temperatures, carbon atoms are more easily dissociated compared with boron atoms; namely, the diffusivity of carbon from B 4 C crystal also is higher than that of boron [25]. In addition, because Zr is a more reactive metal than Si, any free Zr will preferentially react with C species to form Zr compounds and leave unreacted Si. Therefore, it can be concluded that C dissociated from B 4 C rapidly diffused into Zr 3 O, and reduced Zr 3 O to form ZrC, forming B-rich boron carbide at about 960 • C. In the temperature range of 1000 • C-1200 • C, several exothermic peaks were very close, which may be caused by continuous occurrence of several exothermic reactions. The corresponding quenched product, as shown in Figure 5e, suggests that the phase type at 1100 • C was similar to that at 960 • C, but the content of Zr 3 O significantly decreased and ZrC increased instead, illustrating that the transformation from Zr 3 O to ZrC was continuing. With quenching temperature rising above 1200 • C, two new phases of ZrB 2 and Zr 2 Si were detected with the disappearance of Si phase, as shown in Figure 5f. Wu et al. [26] suggests that B and C atoms from B 4 C diffuse faster than Zr and Si. Additionally, Si has a lower reactivity than Zr, so ZrB 2 and ZrC preferentially formed by reaction between Zr 3 O and B-rich or residual B 4 C. However, at this temperature, Si was also replaced by Si compounds. Similar to the Hf-Si system [27], the predominance of specific zirconium silicides at different temperatures logically depends on the activity ratio a Zr /a Si . At a high temperature of 1200 • C (compared with 1100 • C), Zr 2 Si may form as a result of a diffusion reaction between Zr 3 O and Si. However, at 1300 • C, Zr 2 Si transformed into ZrSi and disappeared at higher temperature of 1500 • C, which indicates that Zr x Si y was unstable in the presence of B 4 C and Zr at high temperatures (above 1300 • C). In addition, the peak of Zr 3 O completely disappeared, and ZrB 2 became the main phase at temperatures from 1300 • C to 1500 • C. Additionally, SiC appeared with a disappearance of Zr x Si y peaks and a decrease in the intensities of ZrC peaks, based on the XRD diffractograms. This phenomenon supports the conclusion that the formation of SiC derived from the reaction between ZrC, ZrSi, and residual B 4 C in this temperature region.
Materials 2022, 15, x FOR PEER REVIEW B-rich or residual B4C. However, at this temperature, Si was also replaced by pounds. Similar to the Hf-Si system [27], the predominance of specific zirconium s at different temperatures logically depends on the activity ratio aZr/aSi. At a high ature of 1200 °C (compared with 1100 °C), Zr2Si may form as a result of a diffusion r between Zr3O and Si. However, at 1300 °C, Zr2Si transformed into ZrSi and disap at higher temperature of 1500 °C, which indicates that ZrxSiy was unstable in the p of B4C and Zr at high temperatures (above 1300 °C). In addition, the peak of Zr3 pletely disappeared, and ZrB2 became the main phase at temperatures from 130 1500 °C. Additionally, SiC appeared with a disappearance of ZrxSiy peaks and a d in the intensities of ZrC peaks, based on the XRD diffractograms. This phenomen ports the conclusion that the formation of SiC derived from the reaction betwe ZrSi, and residual B4C in this temperature region.   B-rich or residual B4C. However, at this temperature, Si was also replaced by Si compounds. Similar to the Hf-Si system [27], the predominance of specific zirconium silicides at different temperatures logically depends on the activity ratio aZr/aSi. At a high temperature of 1200 °C (compared with 1100 °C), Zr2Si may form as a result of a diffusion reaction between Zr3O and Si. However, at 1300 °C, Zr2Si transformed into ZrSi and disappeared at higher temperature of 1500 °C, which indicates that ZrxSiy was unstable in the presence of B4C and Zr at high temperatures (above 1300 °C). In addition, the peak of Zr3O completely disappeared, and ZrB2 became the main phase at temperatures from 1300 °C to 1500 °C. Additionally, SiC appeared with a disappearance of ZrxSiy peaks and a decrease in the intensities of ZrC peaks, based on the XRD diffractograms. This phenomenon supports the conclusion that the formation of SiC derived from the reaction between ZrC, ZrSi, and residual B4C in this temperature region.  The typical SEM images of the raw powder and the products at various temperatures of 800 °C, 960 °C, 1200 °C, 1300 °C, and 1500 °C are shown in Figure 6a-f, respectively. As The typical SEM images of the raw powder and the products at various temperatures of 800 • C, 960 • C, 1200 • C, 1300 • C, and 1500 • C are shown in Figure 6a-f, respectively. As indicated, the surface of agglomerated particles quenched at 800 • C seemed to be denser and finer compared with that of spray-dried particles. The reasons may be one of the following: one is that the particle size of ZrH 2 was further decreased due to its decomposition into ZrH 1.66 above 700 • C (as shown in Figure 5 b,c); the other is that, during ball mixing, some Si particles were in high-energy state and had a higher reactivity, except for the size reduction, which causes the small-sized Si with the low melting point to melt slightly first, and fill the pores between the particles. Figure 6c shows the internal microstructures of agglomerated particles at 960 • C; it can be observed that the components kept the original irregular block shape. The EDS embedded in Figure 6c revealed that the surface of B 4 C was rich in a considerable amount of C (80.71 at.%, which is higher than 32.42 at.% of C in raw B 4 C particle), which agrees with the results in Reference [28]. The study suggested that the carbon photoabsorption spectrum of B 4 C changed when the temperatures were higher than 1000 K. At this moment, the surface of B 4 C was covered with a graphite layer, leaving B-rich boron carbide buried beneath the graphite surface. The above findings and suggestions support the results obtained in Figure 6c of this study, e.g., the rapid diffusion of carbon from B 4 C and the prior formation of ZrC. Figure 6d shows typical microstructure of the products at 1200 • C. As shown, with the temperature increasing, many scaly ZrB 2 particles were found on the surface of light gray compound particles containing Zr. Additionally, the surface of dark gray B 4 C particles became rough, suggesting that B and C atoms form B 4 C diffused out to react and were consumed. Figure 6e,f show typical microstructures of the products at 1300 • C and 1500 • C, respectively. Clearly, some partial or complete melts were more visible. Additionally, a large amount of of ZrB 2 particles were observed in the products, especially at 1500 • C, ZrB 2 particles grew into obvious hexagonal shape. Additionally, large grey bulk phase in Figure 6f was identified as ZrSi, and ZrB 2 was traced to grow from ZrSi particles (as shown in the image embedded in Figure 6f). All microstructures observations are consistent with the XRD identification results. From the above results, the possible reaction process of system during plasma spraying could be inferred, which are shown in Equations (2)-(8), as follows: The information obtained from the DSC experiments is not only helpful for the understanding of reaction process during plasma spraying, but also provides a fact that the objective product of ZrB 2 -SiC-ZrC can be synthesized by the ZrH 2 -Si-B 4 C system. The X-ray diffraction patterns of the coatings prepared by RPS using ZrH 2 -Si-B 4 C system are shown in Figure 7. Similar to the DSC results, ZrB 2 , SiC, and ZrC phases existed in the coating, while a small amount of ZrO 2 remained, and no other unreacted or intermediate phases were found. It is considered that the reactions in RPS were almost complete according to Equation (1). In this case, the formation of ZrO 2 was inevitable, since the RPS process was performed in air, and the air was also involved under the disturbing effect of the plasma jet on its surrounding air. Consequently, oxidation behavior occurred through the chemical reactions between the liquid droplet and oxygen during the coating process or after deposition onto the substrate. Besides, the change of ZrO 2 peak intensity with increasing the spray distance was more noticeable than that with increasing arc current, which may be attributed to the longer dwell time in the plasma jet. In addition, the relative intensity of SiC was low in the coating. On the basis of previous works [4,29], SiC can decompose during spraying, which can be reduced by a eutectic phase formed between ZrB 2 and SiC. In the present experimental work, the synthesis of ZrB 2 -SiC-ZrC coating by ZrH 2 -Si-B 4 C powder was carried out under the combined effects of plasma jet and reaction heat release of ZrH 2 -Si-B 4 C powder, where the higher system temperature may have facilitated the decomposition of SiC. Additionally, a possible formation of a ZrB 2 -SiC mixture was beneficial to prevent the decomposition. However, it should be noted that eutectic-like regions have not been found, this hypothesis still needs to be verified. It is also noticeable from Figure 7 that the diffraction peaks of the coating were low and broad, which means the phases in the coating had lower crystallinity and smaller grain size due to extremely high cooling rate. This presumption is in accordance with their high magnification SEM and TEM pictures, which are displayed in Figure 8. The SEM image (Figure 8a) reveals that the splats in the reactive plasma sprayed ZrB 2 -SiC-ZrC coating seems to be amorphous melt, and TEM image (Figure 8b) indicates that the coatings were composed of a large number of nanosized grains. The information obtained from the DSC experiments is not only helpful for the understanding of reaction process during plasma spraying, but also provides a fact that the objective product of ZrB2-SiC-ZrC can be synthesized by the ZrH2-Si-B4C system. The X- tic-like regions have not been found, this hypothesis still needs to be verified. It is also noticeable from Figure 7 that the diffraction peaks of the coating were low and broad, which means the phases in the coating had lower crystallinity and smaller grain size due to extremely high cooling rate. This presumption is in accordance with their high magnification SEM and TEM pictures, which are displayed in Figure 8. The SEM image ( Figure  8a) reveals that the splats in the reactive plasma sprayed ZrB2-SiC-ZrC coating seems to be amorphous melt, and TEM image (Figure 8b) indicates that the coatings were composed of a large number of nanosized grains.   Figure 9 is the surface SEM images of the coating. The desirable surface morphology of the ZrB2-SiC-ZrC coating, which has no obvious macro defects or cracks, was accomplished by utilizing a similar CTE principle between the substrate and the coating to minimize the effect of the mismatched CTE. The microstructural characterization of all of the coatings demonstrates that, except for a few large stacking regions which are composed of insufficiently molten powders, more dense, disk-like or pancake splats consisting of fully molten regions were found in the sprayed coating, and formed a flattened, smooth area. This splat pattern gave rise to coatings endowed with higher adhesive and cohesive noticeable from Figure 7 that the diffraction peaks of the coating were low and broad which means the phases in the coating had lower crystallinity and smaller grain size due to extremely high cooling rate. This presumption is in accordance with their high magnification SEM and TEM pictures, which are displayed in Figure 8. The SEM image ( Figure  8a) reveals that the splats in the reactive plasma sprayed ZrB2-SiC-ZrC coating seems to be amorphous melt, and TEM image (Figure 8b) indicates that the coatings were com posed of a large number of nanosized grains.   Figure 9 is the surface SEM images of the coating. The desirable surface morphology of the ZrB2-SiC-ZrC coating, which has no obvious macro defects or cracks, was accom plished by utilizing a similar CTE principle between the substrate and the coating to min imize the effect of the mismatched CTE. The microstructural characterization of all of the coatings demonstrates that, except for a few large stacking regions which are composed of insufficiently molten powders, more dense, disk-like or pancake splats consisting o fully molten regions were found in the sprayed coating, and formed a flattened, smooth area. This splat pattern gave rise to coatings endowed with higher adhesive and cohesive  Figure 9 is the surface SEM images of the coating. The desirable surface morphology of the ZrB 2 -SiC-ZrC coating, which has no obvious macro defects or cracks, was accomplished by utilizing a similar CTE principle between the substrate and the coating to minimize the effect of the mismatched CTE. The microstructural characterization of all of the coatings demonstrates that, except for a few large stacking regions which are composed of insufficiently molten powders, more dense, disk-like or pancake splats consisting of fully molten regions were found in the sprayed coating, and formed a flattened, smooth area. This splat pattern gave rise to coatings endowed with higher adhesive and cohesive strength [6,30]. It should be noted that the splats increase gradually with the increase in spraying currents and distances. This might be related with an increase in particle temperature, since an increase in spraying current increases the heat energy of the plasma jet and an increase in spraying distances causes the longer dwell time in plasma jet [31]. Additionally, a few splash structures were also observed, and the morphology of solidifying droplets was formed, as depicted in Figure 9d. This splash structure resulted from the fact that in-flight particles, on impact with the substrate, fragmented into smaller multiple droplets, and then solidified rapidly. Additionally, some microcracks and pores were distributed on surface of the coating. The formation of the microcracks may be attributed to be the shrinkage stress caused by rapid cooling of the coating. Incomplete overlapping between the splats might result in the appearance of some pores.
ditionally, a few splash structures were also observed, and the morphology of solidifying droplets was formed, as depicted in Figure 9d. This splash structure resulted from the fact that in-flight particles, on impact with the substrate, fragmented into smaller multiple droplets, and then solidified rapidly. Additionally, some microcracks and pores were distributed on surface of the coating. The formation of the microcracks may be attributed to be the shrinkage stress caused by rapid cooling of the coating. Incomplete overlapping between the splats might result in the appearance of some pores.  Figure 10 shows the cross-section of the coating samples observed by SEM. It can be seen that the coating shows uniform thickness (in the range of 250-300 m) and good interfacial adhesion. It can be seen that internal microstructure of the coating consists of two different regions which are easily distinguishable: loose and high pore region and compact lamellar region. In high magnification of Figure 10d, the intimate lamella is more noticeable; however, a small number of micropores are also clearly visible. The reason for the formation of the micropores may be explained: the molten droplets were deposited onto the substrate and splashed to form the shadowing effect, resulting in the enwrapped gas not being expelled promptly.  Figure 10 shows the cross-section of the coating samples observed by SEM. It can be seen that the coating shows uniform thickness (in the range of 250-300 µm) and good interfacial adhesion. It can be seen that internal microstructure of the coating consists of two different regions which are easily distinguishable: loose and high pore region and compact lamellar region. In high magnification of Figure 10d, the intimate lamella is more noticeable; however, a small number of micropores are also clearly visible. The reason for the formation of the micropores may be explained: the molten droplets were deposited onto the substrate and splashed to form the shadowing effect, resulting in the enwrapped gas not being expelled promptly.
As is well known, RPS is a process that utilizes the heat of high-temperature plasma to initiate combustion synthesis reactions [32]. In this process, once the exothermic reactions are ignited, self-propagating proceeds quickly, and the temperature of the system increases significantly in a short time. The powders, entering the plasma jet, will form a molten state, will convert the reactant particles into final products during in-flight processes, and then will impinge on the substrates as splats. With time, the continuous deposition of splats results in the formation of dense lamella structures. However, in fact, the powders appear to be emanative after they are ejected from the plasma nozzle, and thus some of the powders may be scattered outside the plasma jet and will form a semi-molten state to complete the synthesis reaction during spraying, and they subsequently impact on the substrates as a loose, stacked structure. The phenomenon will be strengthened if large-sized powder particles are involved, outside of the plasma jet. The microstructure characteristic in Figure 10  As is well known, RPS is a process that utilizes the heat of high-temperature plasma to initiate combustion synthesis reactions [32]. In this process, once the exothermic reactions are ignited, self-propagating proceeds quickly, and the temperature of the system increases significantly in a short time. The powders, entering the plasma jet, will form a molten state, will convert the reactant particles into final products during in-flight processes, and then will impinge on the substrates as splats. With time, the continuous deposition of splats results in the formation of dense lamella structures. However, in fact, the powders appear to be emanative after they are ejected from the plasma nozzle, and thus some of the powders may be scattered outside the plasma jet and will form a semi-molten state to complete the synthesis reaction during spraying, and they subsequently impact on the substrates as a loose, stacked structure. The phenomenon will be strengthened if large-sized powder particles are involved, outside of the plasma jet. The microstructure characteristic in Figure 10 supports the results of the surface observation. Figure 11 illustrates the microhardness results of ZrB2-SiC-ZrC coatings. All the microhardness values of the coatings are higher than 1000 HV1, although no obvious variation with varying spraying currents and distances can be observed, and are statistically identical to each other, considering the standard deviation of the data. The attainment of high hardness is presumably due to the more compact structure of the coating deposited by RPS and the high hardness of the hard phases, such as ZrB2, SiC, and ZrC.  Figure 11 illustrates the microhardness results of ZrB 2 -SiC-ZrC coatings. All the microhardness values of the coatings are higher than 1000 HV 1 , although no obvious variation with varying spraying currents and distances can be observed, and are statistically identical to each other, considering the standard deviation of the data. The attainment of high hardness is presumably due to the more compact structure of the coating deposited by RPS and the high hardness of the hard phases, such as ZrB 2 , SiC, and ZrC.

Conclusions
Dense and well-bonded ZrB2-SiC-ZrC composite coatings were successfully p in situ by reactive plasma sprayed ZrH2-Si-B4C powders. Reaction behavior in th Si-B4C system was investigated through phase analysis and microstructures of qu

Conclusions
Dense and well-bonded ZrB 2-SiC-ZrC composite coatings were successfully prepared in situ by reactive plasma sprayed ZrH 2 -Si-B 4 C powders. Reaction behavior in the ZrH 2 -Si-B 4 C system was investigated through phase analysis and microstructures of quenched samples by differential scanning calorimetry. The phase transformation sequence in the ZrH 2 -Si-B 4 C system could be described as ZrH 1.66 and Zr 3 O at 700 • C, ZrC at 960 • C, ZrB 2 and Zr 2 Si at 1200 • C, ZrSi at 1300 • C, and SiC at 1500 • C. The prior formation of ZrC over ZrB 2 was due to the fact that the diffusion rate of C atoms was higher than that of B atoms in B 4 C. As the temperature increased, SiC were finally formed by the reaction of ZrC with ZrSi and B 4 C. The phase composition of RPS coating consisted of ZrB 2 , SiC, ZrC, and a small quantity of ZrO 2 . Microstructure characterization exhibited more dense disk-like or pancake splats, except for a few large stacking structures. By utilizing a similar CTE principle between the substrate and the coating, the absence of macro defects and cracks, and good interfacial bonding, were obtained. In contrast, the microhardness value of the coating exhibited few variations with various spraying currents and distances. In this work, all obtained results indicated that it was feasible to fabricate ZrB 2 -SiC-ZrC ceramics coatings, without an adhesive layer, on ZrC ceramic substrate by the RPS method. For hightemperature application coatings, research on the mechanical behaviors and the failure mechanisms of ZrB 2-SiC-ZrC coatings in high-temperature environments must urgently be carried out in future work.