Abstract: The in situ nano Ta4HfC5 reinforced SiBCN–Ta4HfC5 composite ceramics were prepared by a combination of two-step mechanical alloying and reactive hot-pressing sintering. The microstructural evolution and mechanical properties of the resulting SiBCN–Ta4HfC5 were studied. After the first-step milling of 30 h, the raw materials of TaC and HfC underwent crushing, cold sintering, and short-range interdiffusion to finally obtain the high pure nano Ta4HfC5. A hybrid structure of amorphous SiBCN and nano Ta4HfC5 was obtained by adopting a second-step ball-milling. After reactive hot-pressing sintering, amorphous SiBCN has crystallized to 3C-SiC, 6H-SiC, and turbostratic BN(C) phases and Ta4HfC5 retained the form of the nanostructure. With the in situ generations of 2.5 wt% Ta4HfC5, Ta4HfC5 is preferentially distributed within the turbostratic BN(C); however, as Ta4HfC5 content further raised to 10 wt%, it mainly distributed in the grain-boundary of BN(C) and SiC. The introduction of Ta4HfC5 nanocrystals can effectively improve the flexural strength and fracture toughness of SiBCN ceramics, reaching to 344.1 MPa and 4.52 MPa·m1/2, respectively. This work has solved the problems of uneven distribution of ultra-high temperature phases in the ceramic matrix, which is beneficial to the real applications of SiBCN ceramics.
Keywords: Ta4HfC5; SiBCN; microstructure evolution; mechanical properties
With the development of aerospace technology, more stringent property requirements are put forward for advanced structural-functional integration of ceramic materials. Among many high-temperature structural materials, SiBCN non-oxide ceramics have attracted considerable attention due to their light-weight, high specific strength, excellent thermal stability, and resistance to thermal shock, oxidation, and ablation [1–3].
Polymer/precursor derived ceramics (PDCs), one of the very first routes to prepare SiBCN materials, remain amorphous nature at least up to 1400 ℃, and do not undergo microstructural changes at this temperature . The chemistry, microstructure, and properties of the PDCs SiBCN can be tailored effectively by controlling the initial reagent structure, chemical reaction, and processing parameters . However, this method also has some nonnegligible shortcomings that limit its wide applications. The starting materials (such as SiCl3CH3, and BCl3) are usually expensive and some organic solvents used are harmful to human health and the environment . During pyrolysis, the gas release and mass loss require careful monitoring to avoid porous structure and to carefully control shrinkage (even microcracks) [5–7].
Mechanical alloying and solid-state sintering techniques (also referred to as inorganic method) to obtain dense nano or amorphous SiBCN monoliths and their ceramic composites were proposed . One of the significant advantages of this processing route is to provide centimeter-sized samples convenient for evaluating the basic mechanical and thermophysical properties [7,8]. However, the previous studies showed that the heterogeneous microstructure leading to poor mechanical properties of the SiBCN ceramics are needed to be further optimized for the real applications at high temperatures [9,10].
The addition of ultra-high temperature ceramics (UHTCs) into the SiBCN ceramic matrix is believed to be an effective strategy to improve the mechanical performance and high temperature properties of SiBCN ceramics [11–14]. In previous attempts, the microstructural evolution and thermal stability of the porous PDCs ZrB2–SiBCN and HfN–SiBCN composite ceramics have been investigated [15,16]. However, due to the limitation of sample size, the mechanical properties of these composite ceramics have not been evaluated. In another contribution, Liang et al. [17–19] have introduced some LaB6 into the SiBCN matrix by the inorganic method, obtaining desirable mechanical properties compared with pure SiBCN. However, during the sintering process, some LaB6 has reacted with SiBCN ceramic matrix to generate La2B2C6, and consequently, the strengthening and toughening effects of LaB6 were significantly reduced. Further, Miao et al. [20–21] have prepared SiBCN–ZrB2 composite ceramics by mechanical alloying combined with the sol-gel method. The results suggested that the in situ ZrB2 can improve the mechanical properties of SiBCN ceramic materials to some extent; however, the particle size of the in situ ZrB2 has grown up to a maximum value of ~1 μm.
Generally, the following factors should be considered when the external UHTCs are selected as reinforcements in SiBCN ceramics: (i) They should not react with the SiBCN matrix to form strong bonding interfacial structure during annealing or sintering; (ii) they could be uniformly dispersed in the ceramic matrix in the form of nanocrystals; and (iii) they are supposed to have good physical-chemical compatibility with SiBCN.
The chemically stable Ta4HfC5, the current highest melting point compound, was theoretically forecasted by Agte and Alterthum.  and the experiment confirmed by Andrievskii et al. . Up to the present, various strategies are used to the synthesis of the high pure Ta4HfC5. For example, Simonenko et al.  have reported a sol-gel technology for preparing Ta4HfC5 using metal oxide and gel of polymeric carbon source. The polymeric carbon source undergoes preliminarily annealing to provide a uniform dispersion of metal oxide–carbon mixture, and then further heating to 1200–1500 ℃ to obtain pure Ta4HfC5. Lu et al.  used the Hf and Ta containing polytan-tahafnoxane modified with allyl-functional novolac resin to produce Ta4HfC5 polymer (polytantahafno-xanesal). The resulting polytantahafnoxanesal is subsequently solidified and pyrolyzed at 1400 ℃ to obtain pure Ta4HfC5. These methods are very effective to prepare pure Ta4HfC5 with high thermal stability and chemical stability; however, these processing methods often contain complex synthesis steps leading to high-cost products. The raw materials used are also flammable and present waste disposal problems. For this reason, Gaballa et al.  have successfully prepared nano Ta4HfC5 by mechanical alloying for the first time. However, the formation mechanisms of the mechanical alloying derived nano Ta4HfC5 is still unclear,
Thus, in this work, initially, the Ta4HfC5 nanocrystals were prepared by the mechanical alloying method and the formation mechanism of Ta4HfC5 during mechanical alloying was analyzed. Then, SiBCN–Ta4HfC5 amorphous-nanocrystalline composite ceramic powder was prepared by the second step of mechanical alloying. Afterward, the hot-pressing sintering technology was adopted to consolidate SiBCN–Ta4HfC5 composite ceramics. As expected, the homogeneous distribution of nano Ta4HfC5 within the SiBCN ceramic matrix strongly improved the mechanical properties of the composite ceramics. The formation mechanisms of mechanical allying derived nano Ta4HfC5, and the correlation of microstructural evolution and mechanical properties of the resulting SiBCN–Ta4HfC5 composite ceramics were illustrated.
2 Materials and methods
2. 1 Raw materials
Hexagonal boron nitride (h-BN, 99.0% purity, 0.6 μm, purchased from Advanced Technology &Materials Co., Ltd.), graphite (99.5% purity, 8.7 μm, purchased from Qingdao Huatai Lubricant Sealing S&T Co., Ltd.), cubic silicon (c-Si, 99.9% purity, 9.0 μm, purchased from China New Metal Materials Technology Co., Ltd.), hafnium carbide (HfC, 99.0% purity, 1 μm, purchased from Shanghai Puwei Applied Materials Technology Co., Ltd.), and tantalum carbide (TaC, 99.0% purity, 1 μm, purchased from Shanghai Puwei Applied Materials Technology Co., Ltd.) were used here as received.
2. 2 Synthesis of Ta4HfC5 nanocrystals
Ta4HfC5 nanocrystals were prepared by the mechanical alloying method via a P4 high-energy ball-miller manufactured by FRITSCH company. Firstly, TaC and HfC mixed powders with a molar ratio of 4:1 were poured into the ball milling tanks filled with argon and then subjected to different time of milling. The ball-to-powder mass ratio was set as 20:1 and the effective ball-milling time was x h (x = 0.5, 1, 1.5, 2, 3, 5, 10, 20, 30). The main disc was rotated at 350 rpm and the vials were rotated at 800 rpm in reverse.
2. 3 Synthesis of SiBCN–Ta4HfC5 powder and bulk ceramics
In this stage, c-Si, h-BN, graphite, and Ta4HfC5 nanocrystals were put into ball milling tanks. Among them, the molar ratio of Si:BN:C is 2:1:3, and Ta4HfC5 nanocrystals account for 0 wt%, 2.5 wt%, 5 wt%, 10 wt%, and 15 wt% of the total mass of powder in the ball milling tank, respectively. During the milling process, the rotation speed of the main disc was set at 350 rpm and the vials were rotated at 600 rpm in reverse. In this scene, the effective milling time of 20 h could ensure a well-defined microstructure of the mixtures [28,29]. The powders mixed, stored, and transported were conducted in high pure Ar.
SiBCN–Ta4HfC5 mixed powders were loaded into a graphite mould with a diameter of 36 mm and then sintered via the hot-pressing sintering system (50-250T model, AVS company, USA). The sintering was conducted at a temperature of 1900 ℃ and a holding time for 1 h under an axial pressure of 60 MPa. The content of Ta4HfC5 in SiBCN–Ta4HfC5 composite ceramics is 0 wt%, 2.5 wt%, 5 wt%, 10 wt%, and 15 wt%.
2. 4 Characterization
The phase composition of the samples was measured by X'PERT X-ray diffractometer purchased from the Panalytical Company of Netherlands. The scanning speed is 10 (°)/min and the scanning range is 10°–90°. The morphologies of the samples were observed using the NanoLab 600i scanning electron microscope (SEM) produced by the FEI company. Tecnai G2 F20 transmission electron microscope was used to observe the microstructure of Ta4HfC5 and SiBCN–Ta4HfC5 powders. The microstructure and element distribution of the SiBCN–Ta4HfC5 composite ceramics were investigated using a Talos f200x transmission electron microscope (TEM) produced by the FEI company. The Raman spectra were collected via an inVia-Reflex testing system manufactured by RENISHAW company, with the excitation wavelength of 785 nm.
3 Results and discussion
3. 1 Synthesis and characterization of nano Ta4HfC5 powder
Before mechanical alloying, the powder mixture of the 4TaC–HfC shows sharp diffraction peaks (Fig. 1). After 0.5 h of milling, the intensities of the corresponding diffraction peaks are reduced and the full width at half maximum (FWHM) becomes wider. This should be argued to the grain refinement of TaC and HfC . As the milling time further raises to 1–5 h, TaC and HfC gradually undergo a solid solution reaction to form Ta4HfC5. Interestingly, the single-phase Ta4HfC5 solid solution is obtained after mechanical alloying for 10 h. Further prolonging the milling time to 20–30 h, the diffraction peaks of the as-prepared Ta4HfC5 gradually decrease, implying the grain refinement of Ta4HfC5.
Fig. 1 XRD patterns of the 4TaC–HfC powder mixture after mechanical alloying for different time.
After mechanical alloying, the diffraction peaks of the as-prepared Ta4HfC5 shifting to the small angle direction result from the radius difference between Hf and Ta, which can be depicted by the Vegard's law:
a＝ a1c1+ a2c2 (1)
where a, a1, and a2 represent the lattice constant of the new forming solid solution and the two starting components, respectively. The c1 and c2 represent the concentration of the two starting components, respectively [31–33]. The radius of the Hf is bigger than that of Ta. Thus, the short-range diffusion of Hf atoms into the interstitial sites of Ta can lead to the expansion of the TaC lattice (Fig. 2).
Fig. 2 XRD patterns of pure TaC powder and 4TaC–HfC powder mixture after mechanical alloying for 30 h, Inset is the magnified image of 65°–75°.
TEM analysis confirms the raw materials of TaC and HfC powders both have relatively large grain size (Figs. 3(a) and 3(b)). Besides, the corresponding SAED pattern displays some diffraction spots and rings of HfC and TaC (Fig. 3(c)). After ball milling for only 0.5 h, the grain size of TaC and HfC decreases obviously and some diffraction spots of TaC and HfC have disappeared (Figs. 3(d) and 3(f)). Notably, at this milling time, mechanical alloying has induced many structural defects in both TaC and HfC, which can provide conveniently channel for Hf atoms diffused inward. With progressing milling of 30 h, only pure Ta4HfC5 nanocrystals are obtained without any trace of HfC or below the limits of detection (Figs. 3(g)–3(i)). For pure TaC ball-milled for 30 h (Figs. 3(j)–3(l)), SAED pattern with bright-spots and narrow-rings verifies the well-defined crystalline phase of TaC. The interplanar crystal spacing of TaC is slightly smaller than that of Ta4HfC5, which is consistent with XRD analysis results.
Fig. 3 TEM, HRTEM, and SAED images of the pure TaC powder and 4TaC–HfC powder mixture after mechanical alloying for different time. (a–c) 4TaC–HfC, 0 h; (d–f) 4TaC–HfC, 0.5 h; (g–i) 4TaC–HfC, 30 h; (j–l) pure TaC powder, 30 h.
The SEM surface morphologies in Fig. 4(a) further ensure that the raw materials of the 4TaC–HfC powder mixture have a larger particle size. Nevertheless, after mechanical alloying for 0.5 h, the particle size of the powder mixture decreases obviously, which is mainly composed of nanoparticles (Fig. 4(b)). With the progress of mechanical alloying (1.5–30 h), these particles are continuously deformed, crushed, and cold sintered; however, the morphologies of the resulting particles are almost unchanged. The EDS maps in Figs. 4(f)–4(i) confirm that the Ta, Hf, and C atoms distribute uniformly after 30 h of milling, suggesting a good solid solubility of Hf into TaC.
Fig. 4 SEM surface images of 4TaC–HfC mixed powders after mechanical alloying for different time. (a) 0 h; (b) 0.5 h; (c) 1.5 h; (d) 20 h; (e, f) 30 h; (g–i) the corresponding element maps, and the inserted are the magnifying SEM images.
The formation mechanisms of the mechanical alloying derived nano Ta4HfC5 are elaborated in Fig. 5. During the mechanical alloying process, the sample particles are continuously deformed, crushed, and cold sintered under mechanical impact conditions, which led to grain refinement and microstrain occurring in the crystal grains. The dislocation density (ρd) can be represented by the following equation :
where ε is the microstrain, D is the grain size, and b is the Burgers vector. According to Eq. (2), larger microstrain and smaller grain sizes lead to higher dislocation density in the crystal. Besides, the decrease in grain size increases the surface area per unit volume. Higher surface energy provides a driving force for diffusion, and more defects provide channels for atom diffusion. Therefore, a uniform solid solution of Ta4HfC5 was formed.
Fig. 5 Formation mechanisms of the nano Ta4HfC5 during mechanical alloying.
Figure 6 shows the XRD patterns of the in situ SiBCN–Ta4HfC5 amorphous-nanocrystalline composite powder with different Ta4HfC5 contents. After 20 h of mechanical alloying, the lattice structure of c-Si, h-BN, and graphite was destroyed to form amorphous SiBCN. Besides, the broad diffraction peaks of Ta4HfC5 were observed, and these peak intensities gradually increase with the increase of Ta4HfC5 content. Confidentially, the nano Ta4HfC5 does not react with other components of SiBCN and still presents the form of nanocrystals (3–5 nm) in the amorphous matrix after mechanical alloying (Fig. 7).
Fig. 6 XRD patterns of in situ SiBCN–Ta4HfC5 amorphousnanocrystalline composite powder with different Ta4HfC5 contents.
Fig. 7 TEM and HRTEM images of the (a, b) SiBCN amorphous powder and (c, d) SiBCN–Ta4HfC5 amorphousnanocrystalline composite powder.
3. 2 Microstructural evolution and mechanical properties of as-sintered SiBCN–Ta4HfC5 composite ceramics
After reactive hot-pressing sintering, SiBCN–Ta4HfC5 composite ceramics mainly consist of BN(C), β-SiC, α-SiC, and Ta4HfC5 (Fig. 8). The intensity of the diffraction peaks of the Ta4HfC5 is positively correlated with the content of Ta4HfC5. With the increase of Ta4HfC5 content, the diffraction peaks of α-SiC decrease and the diffraction peaks of β-SiC increase. This means that the in situ Ta4HfC5 affects the β→α transition of SiC.
Fig. 8 XRD patterns of the as-sintered SiBCN–Ta4HfC5 composite ceramics with different Ta4HfC5 contents.
Figure 9 exhibits the Raman spectra of the as-sintered SiBCN–Ta4HfC5 composite ceramics with various Ta4HfC5 content. Two Raman peaks in a range of 60– 260 cm–1 are gradually enhanced with the increased content of the in situ Ta4HfC5. The emergence of Raman peak at 1374 cm–1 results from the D-side peak edge of graphite and the scattering peak of h-BN . However, the generation of Raman peak at 1588 cm–1 should be assigned to the G-side peak of graphite .
Fig. 9 Raman spectra of the as-sintered SiBCN–Ta4HfC5 composite ceramics with different Ta4HfC5 contents.
After hot-pressing sintering, amorphous SiBCN has crystallized and transformed into turbostratic BN(C) and nano SiC (Fig. 10). For pure SiBCN, BN(C) phase is distributed at the grain boundaries of SiC grains. Besides, the Si elements are mainly distributed in SiC grains, and B, C, and N elements are distributed in the form of the BN(C) phase at SiC grain boundaries. With a 2.5 wt% Ta4HfC5 addition, most of the Ta4HfC5 (~10 nm) are distributed in the BN(C) phase in the form of nanocrystals (Fig. 11). The Ta4HfC5 grains are thereby separated by BN(C). This structure is beneficial to the grain refinement of Ta4HfC5 and constrains short-range diffusion of the atoms. The EDS maps of the selective region in Ta4HfC5–BN(C) clearly show that the Ta, Hf, and some C elements are distributed in the Ta4HfC5 phase, and the rest of the C elements are distributed in BN(C) phase (Fig. 12). With 10 wt% Ta4HfC5 addition, a part of nano Ta4HfC5 (~10 nm) is still uniformly distributed in BN(C), and another nano Ta4HfC5 >10 nm is randomly distributed in the ceramic matrix (Fig. 13). The microstructural evolution diagram of the SiBCN–Ta4HfC5 composite ceramics with different Ta4HfC5 contents is displayed in Fig. 14.
Fig. 10 TEM and HEM images of the pure SiBCN ceramics. (a, b) Bright-field images; (c–h) HAADF–STEM image and corresponding EDS maps.
Fig. 11 TEM and HRTEM images of the SiBCN–Ta4HfC5 composite ceramics with 2.5 wt% Ta4HfC5. (a, b) Bright-field images; (c) magnifying TEM image of area A; (d) HRTEM image of area A.
Fig. 12 Selective EDS maps of Ta4HfC5 distributed within BN(C) region for composite ceramics with 2.5 wt% Ta4HfC5.
Fig. 13 TEM images and EDS maps of the SiBCN–Ta4HfC5 composite ceramics with 10 wt% Ta4HfC5. (a) Bright-field TEM image; (b) STEM image; (d–j) EDS maps.
Fig. 14 Schematic illustrating the microstructure evolution of Ta4HfC5 distribution in composite ceramics with different Ta4HfC5 addition.
The flexural strength and fracture toughness of the as-sintered SiBCN–Ta4HfC5 composite ceramics are shown in Fig. 15. The introduction of Ta4HfC5 nanocrystals has effectively improved the flexural strength and fracture toughness of the SiBCN ceramics. For pure SiBCN ceramics, they only show flexural strength of 156.1 MPa and the fracture toughness of 1.82 MPa·m1/2. With 10 wt% Ta4HfC5 addition, the composite ceramics possess optimized flexural strength reaching 344.1 MPa, and the composite ceramics with 5 wt% Ta4HfC5 obtain fracture toughness of 4.52 MPa·m1/2.
Fig. 15 Flexural strength and fracture toughness of the as-sintered SiBCN–Ta4HfC5 composite ceramics with different contents of Ta4HfC5.
In this study, SiBCN–Ta4HfC5 composite ceramics were prepared by two-step mechanical alloying combined with reactive hot-pressing sintering. After the above analysis, the following conclusions can be drawn:
1) In the initial step of mechanical alloying of 30 h, the 4TaC–HfC powder mixture is crushed, cold sintered, and interdiffused, finally forming Ta4HfC5 nanocrystalline. After the second step of milling for 20 h, a hybrid structure of amorphous SiBCN and nano Ta4HfC5 can be obtained. The lattice structures of C–Si, h-BN, and graphite were destroyed to form amorphous structures. However, Ta4HfC5 is uniformly distributed in amorphous powder in the form of nanocrystals.
2) After reactive hot-pressing sintering, SiBCN–Ta4HfC5 composite ceramics mainly contain Ta4HfC5, BN(C), β-SiC, and α-SiC. The Ta4HfC5 still exists in the form of nanocrystalline and does not react with the SiBCN matrix composition. With only 2.5 wt% Ta4HfC5 addition, nano Ta4HfC5 is preferentially distributed in BN(C) phase; however, it tends to both distribute in BN(C) phase and ceramic matrix when 10 wt% Ta4HfC5 is adopted.
3) The introduction of the Ta4HfC5 nanocrystals can effectively improve the flexural strength and fracture toughness of SiBCN ceramics due to the grain refinement and uniform distribution of nano Ta4HfC5. SiBCN ceramics with 10 wt% Ta4HfC5 present optimized flexural strength of 344.1 MPa, and composite ceramics with 5 wt% Ta4HfC5 obtain fracture toughness of 4.52 MPa·m1/2.