Microstructure and properties of nano-laminated Y3Si2C2 ceramics fabricated via in situ reaction by spark plasma sintering

A nano-laminated Y3Si2C2 ceramic material was successfully synthesized via an in situ reaction between YH2 and SiC using spark plasma sintering technology. A MAX phase-like ternary layered structure of Y3Si2C2 was observed at the atomic-scale by high resolution transmission electron microscopy. The lattice parameters calculated from both X-ray diffraction and selected area electron diffraction patterns are in good agreement with the reported theoretical results. The nano-laminated fracture of kink boundaries, delamination, and slipping were observed at the tip of the Vickers indents. The elastic modulus and Vickers hardness of Y3Si2C2 ceramics (with 5.5 wt% Y2O3) sintered at 1500 °C were 156 and 6.4 GPa, respectively. The corresponding values of thermal and electrical conductivity were 13.7 W·m-1·K-1 and 6.3×105 S·m-1, respectively.


Introduction 
Rare earth silicide carbides (RE 3 Si 2 C 2 , RE = Y, La-Nd, Sm, Gd-Tm) belong to a new group of ternary layered structure materials, which were first developed by Gerdes et al. [1,2]. The crystal structure of these compounds shows an orthorhombic subcell and consists of at least two different superstructures [2]. All the RE 3 Si 2 C 2 compounds were reported to have metallic conductivity and their magnetic ordering temperatures are lower than 60 K [1]. Y 3 Si 2 C 2 is one of a typical representative member of the RE 3 Si 2 C 2 group. In the Y 3 Si 2 C 2 structure, the c axis of the subcell is doubled, and thus it crystallizes in the body-centered orthorhombic system with space group Imma (No. 74) [3]. On the other hand, in Y 3 Si 2 C 2 structure, Y atoms form two-dimensionally arranged infinite sheets of edge-sharing octahedra containing C 2 pairs, wherein zig-zag chains of Si atoms are interleaved [2].
Even though the crystal structure of Y 3 Si 2 C 2 is different than hexagonal structure (space group P6 3 /mmc) of typical MAX phases (where M is an early transition metal, A is an A-group element, and X is carbon or nitrogen [4]), both phases have similar characteristics of layered structure and anisotropic bonding. Therefore, they both exhibit anisotropic electrical conductivity, anisotropic mechanical properties, and a low shear deformation resistance [3,[5][6][7]. Zhou et al. [3] theoretically predicted that the bulk modulus and shear modulus of Y 3 Si 2 C 2 are 93 and 50 GPa, respectively. Moreover, it was concluded that it is a soft ceramic material (Vickers hardness of 6.9 GPa) with good damage tolerance, due to the low shear deformation resistance and low Pugh's ratios (G/B = 0.537, where G is the shear modulus and B is the bulk modulus). Furthermore, the calculated volume expansion upon oxidation of Y 3 Si 2 C 2 was found to be ~26%, which could potentially lead to the sealing of the cracks between silicon carbide fibers (SiC f ) and SiC matrix. Therefore, Y 3 Si 2 C 2 may be a promising interphase material for SiC fiber-reinforced SiC matrix (SiC f /SiC) composite, because of its easy cleavage, low shear deformation resistance, and low volume expansion upon oxidation [3].
On the other hand, Y 3 Si 2 C 2 is inert when in contact with SiC at temperatures up to 1560 ℃, while a liquid phase can be formed at temperatures above 1560 ℃ via a ternary eutectic reaction, according to the calculated Y-Si-C ternary phase diagram [8]. Thus, Y 3 Si 2 C 2 was successfully used as the sintering additive for SiC and/or SiC/Al 4 SiC 4 systems [9,10]. The presence of a liquid phase not only effectively promotes the densification of SiC and/or SiC/Al 4 SiC 4 , but also improves the fracture toughness of ceramics by optimizing the grain boundary structure. Most importantly, Y 3 Si 2 C 2 can decompose to SiC and Y 2 O 3 (might act as the sintering additives for SiC) at ~1600 . ℃ Therefore, Y 3 Si 2 C 2 was also successfully used as a transition phase to achieve the seamless joining of SiC ceramics [11]. The joining mechanism was identified as follows: First, the laminated Y 3 Si 2 C 2 structure was formed by the in situ reaction between Y coatings with a thickness of 500 nm and SiC matrix in the joining layer at 1400 ℃. When the joining temperature increased to 1900 , ℃ Y 3 Si 2 C 2 disappeared owing to its decomposition at high temperatures. More recently, high-entropy RE 3 Si 2 C 2 /rare earth oxides with strong electromagnetic (EM) wave absorption capability and wide efficient absorption bandwidth were successfully synthesized. This can significantly broaden the applications potential of RE 3 Si 2 C 2 materials [12]. Due to its layered structure, good oxidation resistance, metallic conductivity as well as excellent EM wave absorption capability, dense Y 3 Si 2 C 2 ceramics might potentially be used as a stealth material and an on-beam-line higher-order-mode (HOM) load.
Even though Y 3 Si 2 C 2 has been demonstrated as the promising sintering additive and joining material for SiC-based advanced ceramics, the synthesis method and basic properties (besides electrical and magnetic properties) of Y 3 Si 2 C 2 bulk ceramics have not yet been investigated. The only reported technique to synthesize Y 3 Si 2 C 2 bulk ceramics consisted of arc-melting of the cold-pressed pellets containing the mixture of Y, Si, and C, followed by their annealing in an evacuated silica tube for 30 days at 900 [1]. This process was found to be ℃ extremely time consuming, because Y ingots were used as raw materials and the reaction temperature was as low as 900 . Spark plasma sintering (SPS) is an effective ℃ consolidation technology, which enables densification of ceramics at lower sintering temperatures and shorter time when compared to conventional methods. This is attributed to the presence of high-density electric current, which can promote mass diffusion [13,14].
Therefore, the Y 3 Si 2 C 2 nano-laminated bulk ceramics was successfully fabricated by the in situ reaction via SPS in this study. Furthermore, the phase composition, microstructure, mechanical properties, as well as electrical and thermal conductivities of Y 3 Si 2 C 2 were investigated. Both the Vickers hardness and the elastic modulus of the as-produced Y 3 Si 2 C 2 ceramics were in good agreement with the reported theoretically calculated results.

1 Preparation of Y 3 Si 2 C 2
YH 2 powder (Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) with a purity of 99.5% and a mean particle size of 75 μm, and β-SiC powder (99.5%, Eno Material Co., Ltd., Qinhuangdao, China) with a mean particle size of 0.5 μm, were used as raw materials. For the formation of Y 3 Si 2 C 2 , the YH 2 and SiC powders were mixed in a stoichiometric ratio of 3.05:2. The in situ reaction sintering process was performed in an SPS furnace (HPD 25/1, FCT systems, Germany) at the temperature range of 1300-1500 ℃ for 30 min under a uniaxial pressure of 30 MPa in an Ar atmosphere. The heating and cooling rates were 50 ℃·min 1 . The as-obtained Y 3 Si 2 C 2 ceramics surfaces were polished using the final 1 µm diamond suspension.

2 Material characterization
The phase compositions of the samples were identified by X-ray diffraction (XRD, D8 Advance, Bruker AXS, Germany) with Cu Kα radiation (λ = 1.5406 Å) under an operating voltage of 40 kV and a current of 40 mA at a step scan of 0.02 (°)/2θ and a step time of 0.2 s. The quantitative phase composition and lattice parameters of the Y 3 Si 2 C 2 phase were determined by Rietveld refinement using the TOPAS software.
The surface and fracture micromorphologies of the specimens were studied by the scanning electron microscope (SEM, Quanta 250 FEG, FEI, USA), equipped with an energy dispersive spectroscopy (EDS) detector. The phase distributions were characterized by electron back-scattered diffraction (EBSD) using a thermal field emission electron scanning microscope (Verios G4 uc, Thermo Scientific, USA), equipped with an EBSD apparatus operating at 20 kV accelerating voltage. For the EBSD analysis, the samples were polished with the final 1 μm diamond suspension, followed by etching using an ion beam (BIB, TIC 3X, Leica, Germany) for 3 h [15]. The microstructure and phase compositions were investigated by a transmission electron microscope (TEM, Talos™ F200x, Thermo Fisher Scientific, USA) system equipped with EDS system. Thin foils for TEM observations were prepared by focused ion beam (FIB, Auriga, Carl Zeiss, UK) technique.

3 Measurement of properties
Apparent density (ρ) of the samples was determined by the Archimedes' method. Elastic modulus was measured using a nanoindentation system (Hysitron PI85, Bruker, USA) on the polished surfaces of the specimens. Hardness of the materials was measured using a Vickers diamond indenter (HVs-1000 Digital Micro Vickers Hardness Tester, Beijing Times Mountain Peak Technology Co., China) under a load of 0.5, 2, and 5 N, respectively. The dwell time at a maximum load was always 10 s. At least 20 indents were measured for each specimen. Electrical resistivity of the samples was determined using a four-probe resistance tester (Cresbox, Napson Co., Japan). The thermal diffusivity coefficient (α) and specific heat capacity (C p ) were measured by laser flash method using a Netzsch LFA 457 apparatus (LFA, NETZSCH-Gerätebau GmbH, Germany). The thermal conductivity (κ, W·m 1 ·K 1 ) was calculated according to Eq. (1) [16]: 3 Results and discussion Figure 1 shows the XRD patterns of the samples sintered at different temperatures. Y 3 Si 2 C 2 (JCPDS No. 70-2799) was the predominant phase in all materials, while some Y 2 O 3 (JCPDS No. 83-0927) as an impurity phase was also detected. Rietveld refinement technique was applied to determine the fundamental parameters. The amount of the predominant Y 3 Si 2 C 2 phase was 88.4, 94.3, and 94.5 wt% for the samples sintered at 1300, 1400, and 1500 ℃, respectively. The corresponding amount of the Y 2 O 3 phase was 11.6, 5.7, and 5.5 wt%, respectively. Figure 2 shows the typical Rietveld refinement of XRD pattern of the sample sintered at 1500 ℃. No significant change in the lattice parameters of Y 3 Si 2 C 2 phase was observed for the samples sintered at different temperatures (Table 1). This indicated that the lattice parameters of the Y 3 Si 2 C 2 phase were not affected by the sintering temperature. Furthermore, the lattice parameters obtained from the Rietveld refinement are in good agreement with those determined by both the experimental measurements [2] and the calculation results [3], as presented in Table 1. The reliability of the refinement was confirmed by the reliability factors (R wp ), which were 9.1%, 9.0%, and 8.6% for the samples sintered at 1300, 1400, and 1500 ℃, respectively (Table 1). In addition, the crystal parameters refined from XRD of the powder samples were similar to the bulk samples (Table 1 and Fig. 2) sintered at 1500 ℃. It indicated that the influence of possible residual stresses on the lattice parameters in bulk samples on the XRD refinement was minimal.    It is believed that the Y 3 Si 2 C 2 phase was formed by the reaction between SiC and Y, which was generated from the decomposition of YH 2 at a low temperature (~650 ℃) [17], according to Eqs. (2) and (3): According to Reaction (4), the formation of Y 2 O 3 can be attributed to the presence of oxygen, which was introduced into the samples during powder homogenization process or during sintering at high temperatures (as a trace impurity in Ar atmosphere). Since it is expected that the amount of Y decreased with the increasing sintering temperature from 1300 to 1400 ℃, the amount of Y 2 O 3 impurity decreased from 11.6 to 5.7 wt%, accordingly. However, the amount of Y 2 O 3 in the sample sintered at 1500 ℃ was only slightly lower than that of the sample sintered at 1400 ℃. This suggested that the reaction between YH 2 and SiC was almost completed at 1500 ℃.
The theoretical density of the samples was calculated by the rule of mixture, taking the actual amount of the individual phases into account [18]. The theoretical density of Y 3 Si 2 C 2 (4.547 g·cm 3 ) and Y 2 O 3 (5.02 g·cm 3 ) was used. The calculated theoretical density of the bulk samples was 4.596, 4.574, and 4.565 g·cm 3 for the samples sintered at 1300, 1400, and 1500 ℃, respectively. Thus, the relative densities of as-obtained ceramics were 98.0% (1300 ℃), 99.0% (1400 ℃), and 99.5% (1500 ℃), respectively. It is clear that the highly dense Y 3 Si 2 C 2 ceramic materials were successfully obtained by the in situ solid state reaction in a significantly shorter time when compared to the previous study, in which Y ingots were used [1]. The use of pulsed current sintering probably improved the mass diffusion and promoted the solid-state reaction, which enable densification to be completed in a relatively short period of time [19]. In addition, YH 2 is more stable than Y, which easily oxidizes to Y 2 O 3 during the experimental process. The oxidized layer of Y 2 O 3 would inhibit the diffusion and reaction between Y and SiC. Thus, the use of YH 2 as a raw material instead of Y probably accelerated the diffusion process, but also decreased the synthesis temperature due to the decomposition of YH 2 at a relatively low temperature (650 ℃ ) [17]. Furthermore, the presence of YH 2 hydride powder may have facilitated the nucleation of  [20]. Figure 3 shows the microstructure of the samples sintered at different temperatures, detected by EBSD. Figures 3(a)-3(c) present the diffraction pattern quality quantified using the "band contrast", while Figs. 3(d)-3(f) show the phase distribution of Y 3 Si 2 C 2 (in red) and Y 2 O 3 (in blue). The elongated, plate-like morphology of Y 3 Si 2 C 2 was clearly identified. The phase fraction of Y 3 Si 2 C 2 measured in the observed area increased with increasing sintering temperature: 80% (1300 ℃), 84% (1400 ℃), and 91% (1500 ℃). At the same time, the grain size distribution is shown in Figs. 3(g)-3(i). The mean grain size of the materials increased from 3.9 μm (1300 ℃) to 8.8 μm (1500 ℃). The abnormal grain growth was obviously observed when the sintering temperature was increased to 1400 and 1500 ℃. Figures 4(a)-4(c) show the fracture surfaces of Y 3 Si 2 C 2 sintered at 1300, 1400, and 1500 ℃ , respectively. The failure mode was mainly intragranular, because of a low shear deformation resistance of Y 3 Si 2 C 2 [3]. Some pores were observed for the sample sintered at 1300 ℃ (Fig. 4(a)). On the other hand, almost fully dense Y 3 Si 2 C 2 was observed after sintering at 1400 ℃ (Fig. 4(b)) and 1500 ℃ (Fig. 4(c)).  TEM analysis was carried out to observe the atomic-scale microstructure of the Y 3 Si 2 C 2 sintered at 1500 ℃. Figures 5(a)-5(e) show a high angle annular dark field (HAADF) image and the corresponding elemental distribution of Y, C, O, and Si, respectively. The semi-quantitative EDS analysis confirmed the presence of Y 3 Si 2 C 2 and Y 2 O 3 , which correspond to the points 1 and 2 in Fig. 5(a), respectively. The EDS results are presented in Table 2. The presence of a small amount of oxygen was not confirmed. The  atomic-scale microstructure along the [001] zone axis was confirmed by HRTEM and corresponding SAED pattern shown in Figs. 5(f) and 5(g). The layered atomic stacking can be clearly seen in the HRTEM image. The lattice fringe spacing of 0.786 nm can be assigned to the (020) planes of Y 3 Si 2 C 2 , as shown in Fig. 5(g). The corresponding SAED pattern also confirmed the orthorhombic crystal structure of Y 3 Si 2 C 2 (Fig. 5(f)). The lattice parameters were derived to be a = 8.44 Å and b = 15.72 Å, which are in good agreement with those determined from the XRD pattern (Table 1). The properties of the as-sintered Y 3 Si 2 C 2 are listed in Table 3, along with the properties of some typical ternary carbides. Both the elastic modulus and the Vickers hardness of the Y 3 Si 2 C 2 materials decreased with increasing sintering temperature. This was probably caused by a decreasing amount of Y 2 O 3 in the materials with increasing temperature. The elastic modulus and Vickers hardness of Y 2 O 3 are ~180 and 7.6 GPa [21][22][23], respectively, which are slightly higher than the calculated values for Y 3 Si 2 C 2 [3]. Moreover, the grain size increased with increasing sintering temperature, and thus the Vickers hardness also decreased with the increase in the sintering temperature according to the Hall-Petch relationship. The elastic modulus of the sample sintered at 1500 ℃ was close to the calculated values reported by Zhou et al. [3]. The Vickers hardness of the sample sintered at 1500 ℃ was 7.2±0.8, 6.5±0.5, and 6.4±0.4 GPa for the indentation load of 0.5, 2, and 5 N, respectively. These values are in good agreement with the reported calculated value of 6.9 GPa [3].
The shape of Vickers indents was irregular with the exfoliated surfaces and deformed particles, which is like the typical indent shape of Ti 3 SiC 2 MAX phase [24]. A typical surface morphology at the tip of a Vickers indent for the sample sintered at 1500 ℃ is shown in Figs. 6(a) and 6(b). Interestingly, in the case of basal Y 3 Si 2 C 2 plane oriented parallel to the indentation load, typical nano-laminated fracture was observed, owing to the kink boundaries, delamination, and slipping ( Fig. 6(a)). Such behavior is commonly  observed for the group of MAX phases, which belong to typical damage tolerant ceramics [27]. On the other hand, when the basal plane of Y 3 Si 2 C 2 was oriented in a direction perpendicular to the indentation load, the exfoliation and sharp steps-like fracture caused by crack deflection inside the Y 3 Si 2 C 2 grains were observed ( Fig. 6(b)). The fracture energy can be consumed by virtue of crack deflection. A typical nano-laminated MAX phase-like structure of Y 3 Si 2 C 2 is shown in Fig. 6(c), which can be easily recognized by its cleavage nature. A relatively low Vickers hardness and typical nano-laminated fracture behavior indicated that Y 3 Si 2 C 2 belongs to the group of soft ceramics. Zhou et al. [3] reported that the low shear deformation resistance along the (010) [101] slip system could be attributed to the weak metallic bonding between Y 2 -C. Thermal conductivity of the Y 3 Si 2 C 2 samples decreased from 16.1 to 13.7 W·m 1 ·K 1 with the increase in the sintering temperature from 1300 to 1500 ℃. This was observed despite the fact that the grain size increased with increasing sintering temperature (Fig. 3), which usually leads to the improved thermal conductivity due to the decreased phonon scattering by the grain boundaries. Therefore, the decreased thermal conductivity with increasing sintering temperature in this study can be attributed to the presence of Y 2 O 3 (27 W·m 1 ·K 1 ) [28], whose content also decreased with increasing sintering temperature.
The electrical resistivity of the samples is presented in Table 3. The corresponding electrical conductivity of the samples sintered at 1300, 1400, and 1500 ℃ were 7.6×10 5 , 7.6×10 5 , and 6.3×10 5 S·m 1 , respectively. The electrical conductivity of Y 3 Si 2 C 2 was mainly affected by Y1 4de g , Y2 4dt 2g , C 2p x , and C 2p z states (x and z are inclined to the x and z axis, respectively, at about 45°), according to the analysis of the projected density of states and the decomposed distribution of electron density [3].

Conclusions
The highly dense Y 3 Si 2 C 2 ceramic material was fabricated by in situ solid state reaction between YH 2 and SiC via SPS. The as-obtained Y 3 Si 2 C 2 ceramic exhibited a nano-laminated structure, which was confirmed by HRTEM analysis. The lattice parameters were derived as a = 8.4418 Å, b = 15.6671 Å, and c = 3.863 Å by the Rietveld refinement of XRD patterns. The experimentally measured elastic modulus (156 GPa) and Vickers hardness (6.4 GPa) of the fabricated ceramics are in good agreement with the reported theoretically calculated values. Typical nano-laminated fracture www.springer.com/journal/40145 behavior was observed at the tip of Vickers indents, which indicated that Y 3 Si 2 C 2 belongs to the group of soft ceramics. The thermal and electrical conductivity of the sample sintered at 1500 ℃ was 13.7 W·m 1 ·K 1 and 6.3×10 5 S·m 1 , respectively. The proposed synthesized strategy could potentially be used to fabricate other RE 3 Si 2 C 2 phases.
properties of the monolithic CVD β-SiC materials joined with a pre-sintered MAX phase Ti 3  Open Access This article is licensed under a Creative Commons Attribution 4.0 International License, which permits use, sharing, adaptation, distribution and reproduction in any medium or format, as long as you give appropriate credit to the original author(s) and the source, provide a link to the Creative Commons licence, and indicate if changes were made. The images or other third party material in this article are included in the article's Creative Commons licence, unless indicated otherwise in a credit line to the material. If material is not included in the article's Creative Commons licence and your intended use is not permitted by statutory regulation or exceeds the permitted use, you will need to obtain permission directly from the copyright holder.