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

Abstract: 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 ℃ 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.

Keywords: Y3Si2C2; rare earth silicide carbides; spark plasma sintering (SPS); ternary layered structure ceramic; properties 

1 Introduction

Rare earth silicide carbides (RE3Si2C2, 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 RE3Si2C2 compounds were reported to have metallic conductivity and their magnetic ordering temperatures are lower than 60 K [1]. Y3Si2C2 is one of a typical representative member of the RE3Si2C2 group. In the Y3Si2C2 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 Y3Si2C2 structure, Y atoms form two-dimensionally arranged infinite sheets of edge-sharing octahedra containing C2 pairs, wherein zig-zag chains of Si atoms are interleaved [2]

Even though the crystal structure of Y3Si2C2 is different than hexagonal structure (space group P63/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 [3,5–7]. Zhou et al. [3] theoretically predicted that the bulk modulus and shear modulus of Y3Si2C2 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 Y3Si2C2 was found to be ~26%, which could potentially lead to the sealing of the cracks between silicon carbide fibers (SiCf) and SiC matrix. Therefore, Y3Si2C2 may be a promising interphase material for SiC fiber-reinforced SiC matrix (SiCf/SiC) composite, because of its easy cleavage, low shear deformation resistance, and low volume expansion upon oxidation [3]

On the other hand, Y3Si2C2 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, Y3Si2C2 was successfully used as the sintering additive for SiC and/or SiC/Al4SiC4 systems [9,10]. The presence of a liquid phase not only effectively promotes the densification of SiC and/or SiC/Al4SiC4, but also improves the fracture toughness of ceramics by optimizing the grain boundary structure. Most importantly, Y3Si2C2 can decompose to SiC and Y2O3 (might act as the sintering additives for SiC) at ~1600 . ℃ Therefore, Y3Si2C2 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 Y3Si2C2 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 , ℃ Y3Si2C2 disappeared owing to its decomposition at high temperatures. More recently, high-entropy RE3Si2C2/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 RE3Si2C2 materials [12]. Due to its layered structure, good oxidation resistance, metallic conductivity as well as excellent EM wave absorption capability, dense Y3Si2C2 ceramics might potentially be used as a stealth material and an on-beam-line higher-order-mode (HOM) load. 

Even though Y3Si2C2 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 Y3Si2C2 bulk ceramics have not yet been investigated. The only reported technique to synthesize Y3Si2C2 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 sint ℃ ering (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 Y3Si2C2 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 Y3Si2C2 were investigated. Both the Vickers hardness and the elastic modulus of the as-produced Y3Si2C2 ceramics were in good agreement with the reported theoretically calculated results. 

2 Experimental 

2. 1 Preparation of Y3Si2C2

YH2 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 Y3Si2C2, the YH2 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 Y3Si2C2 ceramics surfaces were polished using the final 1 µm diamond suspension.

2. 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 Y3Si2C2 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. 

2. 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 (Cp) 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]:

κ = αρCp      (1) 

3 Results and discussion 

Figure 1 shows the XRD patterns of the samples sintered at different temperatures. Y3Si2C2 (JCPDS No. 70-2799) was the predominant phase in all materials, while some Y2O3 (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 Y3Si2C2 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 Y2O3 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 Y3Si2C2 phase was observed for the samples sintered at different temperatures (Table 1). This indicated that the lattice parameters of the Y3Si2C2 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 (Rwp), 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. 

Fig. 1 XRD patterns of the samples sintered at various temperatures by reactive SPS.

Fig. 2 Rietveld refinement of the XRD patterns of the samples synthesized at 1500 ℃: (a) bulk and (b) powder. 

Table 1 Experimental lattice parameters of Y3Si2C2 derived from Rietveld refinement and selected area electron diffraction (SAED) patterns, and their comparison with the calculated and experimental values reported in literature

It is believed that the Y3Si2C2 phase was formed by the reaction between SiC and Y, which was generated from the decomposition of YH2 at a low temperature (~650 ℃) [17], according to Eqs. (2) and (3):

YH2 → Y+H2↑   (2) 

Y+SiC → Y3Si2C2   (3) 

Y+O2 → Y2O3   (4) 

According to Reaction (4), the formation of Y2O3 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 Y2O3 impurity decreased from 11.6 to 5.7 wt%, accordingly. However, the amount of Y2O3 in the sample sintered at 1500 ℃ was only slightly lower than that of the sample sintered at 1400 ℃. This suggested that the reaction between YH2 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 Y3Si2C2 (4.547 g·cm-3) and Y2O3 (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 Y3Si2C2 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, YH2 is more stable than Y, which easily oxidizes to Y2O3 during the experimental process. The oxidized layer of Y2O3 would inhibit the diffusion and reaction between Y and SiC. Thus, the use of YH2 as a raw material instead of Y probably accelerated the diffusion process, but also decreased the synthesis temperature due to the decomposition of YH2 at a relatively low temperature (650 ℃) [17]. Furthermore, the presence of YH2 hydride powder may have facilitated the nucleation of Y3Si2C2 [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 Y3Si2C2 (in red) and Y2O3 (in blue). The elongated, plate-like morphology of Y3Si2C2 was clearly identified. The phase fraction of Y3Si2C2 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 ℃. 

Fig. 3 Microstructures of samples fabricated at various temperatures observed by EBSD. Micrograph in band contrast: (a) 
1300, (b) 1400, (c) 1500 ℃; phase distribution: (d) 1300, (e) 1400, (f) 1500 ℃ (Y3Si2C2 in red and Y2O3 in blue); grain size 
distribution: (g) 1300, (h) 1400, (i) 1500 ℃. 

Figures 4(a)–4(c) show the fracture surfaces of Y3Si2C2 sintered at 1300, 1400, and 1500 ℃ , respectively. The failure mode was mainly intragranular, because of a low shear deformation resistance of Y3Si2C2 [3]. Some pores were observed for the sample sintered at 1300 ℃ (Fig. 4(a)). On the other hand, almost fully dense Y3Si2C2 was observed after sintering at 1400 ℃ (Fig. 4(b)) and 1500 ℃ (Fig. 4(c)). 

Fig. 4 SEM images of the fracture surfaces of Y3Si2C2 sintered by reactive SPS at (a) 1300, (b) 1400, and (c) 1500 ℃.

TEM analysis was carried out to observe the atomic-scale microstructure of the Y3Si2C2 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 Y3Si2C2 and Y2O3, 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 Y3Si2C2, as shown in Fig. 5(g). The corresponding SAED pattern also confirmed the orthorhombic crystal structure of Y3Si2C2 (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). 

Fig. 5 TEM image of Y3Si2C2 ceramics sintered at 1500 ℃: (a) HAADF image and elemental distribution of (b) Y, (c) C, (d) O, (e) Si; (f) SAED pattern for the yellow area in (a); (g) high resolution transmission electron microscopy (HRTEM) image of Y3Si2C2, the insert includes structure models showing the positions of Y, Si, and C. HRTEM image and corresponding SAED parten were confirmed along the [001] zone axis. 

Table 2 EDS results of the spots 1 and 2 in Fig. 5(a) 

The properties of the as-sintered Y3Si2C2 are listed in Table 3, along with the properties of some typical ternary carbides. Both the elastic modulus and the Vickers hardness of the Y3Si2C2 materials decreased with increasing sintering temperature. This was probably caused by a decreasing amount of Y2O3 in the materials with increasing temperature. The elastic modulus and Vickers hardness of Y2O3 are ~180 and 7.6 GPa [21–23], respectively, which are slightly higher than the calculated values for Y3Si2C2 [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]

Table 3 Density, mechanical, thermal, and electrical properties of the as-obtained Y3Si2C2, and their comparison with the reported values of typical ternary layered structural ceramics 

The shape of Vickers indents was irregular with the exfoliated surfaces and deformed particles, which is like the typical indent shape of Ti3SiC2 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 Y3Si2C2 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 Y3Si2C2 was oriented in a direction perpendicular to the indentation load, the exfoliation and sharp steps-like fracture caused by crack deflection inside the Y3Si2C2 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 Y3Si2C2 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 Y3Si2C2 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 Y2–C.  

Fig. 6 Surface morphologies of sample sintered at 1500 ℃ after Vickers indentation test: (a) basal Y3Si2C2 plane oriented 
parallel to the load with the presence of kinks, delamination, and slipping; (b) basal Y3Si2C2 plane oriented perpendicular to the load showing the exfoliation; and (c) high magnification SEM image showing the nano-laminated structure of Y3Si2C2

Thermal conductivity of the Y3Si2C2 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 Y2O3 (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 7.6×105, 7.6×105, and 6.3×105 S·m-1, respectively. The electrical conductivity of Y3Si2C2 was mainly affected by Y1 4deg, Y2 4dt2g, C 2px', and C 2pz' 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]

4 Conclusions 

The highly dense Y3Si2C2 ceramic material was fabricated by in situ solid state reaction between YH2 and SiC via SPS. The as-obtained Y3Si2C2 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 behavior was observed at the tip of Vickers indents, which indicated that Y3Si2C2 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 6.3×105 S·m-1, respectively. The proposed synthesized strategy could potentially be used to fabricate other RE3Si2C2 phases. 

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