Structure and Properties of Spark Plasma Sintered SiC Ceramics with Oxide Additives

Abstract

This article describes spark plasma sintering of ceramics based on silicon carbide with nanoadditives, as follows: MnO_nano 5.5 wt. % + Al₂O_3nano 2.0 wt. % + SiC_nm (37–57 wt. %) + SiC_µm (31–51 wt. %) + SiO_2µm 4.5 wt. %. Sintering was carried out at 2000 °C. The diffraction pattern of the analyzed sample showed the presence of silicon carbide with a hexagonal crystal lattice. Residual amounts of rhombohedral SiC, α-Fe, and a solid solution of silicon in iron were also found. The method of thermogravimetric analysis established the change in mass, heat flow, temperature of the samples, and the change in the partial pressures of gases during the experiment. Samples obtained by SPS show a higher density of the material at the level of 3.3 g/cm³, average mechanical strength of 454 MPa, and microhardness of 35 GPa, compared with samples obtained by liquid-phase sintering. The SPS method also made it possible to obtain materials with a higher density (by 8%) and practically no significant crystal growth compared to samples obtained by liquid phase sintering. The results of the study facilitate the achievement of a combination of new approaches to the design of compositions and the technology of manufacturing SiC ceramics, which significantly expands their areas of application.

1. Introduction

SiC ceramics is becoming a popular material for various technical applications due to its mechanical and thermal properties. This material is used in the production of armor materials and the nuclear fuel, aerospace, and marine industries. It is used in the automotive industry as well as in other modern industries, with ever-higher demands for materials.

Of particular interest are studies on the possibility of obtaining high-tech ceramics with new structural and physico-mechanical properties with the introduction of nanoadditives. The melting temperature of nanoparticles of oxide systems in the process of sintering the SiC ceramics is significantly lower than that of the same particles of micron size; there is an increased chemical activity and a decrease in the interaction temperature during reactions in systems with nanoadditives [1,2,3].

At present, activating additives for sintering materials based on silicon carbide in the form of oxide systems are widely used: MnO-Al₂O₃, Y₂O₃-Al₂O₃, MgO-Y₂O₃-Al₂O₃, etc. During the sintering of ceramics with nano-additives, the effect of inhibiting the growth of micron-sized crystals can be observed due to changes in the surface energy of grain boundaries and their mobility. Diffusion processes of interaction between micro- and nanoparticles proceed with different mass transfer rates. This effect leads to the formation and development of interparticle boundaries and, consequently, to an increase in the density, strength, and other properties of ceramics with nanoadditives [3,4,5].

In recent years, rather widely highly dispersed powders of eutectic compositions consisting of nanoparticles, close in dispersion to the nanolevel with aggregate sizes of ~0.2–0.3 μm, have been used in oxide systems. The use of such additives in the production of ceramics made it possible to significantly reduce the sintering temperature and obtain high-density ceramics with a high level of properties [6]. In this case, it is necessary that the basis of ceramics–aluminum oxide is also a highly dispersed system.

One possibility for improving high-temperature properties is to use additives that can be dissolved in the SiC lattice (solid solution formation). It is shown that AlN and aluminum form a solid solution with SiC. AlN is soluble in 4H and 6H polytypes, as well as in 2H. The bending strength of the material sintered with AlN and Y₂O₃ is 524 MPa at 1000 °C and 377 MPa at 1400 °C in air [7].

Another significant technological problem when working with nanomaterials is the technological stage of charge preparation, which includes uniform mixing of the initial components. It is necessary to achieve a uniform distribution of oxide nanoparticles between SiC particles and obtain a homogeneous charge powder. Only with a uniform distribution of oxide nanoparticles will it be possible to obtain ceramics with a high level of mechanical properties commensurate with the level of properties of hot-pressed silicon carbide materials.

However, it is possible to achieve complete compaction by sintering at high temperatures and short exposure time. One of the ways to solve the problem is to use metal and ceramic nanopowders as modifiers. The most promising additive that makes it possible to effectively control the structure of ceramics based on silicon carbide, among others, is a modifier from SiC nanopowder [8].

The technology of spark plasma sintering (SPS) is a modern promising technology for the consolidation of powder materials. The method is based on passing powerful rectangular DC pulses through the sintered material, which, in addition to resistive heating, leads to the activation of the powder surface due to the formation of spark plasma in the gaps. SPS is a complex multifactorial process. Many processes that occur during plasma sintering are still not clear enough [9,10]. The physical description of SPS is based on the classical theory of sintering; however, several aspects related to the fact that during high-speed heating, the sintering process occurs under essentially nonequilibrium and nonstationary conditions cannot be described on the basis of the classic theory of sintering. The main advantages of SPS technology include uniform distribution of heat over the sample; high density and controlled porosity; pre-treatment by pressure and binders not being required; uniform sintering of homogeneous and dissimilar materials; short cycle time; and minimum grain growth and influence on the microstructure [11,12,13].

Recently, special interest has been shown in the production of porous nanostructured silicon carbide ceramics. Nanoceramics assumes an improvement in properties because of grain size reduction.

Obtaining dense SiC can be achieved with the use of various additives or pressing pressure during sintering. This is possible when using ultradispersed SiC powder and nanodispersed powders of oxide additives, usually Al₂O₃ and MnO. Thus, the aim of this study is the synthesis of ceramics based on silicon carbide by spark plasma sintering with the use of activating additives MnO_nano 5.5 wt. % + Al₂O_3nano 2.0 wt. % + SiC_nm (37–57 wt. %) + SiC_µm (31–51 wt. %) + SiO_2µm 4.5 wt. %, contributing to the improvement in the physical and mechanical properties of the material obtained and a comparative assessment of the physical and mechanical properties of ceramics obtained earlier by the method of liquid-phase sintering with eutectic additives.

Our earlier studies led to the conclusion that the optimum sintering temperature of ceramics based on silicon carbide can be reduced during the liquid-phase formation of a structure with nanoadditives, which made it possible to reduce the temperature to an average of 1800 °C [14,15]. The appearance of a melt of oxide additives in such a system had a positive effect on the sintering process, which contributed to the reorientation of silicon carbide grains and the filling of voids that form at the stage of product molding. Previously, we substantiated the choice of sintering additives. The characteristic covalent type of chemical bonding of silicon carbide SiC affects its density without sintering additives; therefore, the liquid-phase sintering process of SiC ceramics was carried out using the oxides MnO_nano, Al₂O_3nano, and SiO_2μm, with the ratio of components and additives forming the liquid phase: MnOnano 2.5 wt. % + Al₂O_3nano 2.0 wt. % + SiCµm 91 wt. % + SiO_2µm 4.5 wt. % [15]. At the same time, the technology of consolidating powder mixtures was carried out due to ultrasound influence.

It has been established that due to the nonequilibrium state of nanomaterials and the fact that the sintering of silicon carbides requires a high temperature and a certain holding time at this temperature (up to several hours), such as in liquid-phase sintering, the production of dense ceramics is significantly affected by the grain growth factor during long-term holding [16,17,18,19,20]. Since silicon carbide has rather low values of the surface and volume diffusion coefficients, an increase in temperature leads to the decomposition of silicon carbide into elemental components. Therefore, the aim of this study was to apply the method of spark plasma sintering for the synthesis of silicon carbide ceramics and to study the structure and properties of the synthesized materials. It is assumed that it is possible to achieve complete compaction due to sintering at high temperatures and short exposure time, which is ensured by the SPS method. In this work, attention was also paid to minimizing the grain growth factor. To achieve high density and restrain grain growth, the consolidation of structural SiC ceramics was carried out using spark plasma sintering with the use of oxide additives.

2. Materials and Methods

2.1. Design of the Experiment

To achieve the goal, the study was performed according to the flowchart presented in Figure 1.

Based on previous studies [15] and a series of experimental results on the selection of the optimal chemical composition, three potential concentrations were determined:

• S-1—5.5 wt. % MnO_nano + 2.0 wt. % Al₂O_3nano + 57 wt. % SiC_nano + 31 wt. % SiC_micro + 4.5 wt. % SiO_2micro;
• S-2—5.5 wt. % MnO_nano + 2.0 wt. % Al₂O_3nano + 47 wt. % SiC_nano + 41 wt. % SiC_micro + 4.5 wt. % SiO_2micro;
• S-3—5.5 wt. % MnO_nano + 2.0 wt. % Al₂O_3nano + 37 wt. % SiC_nano + 51 wt. % SiC_micro + 4.5 wt. % SiO_2micro.

2.2. Materials

The powders used in this study were as follows: MnO (CAS No. 1344-43-0, nanopowder, 99%, Sigma-Aldrich Taufkirchen, Germany); Al₂O₃ (CAS No. 1344-28-1, nanopowder, Sigma-Aldrich Taufkirchen, Germany); SiC (CAS No. 409-21-2, 60–90 µm powder, nanopowder 200–500 nm, Sigma-Aldrich Taufkirchen, Germany); and SiO₂ (particle size 6–9 µm, KazSilicon LLP Bastobe village, Almaty Region, Republic of Kazakhstan). All materials were used in powder form.

The analysis of oxide powders was previously studied in detail [15]. The average size of the main initial component of the SiC powder was determined with Winner-2005A Intelligent Laser Particle Size Analyzer (Jinan Winner Particle Instrument Stock Co., Jinan, China) with measuring range of 0.1–1000 µm. The average particle size of SiC micropowder was 61.576 µm, and SiC nanopowder was 480 nm (Figure 2).

X-ray phase analysis was performed on a PANalytical X’Pert Pro instrument (Malvern Panalytical Ltd., Almelo, the Netherlands). During the study, a voltage of 40 kV and a current of 30 mA were applied to the anode copper tube (λ = 1.541 Å). The shooting was carried out in the range of angles from 10 to 80 degrees 2θ, the shooting step was 0.02, and the counting time was 0.5 s/step. The phase analysis based on the obtained diffraction lines was performed using the HighScore Plus V.5.2 and Mach 3 R3.043 software package.

Figure 3 shows the diffraction pattern of the analyzed SiC nanopowder. It was established that high-intensity peaks in the X-ray diffraction pattern belong to moissanite (SiC).

Images of silicon carbide nanoparticles obtained on a transmission microscope are shown in Figure 4. It can be seen that the nanopowder has an irregular shape.

2.3. Methods

The microstructure and point elemental analysis of the samples were studied on a scanning electron microscope JSM-6390LV (JEOL, Tokyo, Japan) with an energy dispersion analysis add-on unit (EDX) INCA ENERGY (Oxford Instruments, Abingdon, UK) with a high vacuum resolution of up to 3 nm. Electron diffraction patterns of samples were obtained by transmission electron microscopy (TEM) on JEM-2100 (JEOL, Tokyo, Japan). Powder samples for TEM were applied to a grid of amorphous carbon.

The microhardness of the samples was determined using the Vickers method. A diamond pyramid was used as an indenter; the pressure force was 500 N. The specific surface area was determined on a dispersion analysis device of the PSKh-17 series (“Laboratory of Scientific Instruments” LLC, Moscow, Russia). The gas permeability of the powder layer was determined by the duration of air filtration through it at a fixed initial and final rarefaction in the working volume of the device. Loss on ignition was determined according to the procedure described in ASTM D7348. The apparent density value was determined according to ASTM C20-00.

Determination of the ultimate strength in static bending was carried out on an ultimate tensile testing machine WDW-5 (Shandong Oberthur Group Co., Jinan, China) with a maximum load of 500 kgf. The method is based on the requirements of ASTM C1674-11. Loading the sample with uniformly increasing load to failure with load measurement. The ultimate strength of the sample was determined from the value of the breaking load. The samples were square-section plates 35 mm long, 10 mm wide, and 1 mm thick. The samples were tested at a strain rate of 1 mm/min.

Fracture toughness, in this study, is represented by the critical stress intensity factor K_Ic, which was evaluated by the straight-notched beam bending (SENB) method. The breaking force was determined on samples with a length of l-35 mm using a WDW-5 tensile testing machine (Shandong Oberthur Group Co., Jinan, China). The calculation was carried out according to the formula:

$$\sigma =Y·\frac{3{PlL}^{0.5}}{2{bh}^2},$$

(1)

where Y = 1.93 − 3.07(L/h) + 14.53(L/h)² − 25.11(L/h)³ + 25.80(L/h)⁴; L is the depth of the stress concentrator. The error in determining the value of Y was ±0.2%, and for K_Ic it was ±10%.

Assessment of the elastic properties of ceramics was carried out by the method of determining the dynamic modulus of elasticity, based on the dependence of the velocity of propagation of sound waves in ceramics on the modulus of elasticity. In this case, the resonance of the frequencies applied to the sample and natural bending vibrations of the ceramic rod were used. The sample was fixed on a nickel wire by wrapping and twisting it at each end of the sample. Each end was suspended from a piezoelectric element in such a way that the sample hung horizontally. Next, the frequency generator and oscilloscope were turned on, and the resonant frequency was determined from the corresponding picture on the screen, gradually increasing it on the generator. At resonance, an ellipse appeared on the screen. As the frequency increased or decreased, more or less resonant, the ellipse on the screen disappeared, turning into a narrow luminous strip. There is a relationship between the dynamic modulus of elasticity, mass, and dimensions of the sample and the speed of propagation of sound waves in the sample or the resonant frequency. Using this relationship, the modulus of elasticity (E, Pa) was calculated using the following formula:

$$E=0.9469·{10}^{-10}{\left(\frac{l}{h}\right)}^3\frac{m}{h}{f}_{res}^2,$$

(2)

where l is the sample length, cm; h is the height of the rectangular sample, cm; m is mass of the sample; and f_res is the resonant frequency.

To study the processes of thermal decomposition and synthesis of various compounds, as well as to identify the temperatures of phase transitions, the method of thermogravimetric analysis was used with simultaneous recording of the heating curves of the studied samples (T), mass change (TG), mass change rate (DTG) based on a TGA/DSC 3+ thermogravimetric analyzer (Mettler Toledo, Greifensee, Switzerland) using ThermoStar mass spectrometric gas analyzer (Pfeiffer Vacuum, Vienna, Austria) and modular humidity generator MHG100 (ProUmid, Baden-Württemberg, Germany). The registration of the change in the mass of the sample during the experiment was carried out by built-in ultramicrobalance with a resolution of 0.1 μg and an accuracy of 0.005% in the entire weighing range. Changes in the composition of the gaseous medium in the volume above the test sample were carried out by a ThermoStar quadrupole mass spectrometer (Pfeiffer Vacuum, Vienna, Austria) with Quadera software. In order to reduce the stabilization time of the gas composition in the device chamber and to reduce the amount of atmospheric gases, the N₂ purge rate was 50 mL/min. The maximum heating temperature was up to 1400 °C in order to avoid damage to the heater and thermocouple of the oven of the device. The error in determining the temperature is ±10 °C. Experiment conditions were as follows: sample heating temperature range from 100 to 1400 °C; linear heating rate of 10 °C/min; cooling rate of ~40 °C/min; nitrogen consumption in purge mode of 50 mL/min; and nitrogen consumption during the experiment of 50 mL/min.

2.4. Spark Plasma Sintering

Nanodispersed powders are unstable systems due to the high concentration of crystal lattice defects on the surface and in the bulk of the particle. The particles of such powders are prone to aggregation, which adversely affects the sintering process. Therefore, the introduction of nanoparticles into the composition of the mixture and mixing of its components were carried out in a specially designed impeller-type reactor [15]. Preservation of the nanostructured state of powders is necessary to ensure the sintering of powder systems according to the mechanism of viscous flow of a substance when sintering processes lead to the removal of porosity before the start of crystal growth. In this case, the ceramics are obtained without pores with controlled crystal size and, consequently, properties.

SiC nanopowder was used as a modification additive to effectively control the structure of ceramics based on silicon carbide. To intensify the sintering processes of SiC ceramics, dispersed and chemically pure powders were used to increase the specific surface area and, as a result, reduce the activation energy of diffusion processes.

The introduction of nanoparticles into the composition of the mixture is a difficult task from a technological point of view, since nanoparticles must be evenly distributed over the entire volume of the mixture to exclude their possible coagulation and agglomeration. The efficiency of introduction and uniformity of distribution of MnO, Al₂O₃ nanoparticles in the volume of the mixture of micron SiC, and SiO₂ powders was evaluated by tinting the mixture with the pigment rhodamine. The time and mixing mode were experimentally selected until the suspension was uniformly colored in a liquid medium, using non-polar liquids (distilled and deionized water), in order to prevent coagulation of nanoparticles. Mixing of the charge components was carried out in an impeller-type reactor. The installation is a cylindrical, vertically located stainless steel container. A shaft with blades is installed on a rigid base. During the rotation of the shaft, the charge is mixed (Figure 5).

In the charge, as well as in the initial raw materials, it is important to control the size component of the finely dispersed system. Particle size distribution for S1 sample is presented in Figure 6.

Table 1. Particle size analysis in samples S1–S3.

S1	S2	S3
2.880 nm	2.955 nm	3.004 nm

shows the results of measurements of the dimension of all three samples. According to the results, the average particle size in the charge of sample S1 is 288 nm, while the dimension of SiC nanoparticles in the initial nanopowder was 480 nm. The decrease in the size of nanoparticles apparently occurs because of the process of preparation on a ball and roller mills of both individual components of the charge, and together, with further mixing of the charge components in the reactor by wet mixing in distilled water with simultaneous bubbling under pressure.

To prepare a press mass, a 6–10% solution of polyvinyl alcohol (PVA) was weighed on a scale and poured into an enameled container. The container was placed in a pot of hot water located on an electric stove, the solution was stirred continuously, and it was heated to transparency. The container with the PVA solution was removed from the stove, cooled to room temperature, and then the solution was filtered through a nylon mesh No. 46 to remove pieces of film and impurities.

For binder preparation, PVA and glycerol were poured into a 500-mL polyethylene beaker in a ratio of 7:3, respectively. Glycerol was poured into the PVA solution gradually with constant stirring, and the mixture was thoroughly mixed until a uniform (homogeneous) consistency. The binder was prepared immediately before its introduction into the powder at the rate of 6–10% of the weight of the powder; the shelf life of the binder was not more than 24 h.

A baking sheet, a tray with powders of all compositions weighed and dried at 80 °C for 24 h, a container with the estimated amount of the binder, sieves, and a squeegee were installed in the fume hood. A binder was introduced into the powder in small portions with stirring, in a volume of 1/10 of its original amount. After the introduction of each portion of the ligament, the mass was thoroughly mixed until a homogeneous mass was obtained. A sieve with a mesh size of 0.9–1.0 mm was installed on a baking sheet. The resulting mixture in portions was rubbed three times in a circular motion through a sieve with a mesh size of 0.9 mm or 1.0 mm with a squeegee (Figure 7c). A sieve with a mesh size of 0.6 mm or 0.7 mm was installed. The wiped press mass in portions was rubbed in a circular motion with a squeegee three times through a sieve with a mesh size of 0.6 mm or 0.7 mm. The resulting press mass was evenly distributed on a baking sheet, and the baking sheet was covered with a cloth (coarse calico). The press mass in the trays was placed in a storage cabinet for aging, then the temperature in the storage cabinet was set to no higher than 40 °C. The maturation period of the press mass is from 4 to 6 days. The shelf life of the press mass after aging is not more than 6 days.

Sintering of SiC ceramic samples was carried out on a spark plasma sintering unit CY-SPS-T20 (Zhengzhou CY Scientific Instrument Co., Zhengzhou, China) under a pressure of up to 20 tons and 2000 °C, with holding time of 7 min (Figure 7). The spark plasma sintering system has a three-phase power supply 380 V AC, 50/60 Hz, 0–1500 A DC (digitally controlled), and 0–10 V DC (digitally controlled) built-in precision thermostat. At high heating rate, the error is less than 5 °C. It has temperature accuracy of <0.1 °C. Hydraulic automatic press has maximum load of 20 tons and built-in digital pressure gauge with overpressure alarm.

3. Results and Discussion

A series of samples consisting of MnO_nano 5.5 wt. % + Al₂O_3nano 2.0 wt. % + SiC_nm (37–57 wt. %) + SiC_µm (31–51 wt. %) + SiO_2µm 4.5 wt. % was obtained by SPS, which provides for the preservation of the nanostructure in ceramics. The powder, placed in a graphite mold, was subjected to electric current pulses, which caused the appearance of plasma particles at the contacts. As a result, the entire sample was rapidly heated up to 2000 °C. At the same time, a pressure of 20 tons was applied to the powder through the punch for several minutes.

In this study, the charge for sintering was prepared by multistage methods, including processing in the impeller-type reactor, which protects them from aggregation. Figure 8 shows SEM images of charges S-1, S-2, and S-3. Even at a high magnification of 5000 times, no agglomeration of powder components is observed, and there are no large group accumulations.

The microstructure of the studied ceramics consists mainly of small equiaxed SiC grains with rare inclusions of large ones (Figure 9). Sample S-1 (Figure 9a) with maximum density has characteristic elongated inclusions with a more monolithic (dense) structure. A more detailed study of sample S-1 at higher magnification shows that the microstructure of the sintered sample is characterized by equiaxed grains; the morphology of SiC grains is mainly represented by grains of a regular elongated oval-shaped configuration. The microstructure is characterized by the presence of equiaxed grains, the average size of which is 1 µm. Intercrystalline pores are also observed, predominantly irregular in shape and less than 1 μm in size. There are no defects in the form of extended voids in the bulk of the material.

This nature of the microstructure makes an assumption about the uniform distribution of pressure along the height of the compact, which allows the powder components to collapse; there are no separate voids and extended voids. This suggests that the stage of charge preparation, molding, and sintering is at a sufficient level to complete the material compaction process with a given sintering temperature.

Such a microstructure with elongated oval-shaped grains was obtained on the basis of micro- and nanopowders of the SiC matrix in an average ratio of nanoSiC and microSiC 2:1, and here, a phase transformation takes place. The introduction of SiC nanoparticles into the composition of the material leads to the formation of a specific structure of the material according to the “composite in composite” type. On the one hand, the composite is a dispersion-strengthened nanocomposite of the “micro/nano” type. On the other hand, all ceramic material containing 57 wt. % nanoSiC and 31 wt. % microSiC is composite. Such a complex structure of the composite leads to a significant improvement in the physical and mechanical properties of the material. The spark plasma sintering method makes it possible to synthesize powders that are most active in sintering, since the process proceeds under the most nonequilibrium conditions. Obviously, for the phase transition to occur, it is necessary to have a certain amount of hexagonal silicon carbide (SiC-6H) in the system as a “seed”.

Diffraction patterns after sintering in comparison with silicon carbide nanopowder are shown in Figure 10. As can be seen from the figure, SiC undergoes phase transformations during sintering above 1800 °C. The components during the sintering process mainly pass into the SiC₂ phase in certain ratios, which directly depend on the temperature. The intensity of the main lines of SiC decreases, and peaks of SiC₂ appear. In a complex composite system (MnO, Al₂O₃, SiC, and SiO₂), during sintering and simultaneous pressure, the formation of SiC₂ phases occurs as a result of the interaction of the initial components. Since SiC has a large number of polytypes, the formation of the SiC₂ phase is the result of polymorphism, which is confirmed by the results of many studies during the sintering of SiC powders with oxide additives. In addition, obtaining new silicon carbide phases is possible from SiC clusters, taking into account the dispersion of all constituent components. As mechanisms for obtaining polymorphic varieties of SiC, the crosslinking of precursor nanostructures or the combination of atoms of precursor nanostructures is used. The diffraction pattern shows narrow, intense diffraction peaks, which indicates a well-crystallized and homogeneous phase.

SEM images and the results of elemental analysis of the studied samples are shown in Figure 11. Elemental analysis showed that the samples contain Si, Mn, and Al. The elemental analysis data are in good agreement with the results of the XRD analysis. Sample S1 (Figure 11a) at a higher magnification of 5000 times demonstrates a homogeneous structure.

Based on the results of the study, experimental data on the thermal stability of ceramic samples were obtained. As a result of the experiments, the dependences of the change in mass, heat flux, and temperature of the studied samples, as well as the dependences of the change in the partial pressures of gases during the experiment, were obtained (Figure 12, Figure 13 and Figure 14).

As a result of studying the processes of thermal stability and synthesis of various compounds, it can be noted that the partial composition of the gaseous medium in the reaction chamber with the sample under study did not change significantly throughout the entire time of heating the sample. Analysis of the obtained results showed that the maximum weight loss was registered for samples S-2 and S-3 (0.40% and 0.55% of the total initial weight, respectively). For sample S-1, the weight loss was 0.29%. The analysis of thermal stability in a nitrogen atmosphere for all three selected compositions depends on the quantitative content of the samples. The sample S-1 with the maximum density and mechanical strength shows the maximum thermal stability. With an increase in the pressure of the gas medium, the temperature of the beginning of decomposition of all samples shifts to the region of high temperatures by approximately 400 °C. Figure 12, Figure 13 and Figure 14 show the influence of the pressure of the gaseous medium (N₂) on the temperature at the beginning and end of the decomposition of the SiC ceramic composition. It was found that for all three selected samples, the temperature range from the beginning to the end of decomposition is very narrow. At the maximum heating temperature of 1500 °C, hydrocarbon compounds (M45–M49) are actively released. The shape of the weight loss curves for samples S1 and S3 (Figure 12b and Figure 14b) has the same character, monotonically decreasing in proportion to the temperature level, with a maximum loss from 1450 °C. Sample S-2 in Figure 13b has a different form of the mass loss curve; at the maximum heating temperature of 1500 °C, the mass of the sample shows an increase by an average of 12% from samples S-1 and S-3 at the same temperature, while in sample S-2 is actively emitting carbon dioxide (CO₂). It is likely that the components in this composition of 5.5 wt. % MnO_nano + 2.0 wt. % Al₂O_3nano + 47 wt. % SiC_nano + 41 wt. % SiC_micro + 4.5 wt. % SiO_2micro begin to decompose at high temperatures. The thermal stability of ceramic compositions shifts on average by ∼200 °C to high temperatures with an increase in the pressure of the gas medium.

Figure 15 shows the processed heat flux change curves. Figure 16 shows a general view of the tested samples with lost masses (

Table 2. Weight loss of the studied samples during the experiment.

Sample	Initial Mass, mg	Final Mass, mg	Weight Loss, mg	Weight Loss, %
S-1	49.9798 ± 0.0001	49.7819 ± 0.0001	0.1979	0.40
S-2	50.0100 ± 0.0001	49.7334 ± 0.0001	0.2766	0.55
S-3	52.2000 ± 0.0001	52.0487 ± 0.0001	0.1513	0.29

).

The obtained data on the main physico-mechanical properties of ceramic samples in comparison with the data obtained earlier by the method of liquid-phase sintering are shown in

Table 3. Physico-mechanical properties of SiC ceramic samples obtained by liquid-phase sintering [15] and SPS method.

Name of the Sample	Synthesis Conditions	Mechanical Strength, MPa	Microhardness (HV), GPa	Apparent Density, g/cm3	Open Porosity, %
2.5 wt. % MnOnano + 2.0 wt. % Al2O3nano + 91 wt. % SiCµm + 4.5 wt. % SiO2µm	LPS, 1800 °C, 5000 t	440 ± 9	30 ± 1	3.0 ± 0.1	18.2 ± 0.1
5.5 wt. % MnOnano + 2.0 wt. % Al2O3nano + 57 wt. % SiCnano + 31 wt. % SiCµm + 4.5 wt. % SiO2µm (S-1)	SPS, 2000 °C, 20 t	457 ± 9	38 ± 1	3.3 ± 0.1	16.1 ± 0.1
5.5 wt. % MnOnano + 2.0 wt. % Al2O3nano + 47 wt. % SiCnano + 41 wt. % SiCµm + 4.5 wt. % SiO2µm (S-2)	SPS, 2000 °C, 20 t	450 ± 9	31 ± 1	3.1 ± 0.1	17.0 ± 0.1
5.5 wt. % MnOnano + 2.0 wt. % Al2O3nano + 37 wt. % SiCnano + 51 wt. % SiCµm + 4.5 wt. % SiO2µm (S-3)	SPS, 2000 °C, 20 t	454 ± 9	35 ± 1	3.2 ± 0.1	16.9 ± 0.1

.

The mechanical properties of the synthesized ceramics are greatly influenced by the manufacturing method, even with a pressing force of 20 t. Samples obtained by SPS (S-1, S-2, and S-3) show a higher density of the material at the level of 3.3 g/cm³ (S-3) compared with samples with a pressing force of 5000 tons by liquid-phase sintering. In addition, indicators of mechanical strength and hardness of ceramics increase.

Higher density values indicate the completion of the sintering process at the selected silicon carbide synthesis temperature. Also, the indices of mechanical strength of 457 MPa and microhardness of ceramics of 38 GPa (S-1), obtained by SPS, increase in comparison with LPS 440 MPa and 30 GPa, respectively.

In addition, from the analysis of the results obtained on the main physical and mechanical characteristics of the manufactured ceramic samples, it was revealed that composite ceramics based on silicon carbide has a number of properties that make them suitable for use, including in armor protection: high hardness, strength, density, and heat resistance. The density is 2.5 times lower than that of steel. Dissipation of the energy of a high-speed bullet leads to blunting of its tip and the crushing and fragmentation of ceramics at the point of impact. Therefore, the following properties of the armor material are extremely important: fracture toughness, hardness, and elasticity.

Table 4. Physico-mechanical characteristics of the sample S-1.

Sample	Elastic Modulus E, GPa	Fracture Toughness KIc, MPa∙m1/2
S-1	328 ± 2	3.6 ± 0.1

presents the main results of the physical and mechanical properties of the ceramic sample S-1, which demonstrates the highest performance of the samples.

The speed of the shock wave in the armored material depends on its density and elastic modulus. Often, very brittle materials such as SiC or glass ceramics provide a better shielding effect from simple shock wave propagation calculations than would be expected. The value of the elastic modulus of sample S-1 was 328 GPa, which on average, corresponds to the data obtained for 300–440 GPa. In addition, the high speed of propagation of a wave is important for two reasons: the impact energy is quickly dissipated over a large area, and the degree of destruction of the bullet itself increases. The greater the difference in the speed of propagation of a wave in the material of the bullet and in the armored material, the more the bullet will be destroyed.

Microhardness is also important in determining the initial destruction of the bullet; for the studied samples of materials, the microhardness varies within 30–38 GPa and depends on the hardness of the phases included in their composition and the content of oxides, which also corresponds to the average level of microhardness. Fracture toughness, in this study, is represented by the stress intensity factor K_Ic and is 3.6 MPa∙m^1/2, which is above the average value. According to the results of the study of the physico-mechanical properties of sample S-1, this composition and method for obtaining silicon carbide ceramics can be recommended for use, including in the form of an armor element.

4. Conclusions

Samples of composite ceramics (MnO_nano 5.5 wt. % + Al₂O_3nano 2.0 wt. % + SiC_nm (37–57 wt. %) + SiC_µm (31–51 wt. %) + SiO_2µm 4.5 wt. %) were obtained by the SPS method. The effect of additives on the hardening and structure of the resulting composites was studied. The studies carried out are relevant from the point of view of obtaining materials of a granular structure based on SiC while maintaining high thermomechanical properties.

It is shown that the method of spark plasma sintering, as compared to liquid-phase sintering with eutectic additives, has advantages in obtaining ceramic samples with higher characteristics. This made it possible to obtain SiC materials with a higher density (by 8%) and almost no significant crystal growth compared to samples obtained by liquid-phase sintering. Thus, the advantage of the SPS process is the short process time, which leads to minimal grain growth in the sintered ceramic.

The combined use of oxide eutectic additives and nanodispersed silicon carbide powder makes it possible to create ceramic materials with improved physico-mechanical characteristics. Nanodispersed silicon carbide powder facilitates the obtaining of a dense packing of particles during the formation of samples, which gives the sintered material the highest strength and hardness.

The obtained results on the density of sample S1 in comparison with the studied samples indicate the degree of completion of the sintering process of the studied samples. According to the obtained results of the physico-mechanical properties of ceramic samples, the optimal proportion of components from a series of three concentrations is sample S-1 with a composition of 5.5 wt. % MnO_nano + 2.0 wt. % Al₂O_3nano + 57 wt. % SiC_nano + 31 wt. % SiC_micro + 4.5 wt. % SiO_2micro with a maximum strength of 457 MPa and density of 3.3 g/cm³.

Sample S-1, with the maximum density, has characteristic elongated inclusions with a more monolithic (dense) structure. A more detailed study of sample S-1 at higher magnification shows that the microstructure of the sintered sample is characterized by equiaxed grains. The morphology of SiC grains is mainly represented by grains of a regular elongated oval-shaped configuration, the average size of which is 1 μm. Intercrystalline pores are also observed, predominantly irregular in shape and less than 1 μm in size. There are no defects in the form of extended voids in the bulk of the material.

The combined use of oxide eutectic additives and nanodispersed silicon carbide powder makes it possible to create ceramic materials with improved physico-mechanical characteristics. Nanodispersed silicon carbide powder makes it possible to obtain a dense packing of particles during the formation of samples, which gives the sintered material the highest strength and hardness.

The results of the study made it possible to achieve a combination of new approaches to the design of compositions and technology for the manufacture of SiC ceramics, which significantly expands their areas of application. Analysis of the obtained results of thermal stability showed that the maximum weight loss was registered for samples S-2 and S-3. For sample S-1, the weight loss was 0.29%. The evaluation of the physico-mechanical properties of the obtained samples of silicon carbide suggests that they can be used in various industries, for example, in parts of internal combustion engines and gas turbine engines, heat exchangers operating in an aggressive environment, heaters for obtaining temperatures in furnaces in the range of 1400–1500 °C in the air, as well as armor material.

According to the results of the study of the physico-mechanical properties of sample S-1, this composition and method for obtaining silicon carbide ceramics can be recommended for use, including in the form of an armor element. The exclusion from the technology of expensive oxides of rare earth metals will reduce the cost of the synthesized material.

At the same time, despite the obtained results, some issues require further study and clarification. The search for inexpensive oxides for the most complete sintering of SiC will continue in terms of good wetting of silicon carbide by oxide-containing melts. The introduction of such oxides into the composition of the silicon carbide charge leads to a significant decrease in the sintering temperature of ceramics of various compositions, including those containing aluminum oxide. Therefore, it is necessary to clarify the feasibility of introducing such additives in terms of the price–quality level of sintering temperature reduction. In addition, all these further questions will significantly improve the strength and ductile fracture behavior of ceramics.