INTRODUCTION

Ceramics based on SiC exhibit substantial mechanical strength at high temperatures, as well as wear resistance, low coefficient of thermal expansion, resistance to oxidation at temperatures up to 1500°C, chemical inertness, biocompatibility, corrosion resistance, resistance to radiation effects, and good hardness and heat conductivity parameters [1, 2]. To obtain products with high mechanical properties at elevated temperatures, it is necessary to create nonporous nanocrystalline materials based on SiC [3, 4]. The properties of nanoceramics can substantially differ from the characteristics of traditional ceramics with micrometer-sized grains, and the behavior of characteristics with an increase in the grain size is individual and depends both on the physical nature of the investigated property and on the physicochemical characteristics of the ceramics used. The synthesis of ceramics based on silicon carbide requires high temperatures. Sintering without impurities at temperatures of 2150–2200°C does not allow one to obtain nonporous products both because of low diffusion coefficients and high volatility of SiC at these temperatures [5, 6].

Despite the existence of a wide range of ceramic materials for various purposes, which are developed on the basis of silicon carbide, the high energy and resource consumption reqired for their production, as well as the growing requirements of scientific and technical progress, force the search for ways to create new types of materials with a high level of physical and mechanical characteristics, but a lower sintering temperature. The use of various nanopowders as modifiers is among the main ways to solve these problems [7]. On the other hand, the use of hot pressing with transmission of a high-amperage current through a graphite mold (electroconsolidation) makes it possible to quickly heat the ceramic mixture, which prevents the growth of grains and activates the sintering process [8, 9].

It was established that the powder consolidation rates are substantially increased compared to the rate of traditional sintering both in the case of electric sintering and in the case of microwave sintering [10, 11]. Moreover, consolidation rates can increase by several orders of magnitude for high-intensity processes with high heating rates. This leads to complete compaction of the powders in a very short time while preserving the nanoscale internal structure. Electroconsolidation (electrosintering) makes it possible to obtain consolidated ceramic materials, such as Al2O3, ZrO2, TiC, and WC, without impurities and with minimum grain growth in about 10 min [12, 13], while traditional sintering requires several hours and special additives that deteriorate the properties of the material. At the same time, both electrically conductive and nonconductive powders are electroconsolidated with equally high success (using electrically conductive graphite molds) [14], and microwave sintering is successfully used for sintering both the ceramics and metals [15]. This confirms the urgent need for fundamental research on the consolidation of nanostructured bodies activated by an electromagnetic field, primarily through theoretical analysis and model experiments, which are rarely used now in this field of materials science.

EXPERIMENTAL

The submicrometer SiC powder (0.1–0.3 μm) obtained by the method of self-propagating high-temperature synthesis (SHS) [16] was used as a starting material (Fig. 1).

Fig. 1.
figure 1

Microstructures of (a) the initial SiC mixture, (b) the mixture after 90 min of mechanical activation, and (c) the SBC product.

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We used ZrO2 nanopowders partially stabilized with 3 mol % Y2O3 (Fig. 2) from the Donetsk Institute of Physics and Engineering, National Academy of Sciences of Ukraine, and α-Al2O3 nanopowders with a grain size of 30–50 nm (Fig. 3) from the NANOE Company (France). The characteristics of α-Al2O3 nanoparticles are given in Table 1.

Fig. 2.
figure 2

Microphotograph of (a) ZrO2 nanoparticles and their size distribution (b).

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Fig. 3.
figure 3

Microphotograph of α-Al2O3 nanoparticles.

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Table 1.   Characteristics of α-Al2O3 nanoparticles
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A unique Activator 2S high-energy mill was used for the preparation of powder mixtures. The energy saturation of the grinding process was 3 W/h.

The microstructure was studied by scanning electron microscopy [17, 18]. The X-ray diffraction spectra were taken on an automated DRON-4.0-07 diffractometer with CoKα and CuKα radiation sources. The recording was performed in the step scanning mode within the angle range of 20°–120° with a recording step of 0.1° and an exposure time of 4 s. Nanoindentation was performed with a three-sided Berkovich indenter on a Nano Indenter G200 nanohardness tester manufactured by Agilent Technologies (USA); the indentation depth was 200 nm. The tests were carried out on the sample matrix and inclusions. Figure 4 shows impressions from the indenter after tests on the matrix and inclusions of the 20SiCmicro–80Al2O3 nanosample.Footnote 1

Fig. 4.
figure 4

Loading curves in (a) the matrix and (b) inclusions for 20SiCmicro–Al2O3 nano.

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At least ten imprints were made on each sample at a distance of 15 μm on inclusions and 25 μm from each other on the matrix to avoid the influence of neighboring imprints on the test results, and then the obtained data were averaged. To find the hardness and elastic modulus, the method of Oliver and Pharr [19] was used. A sample of fused quartz was tested before starting the tests of the samples, which serves as a hardness standard during nanoindentation, since it does not have a large-scale effect and deformation hardening of the surface during its mechanical processing. The obtained data for fused quartz (E = 72 GPa and H = 9.2 GPa) confirm that the instrument is reliably calibrated.

The microhardness and crack resistance values were determined by measuring the imprint diagonal and the length of the radial cracks obtained during indentation of a diamond indenter in the form of a four-sided pyramid with an angle of α = 136° near the tip (Vickers pyramid) on a NEXUS 4504 automatic microhardness tester. The load was P = 10 N, and the exposure time was 10 s.

The microhardness determined was determined using the following expression [20]:

$${{H}_{V}} = \frac{{kP}}{{{{{\left( {2a} \right)}}^{2}}}},$$
(1)

where P is the load on the indenter, kg; 2a is the average value of the lengths of both diagonals of the imprint, μm; k is a coefficient that depends on the shape of the indenter (for the Vickers pyramid, k = 1.854).

Fracture viscosity coefficient KIc, which characterizes the crack resistance of the sample, was determined by the following expression [21, 22]

$${{K}_{{{\text{I}}c}}} = \zeta {{\left( {\frac{l}{a}} \right)}^{{ - 0.5}}}{{\left( {\frac{{{{H}_{V}}}}{{EF}}} \right)}^{{ - 0.4}}}\frac{{{{H}_{V}}{{a}^{{0.5}}}}}{F},$$
(2)

where ζ is a dimensionless constant (for ceramics, it has an average value of 0.016); E is the Young modulus, GPa; HV is the hardness, GPa; F is a constant (F ≈ 3); l is the length of the crack from the corner of the imprint of the Vickers pyramid, m; a is the semidiagonal of the imprint of the Vickers pyramid, m.

The following conditions were fulfilled during the calculation:

$$0.25 \leqslant \frac{l}{a} \leqslant 2.5.$$
(3)

The strength limit during bending of the sample under conditions of three-point bending was determined by the following expression [23]:

$${{\sigma }_{{{\text{flex}}}}} = \frac{{3PL}}{{2b{{h}^{2}}}},$$
(4)

where P is the force at the moment of dividing the sample into parts, N; L is the sample length, mm; b is the sample width, mm; h is the sample thickness in the direction parallel to the force applied to the sample, mm.

The heat conductivity was measured using the stationary heat flow method [24]. This method makes it possible to measure the temperature dependence of the heat conductivity in a wide temperature range. A rod whose length was several times larger than the cross size was cut from the finished compressed sample. The resulting rod was placed between the receiver and the heat source, due to which a longitudinal heat flow and a temperature gradient were created in it. The installation was comprised of a cryostat, i.e., a hermetically sealed cell with a sample placed in the container with liquid helium, which was immersed, in turn, in liquid nitrogen. To reduce heat losses, the air from the chamber was evacuated to a pressure of about 10–7 mm Hg. Inside the chamber, the sample under test was placed in the thermostat, to which a heat source (heater) and a differential copper–constantan thermocouple were attached. The temperature was measured with a platinum resistance thermometer. With small differences in the mounting of the sample in the cryostat, it was possible to measure not only the heat conductivity, but also other transport characteristics of the solid body, such as the thermal diffusivity and electrical conductivity.

Silicon carbide refers to compounds that have strong covalent bonds, which complicates mass transfer during sintering without the use of activating impurities and the application of external pressure [25]. To obtain dense SiC ceramics, the method of liquid phase sintering with use of oxide impurities is typically employed [26, 27]. The most important conditions of this method are a decrease in the temperature of the appearance of the liquid phase under the condition of its minimum viscosity and good wettability of the surface of silicon carbide particles [28]. High density indicators were obtained with an oxide impurity content of 10 wt %, and the maximum shrinkage of the samples was observed with an impurity content of 20 wt %.

RESULTS AND DISCUSSION

In the process of heating during electroconsolidation, the following reaction occurs between SiC and oxides:

$${\text{Si}}{{{\text{C}}}_{{{\text{(solid)}}}}}{\text{ + A}}{{{\text{l}}}_{{\text{2}}}}{{{\text{O}}}_{{{\text{3(solid)}}}}} \to {\text{Si}}{{{\text{O}}}_{{{\text{(gas)}}}}}{\text{ + A}}{{{\text{l}}}_{{\text{2}}}}{{{\text{O}}}_{{\text{3}}}}_{{{\text{(gas)}}}}{\text{ + C}}{{{\text{O}}}_{{{\text{(gas)}}}}}{\text{.}}$$
(5)

Therefore, it is necessary to take into account the release of CO gases in the process of hot pressing; for this, it is necessary to hold the reaction in the temperature range of 800–1000°C for up to 10 min in order to release CO gases.

To form a liquid phase, Al2O3 and ZrO2 powders were used as impurities, the mixture of which has a melting temperature of 1860°C and form a liquid phase when mixed in the ratio of 60 and 40 mol %, respectively [29, 30]. Figure 5 shows the structure of SiC ceramics with an admixture of aluminum oxide.

Fig. 5.
figure 5

Structure of hot-pressed SiC–50Al2O3 at a temperature of T = 1600°C, a pressure of 30 MPa, and a residing time of 3 min.

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To determine the optimal content of the admixture of an aluminum oxide nanopowder, the composition with the same content of both components was first studied. The sintering temperature was empirically chosen based on the possibility of obtaining the densest samples. Figure 6 shows the phase diagram of the Al2O3–ZrO2 system.

Fig. 6.
figure 6

Phase diagram of the Al2O3–ZrO2 system.

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The maximum density was obtained with the silicon carbide, aluminum oxide, and zirconium oxide contents corresponding to the eutectic composition (Table 2).

Table 2.   Density of initial mixtures with different compositions
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A study of the physical and mechanical properties of SiC-based materials obtained by liquid-phase sintering with two activating impurities of aluminum oxide and zirconium dioxide nanopowders showed that the density in the case of sintering of the eutectic composition at a temperature of 1860 °C is almost 100%. With a different composition of the initial mixture and in the case of sintering at a temperature below 1860°C, incomplete wetting of the SiC surface by oxides occurs and, as a result, the density of the composite and the level of mechanical properties decrease (Fig. 7).

Fig. 7.
figure 7

Fractogram of the hot-pressed initial SiC–12ZrO2–8Al2O3 mixture at a temperature of 1700°C, a pressure of 30 MPa, and a residing time of 3 min.

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Samples of products based on SiC micropowders with an admixture of aluminum oxide nanopowders had the following maximum mechanical properties:

σflex = 560 ± 10 MPa;

HV = 22.3 ± 0.2 GPa;

KIc = 5.5 ± 0.2 MPa m1/2.

The result of the X-ray phase analysis of one of the samples obtained from a mixture of starting powders (see sample 7 in Table 2) is shown in Fig. 8.

Fig. 8.
figure 8

X-ray diffractograms of SiC–12ZrO2–8Al2O3 ceramics sintered at a temperature of 1860°C, a pressure of 30 MPa, and a residing time of 3 min.

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The composite material obtained using liquid-phase pressing (Fig. 7), has the following physical and mechanical properties:

— flexural strength limit 800–1000 MPa;

— crack resistance 6–8 MPa m1/2;

— hardness 91–93 HRA;

— heat conductivity coefficient 30–35 W/(m K).

In fact, the mechanical strength of ceramics is determined by the physical and mechanical properties of the homogeneous eutectic medium–(silicon carbide micropowder + nanodispersed aluminum and zirconium oxides) system, which itself is a ceramic nanocomposite of the micro/nano type (Fig. 9). In this case, the mechanism of destruction of the material changes fundamentally. The crack passes through a homogeneous eutectic medium, and the intercrystalline fracture actually turns into a transcrystalline fracture.

Fig. 9.
figure 9

Structure of the material of the micro/nano type.

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Moreover, the high strength of the interface with the matrix provides a substantial shielding stress and a sharp increase in both the strength and crack resistance when the dispersed phase is represented by a brittle material, the destruction of which occurs even in the event of a small opening of the crack.

Figure 10 shows the fractogram of the sample obtained on the basis of initial mixtures with an eutectic composition.

Fig. 10.
figure 10

Fractogram of the hot-pressed composite from the initial 80SiC + 20(Al2O3 + ZrO2) mixture at a temperature of 1860°C, a pressure of 30 MPa, and a residing time of 3 min.

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It should be noted that quite interesting results were obtained when the Al2O3 nano–20SiCmicro nanopowder was used as a basis of the composite material.

In this case, the nanohardness and elastic modulus of individual Al2O3 and SiC phases were investigated. The fairly high hardness and elastic modulus values allow us to assume that the resulting composite material will exhibit high wear resistance (Fig. 11, Table 3).

Fig. 11.
figure 11

Spectral analysis of the composite material obtained from the initial Al2O3 nano–20SiC mixture by hot pressing in vacuum at a temperature of 1600°C, a pressure of 30 MPa, and a residing time of 3 min.

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Table 3.   Density of initial mixtures with different compositions
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As can be seen from Fig. 11, agglomerates of aluminum oxide nanopowder are visible in points 2 and 3, which is a result of insufficiently good mixing of the original powders. Nanohardness and elastic modulus measurements of individual phases of the same sample show a large difference in the nanohardness and elastic modulus of the individual phases.

Fractograms and phase characteristics of silicon carbide (Fig. 12) and aluminum oxide (Fig. 13) are given for the 20SiCmicro–80Al2O3 nano composite sintered at a temperature of 1700°C under a pressure of 30 MPa with an exposure time of 3 min.

Fig. 12.
figure 12

(a) Fractogram and (b) characteristics of the silicon carbide phase with a nanohardness of E = 446 ± 15 GPa and an elastic modulus of Н = 30.6 ± 2 GPa.

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Fig. 13.
figure 13

(a) Fractogram and (b) characteristics of the aluminum oxide phase with a nanohardness of E = 348 ± 29 GPa and an elastic modulus of Н = 19.5 ± 3 GPa.

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Thus, it can be assumed that the combination of a matrix of micropowder grains and admixtures of nanopowders of refractory oxides in the initial mixtures allows one to obtain composite materials with high hardness, crack resistance, and elastic modulus charcateristics. However, it is necessary to investigate how heating to temperatures above 1000°C will affect the mechanical properties of composite materials based on silicon carbide. This is very important if the material is intended to be used as a thermal abrasive nozzle or under other heavy-duty conditions.

CONCLUSIONS

The samples prepared using the 80SiC + 20(Al2O3 + ZrO2) eutectic compositions have the highest hardness and crack resistance characteristics. It was determined that the optimal modes of sintering of eutectic composition mixtures are achieved at a temperature of 1860°C, a pressure of 30 MPa, and an exposure time of 3 min. Samples obtained from mixtures of the Al2O3 nano–SiCmicro system by hot pressing at a temperature of 1600°C also have high values of the hardness and elastic modulus. High values of the hardness and elastic modulus cause high wear resistance of the obtained composite material, which will allow it to be used as a material for ceramic inserts of abrasive jet nozzles with a straight channel.

The performed experiments make it possible to state that a high solubility of the solid substance in the liquid phase is not required for high-quality compaction in the process of liquid-phase sintering. On the contrary, it should be avoided, since a high solubility can lead to deformation of the material and its subsequent destruction. Moreover, there is a problem of uniform distribution of modifiers over the surface of their particles in the technology of ceramics synthesis in the case of using nanodisperse powders as starting materials.

To solve this problem and improve the processes of controlling the structure of materials based on silicon carbide, which contain impurities of eutectic compositions, it is advisable to use the method of heterophase chemical deposition directly on silicon carbide particles. This will make it possible to obtain a more perfect microstructure, to effectively carry out the high-temperature compaction process, and to reduce the probability of the reaction of SiC with the melt, which leads to the formation of gaseous products.