Microstructure and Mechanical Properties of Si3N4-Fe3Si Composites Prepared by Gas-Pressure Sintering

Si3N4-Fe3Si composites were prepared using Fe-Si3N4 as the source of Fe3Si by gas-pressure sintering. By adding different amounts of Fe-Si3N4 into the starting powders, Si3N4-Fe3Si composites with various Fe3Si phase contents were obtained. The microstructure and mechanical properties of the composites were investigated. With the increase of Fe-Si3N4 contents, the content and particle size of Fe3Si both increased. When more than 60 wt. % Fe-Si3N4 were added, the abnormal growth of Fe3Si particles occurred and oversized Fe3Si particles appeared, leading to non-uniform microstructures and worse mechanical properties of the composites. It has been found that Fe3Si particles could toughen the composites through particle pull-out, interface debonding, crack deflection, and particle bridging. Uniform microstructure and improved mechanical properties (flexural strength of 354 MPa and fracture toughness of 8.4 MPa·m1/2) can be achieved for FSN40.


Introduction
Si 3 N 4 is one of the most promising engineering ceramics with high strength, high hardness, good oxidation, and corrosion resistance-even at high temperature. It is also an irreplaceable material in modern industries [1][2][3][4]. Numerous works have been done to lower the price of Si 3 N 4 products in order to realize its wide application. However, in some price-sensitive areas, like the Metallurgical industry, the cost of Si 3 N 4 is still too high. To tackle this problem, a novel refractory material, ferro-silicon nitride (Fe-Si 3 N 4 ) has been developed and has successfully replaced the relatively expensive Si 3 N 4 [5,6]. Using the ferro-silicon alloy FeSi75 as the raw material, Fe-Si 3 N 4 powder can be synthesized by direct nitridation [7][8][9] or self-propagating high-temperature synthesis [10][11][12][13] under nitrogen atmosphere. Owing to the catalytic effect of Fe, a lower nitridation temperature and higher reaction rate can be achieved [7,14,15]. FeSi75 is extensively used as a hardener and scavenger in smelting steel, so it can be easily obtained at a low price. And the synthesis process of Fe-Si 3 N 4 is concise, cost effective, and can be operated on a large scale. The above factors lead to the low cost of Fe-Si 3 N 4 products.
Microstructural analysis has shown that Fe-Si 3 N 4 was mainly composed of Si 3 N 4 (~75 wt. %) and a small amount of un-nitrided Fe 3 Si (~15 wt. %) [16,17]. With the main phase being Si 3 N 4 , Fe-Si 3 N 4 inherits the excellent comprehensive properties of pure Si 3 N 4 . What is more, the low melting point Fe-containing phase endows Fe-Si 3 N 4 with good sinterability which resulted from enhanced particle rearrangement and diffusion in the presence of a more liquid phase [6].
Through the reaction bonding of FeSi75 powder compact or the sintering of Fe-Si 3 N 4 , porous Fe-Si 3 N 4 ceramics and Fe-Si 3 N 4 based ceramic composites for the refractory application

Materials and Methods
Fe-Si 3 N 4 powder (Fe content: 12~20 wt. %; Xi'an Aoqin new materials Co., Ltd., Xi'an, China), Si 3 N 4 powder (purity > 99.9%, 0.8 µm, Shanghai Shuitian technology Co., LTD., Shanghai, China), Y 2 O 3 powder (purity >99.9%, 1 µm, Shanghai Shuitian technology Co., LTD., Shanghai, China), and Al 2 O 3 powder (purity >99.9%, 1 µm, Shanghai Shuitian technology Co., LTD., Shanghai, China) were used as starting powders. Y 2 O 3 and Al 2 O 3 act as sintering aids. Under high temperature, Y 2 O 3 and Al 2 O 3 react with SiO 2 or silicon oxynitride, which are always present on the surfaces of Si 3 N 4 powders, to form a liquid phase that is beneficial for densification. The size distribution of the starting powders are given in Figure 1. In order to adjust the content of Fe 3 Si in the final composites of the Si 3 N 4 -Fe 3 Si composites, different ratios between Fe-Si 3 N 4 and Si 3 N 4 powders were used. The compositions of the starting powders are shown in Table 1. The powders were balled milled in alcohol using a planetary mill for 24 h at a rotating speed of 300 rpm. After drying and sieving, the powders were pressed uniaxial in a stainless-steel die at a pressure of 70 MPa and were then cold-isostatically pressed at a pressure of 200 MPa. The sintering of the green bodies was conducted under a nitrogen pressure of 10 MPa at 1800 • C for 2 h. and Fe-Si3N4-ZrO2 [20] refractories. Results showed that, compared with their traditional rivals, Fe-Si3N4 based refractories exhibit high thermal strength, a higher coefficient of thermal conductivity, and better thermal shock resistance [19].
As the above application of Fe-Si3N4 has attracted much attention, the potential of Fe-Si3N4, which can be regarded as Si3N4-Fe3Si composite ceramics, as a thermal structural material has not been researched yet. For a long time, intermetallic compounds have been added to ceramics materials to toughen the brittle ceramics [21][22][23][24]. For instance, using Ni3Al intermetallic compound as a second phase, an Al2O3 composite with high strength and high toughness can be prepared [25]. Toughening effects that result from intermetallic particles can be attributed to crack deflection, crack bridging, plastic deformation [26,27], and residue stress caused by a coefficient of thermal expansion (CTE) mismatch [28]. Among the component of Fe-Si3N4, not only Si3N4, but also Fe3Si-which is an intermetallic compound-have good thermal mechanical properties [29,30]. So, one can infer that with both of its main phases having a high performance, the mechanical properties of Fe-Si3N4 are worth studying.
In this work, Si3N4-Fe3Si composites that contain different amount of Fe3Si with high relative density are successfully prepared by gas-pressure sintering. The microstructure and mechanical properties are studied and the effect of Fe3Si particles on crack propagation behavior is highlighted.

Materials and Methods
Fe-Si3N4 powder (Fe content: 12~20 wt. %; Xi'an Aoqin new materials Co., Ltd., Xi'an, China), Si3N4 powder (purity > 99.9%, 0.8 μm, Shanghai Shuitian technology Co., LTD., Shanghai, China), Y2O3 powder (purity >99.9%, 1 μm, Shanghai Shuitian technology Co., LTD., Shanghai, China), and Al2O3 powder (purity >99.9%, 1 μm, Shanghai Shuitian technology Co., LTD., Shanghai, China) were used as starting powders. Y2O3 and Al2O3 act as sintering aids. Under high temperature, Y2O3 and Al2O3 react with SiO2 or silicon oxynitride, which are always present on the surfaces of Si3N4 powders, to form a liquid phase that is beneficial for densification. The size distribution of the starting powders are given in Figure 1. In order to adjust the content of Fe3Si in the final composites of the Si3N4-Fe3Si composites, different ratios between Fe-Si3N4 and Si3N4 powders were used. The compositions of the starting powders are shown in Table 1. The powders were balled milled in alcohol using a planetary mill for 24 h at a rotating speed of 300 rpm. After drying and sieving, the powders were pressed uniaxial in a stainless-steel die at a pressure of 70 MPa and were then cold-isostatically pressed at a pressure of 200 MPa. The sintering of the green bodies was conducted under a nitrogen pressure of 10 MPa at 1800 °C for 2 h.   The bulk density and the open porosity of the sintered specimens were measured according to the Archimedes principle and can be calculated from the equation: where ρ is the bulk density, m 1 is the dry mass of the samples in air, m 2 is the mass of the specimen when fully impregnated with the water, and m 3 is the impregnated mass whilst suspended in the water. Phase compositions of the samples were identified by X-ray powder diffraction analysis (XRD, Rigaku-D/max-2400; Tokyo, Japan). Microstructures were observed by back scattered electron images (BSE, S-4700, Hitachi, Tokyo, Japan) on the polished surfaces so as to reveal the morphologies and distribution states of the β-Si 3 N 4 grains, the grain boundary phases, and the Fe 3 Si particles. Image analysis was conducted to determine the phase content and the particle size with image analysis software that analyzed ten different back scattered electron (BSE) images. The microstructures of the fracture surfaces of the ceramics were observed using SEM (S-4700, Hitachi, Tokyo, Japan). The elemental composition was analyzed with an energy dispersive X-ray spectrometer (EDS). Indentations were placed on the polished surfaces by a Vickers indenter with a load of 9.8 N holding for 15 s to measure the hardness of the sintered samples, and then the indented surfaces were observed by SEM to examine the crack/microstructure interactions. Flexural strengths of the composites were tested by three-point bending on bars that were 40 mm long, 4 mm wide, and 3 mm thick according to the ASTM-D790 standard, using a 30 mm span and a crosshead speed of 0.5 mm/min. Fracture toughness was evaluated by single-edge notched beam (SENB) according to the ASTM-C1421-01b standard at a span of 20 mm and a crosshead speed of 0.05 mm/min using bar samples that were 30 mm long, 2 mm wide, and 4 mm thick. Fracture toughness was also calculated by the Vickers indentation technique through the following equation: where K IC is the fracture toughness, H is the hardness, E is Young's modulus, φ is the constraint factor (~3), and c and a are half length of the diagonal of the indentation and the average crack length that was introduced by the indentation.

Results and Discussion
The density and open porosity of the Si 3 N 4 -Fe 3 Si composites are listed in Table 2. It can be found that the dense ceramic composites with high densities (≥3.2 g/cm 3 ) and low open porosities (≤2.06%) were prepared successfully by gas-pressure sintering. Figure 2 shows the X-ray diffraction patterns of the sintered samples. The crystalline phases that were identified in these spectrums include β-Si 3  It is obvious that in all of the samples, the main phases were β-Si 3 N 4 and no α-Si 3 N 4 could be detected, indicating a fully α → β phase transformation during the liquid phase sintering. Diffraction peaks around the 2θ of 29.4 • indicate the existence of a Y 2 Si 2 O 7 phase, which was formed at the grain junction of Si 3 N 4 by the reaction between sintering aid Y 2 O 3 and SiO 2 at the surface of Si 3 N 4 powder, which crystallized when it was cooled [31]. Studies have shown that the crystalline Y 2 Si 2 O 7 secondary phase is beneficial for high temperature oxidation resistance and thermal mechanical properties of Si 3 N 4 based materials [32,33]. Fe 3 Si was identified in all of the samples, suggesting that Fe 3 Si was stable at the sintering condition and was successfully introduced into the composites by adding Fe-Si 3 N 4 to the raw materials.   Figure 2 shows the X-ray diffraction patterns of the sintered samples. The crystalline phases that were identified in these spectrums include β-Si3N4 (ICDD PDF Card No. 33-1160), Y2Si2O7 (ICDD PDF Card No. 38-0440), Fe3Si (ICDD PDF Card No. 35-0519), and Al2O3 (ICDD PDF Card No. . It is obvious that in all of the samples, the main phases were β-Si3N4 and no α-Si3N4 could be detected, indicating a fully α → β phase transformation during the liquid phase sintering. Diffraction peaks around the 2θ of 29.4° indicate the existence of a Y2Si2O7 phase, which was formed at the grain junction of Si3N4 by the reaction between sintering aid Y2O3 and SiO2 at the surface of Si3N4 powder, which crystallized when it was cooled [31]. Studies have shown that the crystalline Y2Si2O7 secondary phase is beneficial for high temperature oxidation resistance and thermal mechanical properties of Si3N4 based materials [32,33]. Fe3Si was identified in all of the samples, suggesting that Fe3Si was stable at the sintering condition and was successfully introduced into the composites by adding Fe-Si3N4 to the raw materials. It has been found that ferrous metals (Fe, Ni, Co, etc.) have a high affinity for Si, so that SiC and Si3N4 are reactive to Fe and some of its alloys [34,35]. T. Shimoo and K. Okamura [36] studied the reactions between silicides of Fe and Si3N4. In their work, they found that the silicides with a low Si/metal ratio react with Si3N4 to produce those with a high Si/metal ratio. At 1250 °C under Ar atmosphere, FeSi can be generated through the following overall reaction: However, in our system, the sintering was conducted under a high N2 pressure (10 MPa). Furthermore, in the process of gas-pressure sintering, with an increasing temperature, the pores of the ceramic became closed pores so that the N2 that was produced through the above reaction could not be released, resulting in a local environment with an even higher N2 pressure which meant that It has been found that ferrous metals (Fe, Ni, Co, etc.) have a high affinity for Si, so that SiC and Si 3 N 4 are reactive to Fe and some of its alloys [34,35]. T. Shimoo and K. Okamura [36] studied the reactions between silicides of Fe and Si 3 N 4 . In their work, they found that the silicides with a low Si/metal ratio react with Si 3 N 4 to produce those with a high Si/metal ratio. At 1250 • C under Ar atmosphere, FeSi can be generated through the following overall reaction: However, in our system, the sintering was conducted under a high N 2 pressure (10 MPa). Furthermore, in the process of gas-pressure sintering, with an increasing temperature, the pores of the ceramic became closed pores so that the N 2 that was produced through the above reaction could not be released, resulting in a local environment with an even higher N 2 pressure which meant that the reaction could be suppressed, making Fe 3 Si stable to survive the sintering process. Thermodynamic analysis was done to calculate the Gibbs free energy of the reaction, and the results are present in Figure 3. It can be seen that with an increasing pressure of N 2 , the Gibbs free energy drops gradually. When N 2 pressure was above 13.7 MPa, the Gibbs free energy was greater than 0 kJ·mol −1 , indicating that the reaction could no longer continue. In our sintering condition, the external pressure of N 2 was 10 MPa, therefore when the reaction occured and N 2 was produced, the local pressure of N 2 was much higher than 10 MPa, meaning the reaction could be suppressed or even stopped.
Materials 2018, 11, x FOR PEER REVIEW 5 of 12 the reaction could be suppressed, making Fe3Si stable to survive the sintering process. Thermodynamic analysis was done to calculate the Gibbs free energy of the reaction, and the results are present in Figure 3. It can be seen that with an increasing pressure of N2, the Gibbs free energy drops gradually. When N2 pressure was above 13.7 MPa, the Gibbs free energy was greater than 0 kJ·mol -1 , indicating that the reaction could no longer continue. In our sintering condition, the external pressure of N2 was 10 MPa, therefore when the reaction occured and N2 was produced, the local pressure of N2 was much higher than 10 MPa, meaning the reaction could be suppressed or even stopped.  Figure 4 shows the BSE images and energy-dispersive X-ray spectroscopy (EDS) analysis of the composites. It can be seen from Figure 4a-f that all of the samples were mainly composed of three phases which are indicated by arrows in Figure 4f: the dark-gray columnar grains, the light-gray phases, and the white dispersive particles. EDS analysis (Table 3) confirmed that they are Si3N4, grain boundary phases, and Fe3Si, respectively. It can be clearly noted that, with the increase of Fe-Si3N4 content in the raw materials, the content of the Fe3Si phase in the composites arises. More importantly, the particle size of Fe3Si grows remarkably. In Figure 4a, the particle size of Fe3Si is smaller than 5 μm. However, in Figure 4e, Fe3Si particles that are bigger than 20 μm can be found. This phenomenon was confirmed by particle size measurements through image analysis ( Figure 5). In Figure 5a, it can be found that the volume content rose in approximate linearity with the increase of Fe-Si3N4 content. The volume fraction of Fe3Si for FSN20, FSN40, FSN 60, FSN80, and FSN90 after sintering are 0.7, 1.6, 2.5, 3.3, and 4.1 vol. %. Assuming that Fe-Si3N4 contains 18 wt. % Fe3Si and the density of Fe3Si is 6.34g/cm 3 [37], the content of Fe3Si for each sample before sintering are calculated and the results (see Table 1) are 1.9, 3.9, 5.9, 8.1, and 9.2 vol. % respectively. So, we can estimate that about 60% of Fe3Si are sacrificed. There may be two reasons for the loss of Fe3Si. Firstly, Fe3Si dissolves into the grain boundary phase and EDS results showed that Fe exists in it. Secondly, since our method was based on the image analysis of BSE images, Fe3Si particles that were too small to recognize would have been neglected.
The average particle size of Fe3Si increased gradually from 1.19 μm to 2.75 μm. Figure 5b gives the particle size distribution of Fe3Si, from which it can be discovered that with the increase of Fe-Si3N4, although the average particle size of Fe3Si rose mildly, the abnormal growth of the particles occurred, and the number of big Fe3Si particles grew rapidly. In samples of FSN80 and FSN90, although the average size of Fe3Si remains relatively small, many Fe3Si particles that were bigger than 10 μm can be found frequently, indicating that with the increase of Fe-Si3N4 content in the starting powder, the degree of microstructure inhomogeneity would rise.  Figure 4 shows the BSE images and energy-dispersive X-ray spectroscopy (EDS) analysis of the composites. It can be seen from Figure 4a-f that all of the samples were mainly composed of three phases which are indicated by arrows in Figure 4f: the dark-gray columnar grains, the light-gray phases, and the white dispersive particles. EDS analysis (Table 3) confirmed that they are Si 3 N 4 , grain boundary phases, and Fe 3 Si, respectively. It can be clearly noted that, with the increase of Fe-Si 3 N 4 content in the raw materials, the content of the Fe 3 Si phase in the composites arises. More importantly, the particle size of Fe 3 Si grows remarkably. In Figure 4a, the particle size of Fe 3 Si is smaller than 5 µm. However, in Figure 4e, Fe 3 Si particles that are bigger than 20 µm can be found. This phenomenon was confirmed by particle size measurements through image analysis ( Figure 5). In Figure 5a, it can be found that the volume content rose in approximate linearity with the increase of Fe-Si 3 N 4 content. The volume fraction of Fe 3 Si for FSN20, FSN40, FSN 60, FSN80, and FSN90 after sintering are 0.7, 1.6, 2.5, 3.3, and 4.1 vol. %. Assuming that Fe-Si 3 N 4 contains 18 wt. % Fe 3 Si and the density of Fe 3 Si is 6.34 g/cm 3 [37], the content of Fe 3 Si for each sample before sintering are calculated and the results (see Table 1) are 1.9, 3.9, 5.9, 8.1, and 9.2 vol. % respectively. So, we can estimate that about 60% of Fe 3 Si are sacrificed. There may be two reasons for the loss of Fe 3 Si. Firstly, Fe 3 Si dissolves into the grain boundary phase and EDS results showed that Fe exists in it. Secondly, since our method was based on the image analysis of BSE images, Fe 3 Si particles that were too small to recognize would have been neglected.
The average particle size of Fe 3 Si increased gradually from 1.19 µm to 2.75 µm. Figure 5b gives the particle size distribution of Fe 3 Si, from which it can be discovered that with the increase of Fe-Si 3 N 4 , although the average particle size of Fe 3 Si rose mildly, the abnormal growth of the particles occurred, and the number of big Fe 3 Si particles grew rapidly. In samples of FSN80 and FSN90, although the average size of Fe 3 Si remains relatively small, many Fe 3 Si particles that were bigger than 10 µm can be found frequently, indicating that with the increase of Fe-Si 3 N 4 content in the starting powder, the degree of microstructure inhomogeneity would rise. The difference in particle sizes and their distributions of Fe 3 Si in different samples can be explained by the flow of liquid Fe 3 Si in porous Si 3 N 4 during the sintering process. Since the melting point of Fe 3 Si is about 1280 • C [38], at sintering temperature, the Fe 3 Si is in liquid state and can flow easily. When the content of Fe-Si 3 N 4 was low and the Si 3 N 4 content was high, the frameworks that were formed by the Si 3 N 4 particles were relatively dense. Hence, the melting Fe 3 Si droplets were separated from each other and existed discretely. With the increase of the content of Fe-Si 3 N 4 and the decrease of the content of Si 3 N 4 , the Si 3 N 4 frameworks were weakened, and when the porosity of Si 3 N 4 and the content of Fe 3 Si reached a critical point, percolation occurred and Fe 3 Si droplets joined each other and flowed to form bigger droplets and solidified into solid particles upon cooling.      Figure  5b, 250 data of each particle size of each sample (represented by round dots with different colors) were selected randomly to illustrate the particle size distribution. Table 3. EDS analysis of spots in Figure 4f.  [38], at sintering temperature, the Fe3Si is in liquid state and can flow easily. When the content of Fe-Si3N4 was low and the Si3N4 content was high, the frameworks that were formed by the Si3N4 particles were relatively dense. Hence, the melting Fe3Si droplets were separated from each other and existed discretely. With the increase of the content of Fe-Si3N4 and the decrease of the content of Si3N4, the Si3N4 frameworks were weakened, and when the porosity of Si3N4 and the content of Fe3Si reached a critical point, percolation occurred and Fe3Si droplets joined each other and flowed to form bigger droplets and solidified into solid particles upon cooling.

Element Spot A Spot B Spot C
The mechanical properties of the Si3N4-Fe3Si composites are presented in Table 2 and Figure 6. From Table 2, we can see that the Vickers hardness of the samples decreased with the increase of Fe3Si In Figure 5b, 250 data of each particle size of each sample (represented by round dots with different colors) were selected randomly to illustrate the particle size distribution.
The mechanical properties of the Si 3 N 4 -Fe 3 Si composites are presented in Table 2 and Figure 6. From Table 2, we can see that the Vickers hardness of the samples decreased with the increase of Fe 3 Si content, which is as expected because the hardness of Fe 3 Si is much lower than that of Si 3 N 4 . Figure 6 shows the dependence of the flexural strength and fracture toughness of the Si 3 N 4 -Fe 3 Si composites on the content of Fe-Si 3 N 4 in starting powders. It can be seen that the flexural strength and fracture toughness of FSN20 are 293 MPa and 7.9 MPa·m 1/2 , respectively. The highest flexural strength and fracture toughness were obtained (354 MPa and 8.4 MPa·m 1/2 ) when the content of Fe-Si 3 N 4 increased to 40 wt. %. However, a further increase in Fe-Si 3 N 4 resulted in a gradual degradation of the mechanical properties of the composites. The above results showed that by carefully adjusting the composition of the raw materials, a Si 3 N 4 -Fe 3 Si composite with improved mechanical properties can be obtained. Dense monolithic Si 3 N 4 ceramics that are fabricated using Si 3 N 4 powder with high purity typically show good mechanical properties (three point bending strength of 400~900 MPa and fracture toughness of 3.4~8.2 MPa·m 1/2 [39]). Our results show that by replacing 40 wt. % Si 3 N 4 powder with the cost-effective Fe-Si 3 N 4 powder, composites can obtain mechanical properties-especially fracture toughness that is at the same level with dense monolithic Si 3 N 4 ceramics.
with the cost-effective Fe-Si3N4 powder, composites can obtain mechanical properties-especially fracture toughness that is at the same level with dense monolithic Si3N4 ceramics. In order to research the microstructure-the mechanical properties' relationship of the composites, fracture surface and crack/microstructure interactions were observed. Figure 7 shows the fracture surface of FSN40. In Figure 7a, it can be seen that fracture modes of the composites were in co-action by transgranular and intergranular fracture. Since Si3N4 has its well-known characteristic of self-reinforcement, two kinds of typical toughening mechanisms of crack bridging and crack deflection by Si3N4 can be found in Figure 7b, which resulted from the elongated Si3N4 crystals that were bounded by the weak interface [40]. Apart from the toughening mechanisms that were caused by Si3N4, pull out of Fe3Si particle can also be found in Figure 7b, indicating that Fe3Si also plays a part in toughening the composites. The toughening effect of Fe3Si was further studied by Vickers indentation (Figure 8). Four typical cracks (Figure 8a-d) in sample FSN60 were selected for detailed analysis. The propagation paths of the cracks shown in Figure 8a,b were free of Fe3Si particles, and although crack deflection by Si3N4 crystals can be seen (Figure 8a), the crack lengths was relatively long (24.0 μm and 18.9 um, respectively). In Figure 8c,d, where the cracks interacted with Fe3Si in the form of interface debonding, particle bridging, and crack deflection, the crack length was smaller (15.3 μm and 14.7 um, respectively), which indicates an improved toughness. Since Fe3Si has better plasticity than the brittle Si3N4, when the crack propagated to the vicinity of the Fe3Si particle, the stress concentration In order to research the microstructure-the mechanical properties' relationship of the composites, fracture surface and crack/microstructure interactions were observed. Figure 7 shows the fracture surface of FSN40. In Figure 7a, it can be seen that fracture modes of the composites were in co-action by transgranular and intergranular fracture. Since Si 3 N 4 has its well-known characteristic of self-reinforcement, two kinds of typical toughening mechanisms of crack bridging and crack deflection by Si 3 N 4 can be found in Figure 7b, which resulted from the elongated Si 3 N 4 crystals that were bounded by the weak interface [40]. Apart from the toughening mechanisms that were caused by Si 3 N 4 , pull out of Fe 3 Si particle can also be found in Figure 7b, indicating that Fe 3 Si also plays a part in toughening the composites. with the cost-effective Fe-Si3N4 powder, composites can obtain mechanical properties-especially fracture toughness that is at the same level with dense monolithic Si3N4 ceramics. In order to research the microstructure-the mechanical properties' relationship of the composites, fracture surface and crack/microstructure interactions were observed. Figure 7 shows the fracture surface of FSN40. In Figure 7a, it can be seen that fracture modes of the composites were in co-action by transgranular and intergranular fracture. Since Si3N4 has its well-known characteristic of self-reinforcement, two kinds of typical toughening mechanisms of crack bridging and crack deflection by Si3N4 can be found in Figure 7b, which resulted from the elongated Si3N4 crystals that were bounded by the weak interface [40]. Apart from the toughening mechanisms that were caused by Si3N4, pull out of Fe3Si particle can also be found in Figure 7b, indicating that Fe3Si also plays a part in toughening the composites. The toughening effect of Fe3Si was further studied by Vickers indentation (Figure 8). Four typical cracks (Figure 8a-d) in sample FSN60 were selected for detailed analysis. The propagation paths of the cracks shown in Figure 8a,b were free of Fe3Si particles, and although crack deflection by Si3N4 crystals can be seen (Figure 8a), the crack lengths was relatively long (24.0 μm and 18.9 um, respectively). In Figure 8c,d, where the cracks interacted with Fe3Si in the form of interface debonding, particle bridging, and crack deflection, the crack length was smaller (15.3 μm and 14.7 um, respectively), which indicates an improved toughness. Since Fe3Si has better plasticity than the brittle Si3N4, when the crack propagated to the vicinity of the Fe3Si particle, the stress concentration The toughening effect of Fe 3 Si was further studied by Vickers indentation (Figure 8). Four typical cracks (Figure 8a-d) in sample FSN60 were selected for detailed analysis. The propagation paths of the cracks shown in Figure 8a,b were free of Fe 3 Si particles, and although crack deflection by Si 3 N 4 crystals can be seen (Figure 8a), the crack lengths was relatively long (24.0 µm and 18.9 µm, respectively). In Figure 8c,d, where the cracks interacted with Fe 3 Si in the form of interface debonding, particle bridging, and crack deflection, the crack length was smaller (15.3 µm and 14.7 µm, respectively), which indicates an improved toughness. Since Fe 3 Si has better plasticity than the brittle Si 3 N 4 , when the crack propagated to the vicinity of the Fe 3 Si particle, the stress concentration around the crack tip can be somewhat reduced, thus the tendency to the ripping of the material was inhibited. The CTE of Fe 3 Si (14.4 × 10 −6 K −1 [41]) was much bigger than that of Si 3 N 4 (2.9 × 10 −6 K −1 [39]), so the interface between Fe 3 Si and Si 3 N 4 was under tensile stress at room temperature, resulting in a weak interfacial bonding strength. When under stress, the weak interface between Fe 3 Si and the surrounding phase debonded (Figure 8c,d). This led to a tortuous crack path (Figure 8d) or particle bridging (Figure 8d).  [41]) was much bigger than that of Si3N4 (2.9 × 10 −6 K −1 [39]), so the interface between Fe3Si and Si3N4 was under tensile stress at room temperature, resulting in a weak interfacial bonding strength. When under stress, the weak interface between Fe3Si and the surrounding phase debonded (Figure 8c,d). This led to a tortuous crack path (Figure 8d) or particle bridging (Figure 8d). With the above-revealed toughening mechanism of Fe3Si, the dependence of the flexural strength and fracture toughness of the Si3N4-Fe3Si composites on the content of Fe-Si3N4 in starting powders can be explained. The inherent brittleness of ceramics is determined by its poor plasticity, so that when under stress, little energy can be consumed by the plastic flow. In order to improve the fracture toughness of ceramics, other energy dissipation mechanisms like crack deflection, bridging, or particle pull-out are often utilized in ceramic composites. Compared with the sample of FSN20, FSN40 contains more Fe3Si particles and the particle size of Fe3Si in it remains small ( Figure 5). Large amounts of fine dispersed Fe3Si particles improved the strength and toughness of FSN40. When the content of Fe-Si3N4 in starting powders was more than 60 wt. %, despite the fact that the phase content of Fe3Si increased, the Fe3Si particles suffered severe growth, and large particles that were bigger than 15 μm emerged (Figure 5b). Large Fe3Si particles result in a non-uniform microstructure and may serve as crack origins when the samples are loaded, so the mechanical properties of FSN60, FSN80, and FSN90 are damaged. Narciso [42,43] studied the coefficient of the thermal expansion (CTE) properties of several metal-ceramic composites, and it was found that metals usually have CTE one magnitude higher than that of ceramics. In our study, the interface of Fe3Si and Si3N4 are under residue tensile stress due to the mismatch of CTEs. The residue tensile stress may result in cracks and With the above-revealed toughening mechanism of Fe 3 Si, the dependence of the flexural strength and fracture toughness of the Si 3 N 4 -Fe 3 Si composites on the content of Fe-Si 3 N 4 in starting powders can be explained. The inherent brittleness of ceramics is determined by its poor plasticity, so that when under stress, little energy can be consumed by the plastic flow. In order to improve the fracture toughness of ceramics, other energy dissipation mechanisms like crack deflection, bridging, or particle pull-out are often utilized in ceramic composites. Compared with the sample of FSN20, FSN40 contains more Fe 3 Si particles and the particle size of Fe 3 Si in it remains small ( Figure 5). Large amounts of fine dispersed Fe 3 Si particles improved the strength and toughness of FSN40. When the content of Fe-Si 3 N 4 in starting powders was more than 60 wt. %, despite the fact that the phase content of Fe 3 Si increased, the Fe 3 Si particles suffered severe growth, and large particles that were bigger than 15 µm emerged (Figure 5b). Large Fe 3 Si particles result in a non-uniform microstructure and may serve as crack origins when the samples are loaded, so the mechanical properties of FSN60, FSN80, and FSN90 are damaged. Narciso [42,43] studied the coefficient of the thermal expansion (CTE) properties of several metal-ceramic composites, and it was found that metals usually have CTE one magnitude higher than that of ceramics. In our study, the interface of Fe 3 Si and Si 3 N 4 are under residue tensile stress due to the mismatch of CTEs. The residue tensile stress may result in cracks and voids between Fe 3 Si and Si 3 N 4 , which can act as a crack origin when under stress, and this may be one reason for the low mechanical properties of the composite with a high Fe 3 Si content.

Conclusions
In this study, Si 3 N 4 -Fe 3 Si composites with a different content of the Fe 3 Si phase were fabricated using starting powders of different compositions. The microstructure and mechanical properties of the composites were investigated by various methods. Special attention was placed on the particle size distribution of Fe 3 Si and their effect on the mechanical properties of the composites. The main conclusions can be summarized as follows.
The sintered composites were mainly composed of Si 3 N 4 , Fe 3 Si, and the grain boundary phase. With the increase of Fe-Si 3 N 4 powder from 20 wt. % to 90 wt. % in the starting powder, the Fe 3 Si phase content increased from 0.7 vol. % to 4.1 vol. %, and the particle size increased from 1.2 µm to 2.8 µm. When more than 60 wt. % Fe-Si 3 N 4 was added to the starting powders, the abnormal growth of Fe 3 Si particles occurred and particles bigger than 15 µm were commonly seen, leading to non-uniform microstructures and poor mechanical properties. The dispersive Fe 3 Si particles had a toughening effect on the composites through mechanisms such as particle pull-out, interface debonding, crack deflection, and particle bridging. Owing to the uniform microstructure and Fe 3 Si toughening, FSN40 showed the highest flexural strength and fracture toughness of 354 MPa and 8.4 MPa·m 1/2 , respectively, indicating great potential as thermal structural materials.