Effect of C Addition on the Microstructure and Fracture Properties of In Situ Laminated Nb/Nb5Si3 Composites

//Nbss and α-Nb5Si3 phases were detected. Meanwhile, Nb2C was observed, and the crystal forms of Nb5Si3 changed in the C-doped composites. Furthermore, micron-sized and nano-sized Nb2C particles were found in the Nbss layer. The orientation relationship of Nb2C phase and the surrounding Nbss was [001]Nbss//[010]Nb2C, (200) Nbss//(101) Nb2C. Additionally, with the addition of C, the compressive strength of the composites, at 1400 °C, and the fracture toughness increased from 310 MPa and 11.9 MPa·m1/2 to 330 MPa and 14.2 MPa·m1/2, respectively; the addition of C mainly resulted in solid solution strengthening.


Introduction
The Nb 5 Si 3 intermetallic compound has been considered as a potential material for high-performance structural application due to its high melting point (2520 • C), low density (7.1 g/cm 3 ), and excellent strength retention at elevated temperatures [1][2][3]. However, due to the relatively low fracture toughness of 1-3 MPa·m 1/2 at ambient temperatures [4], a ductile niobium-based solid solution (Nb ss ) was brought into Nb 5 Si 3 to achieve a balance of low fracture toughness and high temperature strength [5][6][7][8][9][10][11][12]. As for the as-cast Nb-Si alloy, with the increase in Si content, the volume fraction of ductile Nb ss phase decreased, thereby significantly lowering the fracture toughness [13,14]. It was reported that the fracture toughness of Nb-10Si and Nb-16Si alloys were 12 MPa·m 1/2 and 4.5 MPa·m 1/2 , respectively [15]. In order to improve the fracture toughness of Nb-Si alloys, a number of reinforcement elements such as Ti, Mo, and B were added as well [16,17]. Wang et al. [18] reported that the fracture toughness of Nb-16Si alloy improved after increasing the Hf content. Furthermore, the addition of B could also enhance the fracture toughness of a Nb-10W-10Si alloy [19].
Preparing unidirectionally solidified alloys and laminated composites has also been shown to be an efficient method of enhancing fracture toughness. Ye et al. [20] found that excellent fracture toughness of 14.5 MPa·m 1/2 and 18.7 MPa·m 1/2 could be exhibited by unidirectionally solidified Nb-Si and Nb-Si-Ti alloys, respectively. A Nb 5 Si 3 /Nb/Nb 5 Si 3 laminate with a relatively high fracture toughness of 7.1-11.5 MPa·m 1/2 was fabricated by hot pressing the Nb 5 Si 3 compacts and Nb foil at 1200 • C for 5 h [21]. However, since the thickness of the Nb 5 Si 3 compacts and Nb foil were 4 mm and 0.25 mm, respectively, the fracture toughness apparently changed into the distance changes of notch from the Nb/Nb 5 Si 3 interface. Thus, the thickness of Nb ss and Nb 5 Si 3 layers should be decreased, Materials 2023, 16 and the in situ laminated Nb/Nb 5 Si 3 composite with micron-sized multi-layer structures was fabricated from the previous work [22]. Furthermore, it was reported that the Nb-16Si-10Mo-15W alloy could be strengthened by addition a solution to B in the Nb 5 Si 3 phases [23][24][25]. Similarly, it was confirmed that a C atom could also dissolve in the Nb 3 Al phase [26][27][28]. Due to fact that the C atom has a smaller radius than the B atom, the strengthening effect of adding a solid solution to C in the Nb 5 Si 3 would be probably better than that of B. Reports regarding the addition of C to a Nb/Nb 5 Si 3 alloy are scarce in the open literature. It is unclear whether and to what extent the addition of C can improve the fracture toughness of a Nb/Nb 5 Si 3 alloy. Therefore, the aim of this work was to prepare in situ laminated Nb/Nb 5 Si 3 composites supplemented with C via spark plasma sintering, evaluate the effect of C on the microstructure and mechanical properties of the composites, and identify the strengthening and toughening mechanisms.

Experimental Procedure
Nb foils (99.99%, 25 µm in thickness), Nb powders (99.99%, 1-3 µm), Si powders (99.99%, 1-3 µm), and C powders (99.99%, 1-3 µm) were the raw materials used. For the experimental work, a three-step procedure was adopted. Firstly, mixtures of molar ratios of Nb-50Si and Nb-40Si-10C were selected to prepare the Nb/Si/(C) slurry via vacuum ball milling for 24 h using ethanol as a milling medium. For convenience of expression, the corresponding prepared materials are called Nb-50Si and Nb-40Si-10C, respectively. Secondly, the Nb foils were covered with mixed Nb/Si/(C) slurry via the dip-coating method and then stacked together and dried at 120 • C for 24 h in a vacuum. Thirdly, the stacked Nb foils were put in a graphite die and sintered at 1750 • C under a pressure of 30 MPa for 30 min in vacuum using a heating rate of 100 • C/min. Finally, the sintered Nb/Nb 5 Si 3 composites were cooled at about a rate of 100 • C/min above 500 • C and then furnace-cooled down to room temperature. A diagram of the preparation process for the niobium-based composites is shown in Figure 1. The sintered material had a porosity of 0.3381% and a density of 8.1621 g/cm 3 .
The phase and crystallinity were analyzed via X-ray diffraction (XRD) using CuK α radiation at 40 kV and 250 mA. The lattice parameters were calculated by using JADE 5 software. The microstructures of the samples were characterized using scanning electron microscopy (SEM), wavelength-dispersive spectroscopy (WDS), and transmission electron microscopy (TEM). The SEM samples were cut via electrical discharge machining (EDM) and polished to a surface finish using 1 µm diamond paste, and the TEM foils were prepared via ion milling. The volume fractions of phase in the composite were calculated via quantitative image analysis using EPMA micrographs; five EPMA images were used for each composite. Fracture toughness was determined via three-point bending (TPB) tests at room temperature. In the TPB tests, a specimen with a dimension of 2.5 mm × 5 mm × 20 mm and a notch introduced perpendicular to the layer direction was cut via electro-discharge machining (EDM), and the cross-head speed was set at a rate of 0.1 mm/min. Compression tests were conducted at 1400 • C at a strain rate of 10 −3 s −1 in a vacuum. The dimension of the compression test specimen was ϕ4 mm × 6 mm, and the loading direction was parallel to the layer direction. Five specimens were tested for each condition, and the average values were recorded. A computer microhardness tester (200HBVS-30) was used to measure the Vickers hardness of the alloy under a load of 15 N, and the test time was 15 s. The same alloy sample was tested 5 times at random locations; a group of 3 samples were tested, and the average value of each measurement was taken as the Vickers hardness of the alloy.  Figure 2 shows XRD patterns of the Nb-50Si and Nb-40Si-10C composites. It found that the obtained Nb-50Si composite exhibited an XRD pattern typical of Nbss α-Nb5Si3. However, following the addition of C, Nb2C and γ-Nb5Si3 were present in Nb-40Si-10C composite. This indicated that the addition of C promoted the formatio the metastable γ-Nb5Si3 phase and high temperature β-Nb5Si3 phase. Additionally should be pointed out that SiC was not identified in the patterns. The PDF card numb of the phases involved in the Figure 1 are shown in Table 1.   Figure 2 shows XRD patterns of the Nb-50Si and Nb-40Si-10C composites. It was found that the obtained Nb-50Si composite exhibited an XRD pattern typical of Nb ss and α-Nb 5 Si 3 . However, following the addition of C, Nb 2 C and γ-Nb 5 Si 3 were present in the Nb-40Si-10C composite. This indicated that the addition of C promoted the formation of the metastable γ-Nb 5 Si 3 phase and high temperature β-Nb 5 Si 3 phase. Additionally, it should be pointed out that SiC was not identified in the patterns. The PDF card numbers of the phases involved in the Figure 1 are shown in Table 1.  It was assumed that little SiC remained in the material. The following induced:

Results and Discussion
The thermodynamic results of reaction (1) are listed in the following r tion (2)) according to the thermodynamic data shown in Table 2.
The molar heat capacities of the various substances in reaction (1) at 140 the standard Gibbs free energy at different temperatures is shown in Figur in Figure 3, with increasing temperature, the standard Gibbs free energy o (1) decreased, and all the values were negative. Therefore, according to the reaction (1) occurred during the sintering process. Therefore, all of the XRD the absence of SiC in the Nb/Nb5Si3 composites. It was assumed that little SiC remained in the material. The following reaction was induced: 11Nb + 3SiC = Nb 5 Si 3 + 3Nb 2 C (1) The thermodynamic results of reaction (1) are listed in the following reaction (reaction (2)) according to the thermodynamic data shown in Table 2.  The standard Gibbs free energy of reaction (1) can be expressed as (3), according to the calculation of the second approximation equation of thermodynamics: The molar heat capacities of the various substances in reaction (1) at 1400-1800 K and the standard Gibbs free energy at different temperatures is shown in Figure 3. As shown in Figure 3, with increasing temperature, the standard Gibbs free energy of the reaction (1) decreased, and all the values were negative. Therefore, according to the above results, reaction (1) occurred during the sintering process. Therefore, all of the XRD patterns show the absence of SiC in the Nb/Nb 5 Si 3 composites.    Table 3) and niobium compound layers can be observed. In the Nb-50Si composite, the average thicknesses of the Nbss layers were decreased from 25 µm to 12.7 µm with increasing sintering time, confirm Si element diffusion from the Nb5Si3 layers (point 2 in Table 3) to the Nbss layers during sintering. Interestingly, with the addition of C, we observed that the microstructures of the composites significantly changed. According to the WDS results shown in Table 3, the Nb2C particle (3-5 µm) was present in the Nbss layers and exhibited a morphology different from the carbide in as-cast Nb-20Ti-12.5C-Mo-Hf alloys [29]. In addition, a lot of fine carbide (nano-sized) was also observed in the Nbss layer. The formation mechanism of the carbide will be discussed later. Secondly, both C-rich (point 4) and C-poor (point 3) Nb5Si3 were observed in the compound layers, which can be attributed to the different diffusion rates of the Si and C atoms in Nb. Due to the lighter atomic mass and smaller atomic radius of C compared to Si, the diffusion rate of C should be higher than Si, leading to the longer diffusion distance of the C element. As a result, the C-rich Nb5Si3 is closer to the Nbss layers. A relatively high oxygen content of 1.3 wt% was detected in the Nbss layer of the Nb-50Si composite. This was due to the fact that the dipping and stacking of the Nb foils was preformed in the air. Due to the fine Nb and Si powders, it is very difficult to avoid the physisorption of oxygen during the synthesis of the composites. However, interestingly, we observed that the O content (point 5) in the Nbss layer decreased to 0.3 wt.% in the Nb-40Si-10C composite.  Table 4.   Table 3) and niobium compound layers can be observed. In the Nb-50Si composite, the average thicknesses of the Nb ss layers were decreased from 25 µm to 12.7 µm with increasing sintering time, confirm Si element diffusion from the Nb 5 Si 3 layers (point 2 in Table 3) to the Nb ss layers during sintering. Interestingly, with the addition of C, we observed that the microstructures of the composites significantly changed. According to the WDS results shown in Table 3, the Nb 2 C particle (3-5 µm) was present in the Nb ss layers and exhibited a morphology different from the carbide in as-cast Nb-20Ti-12.5C-Mo-Hf alloys [29]. In addition, a lot of fine carbide (nano-sized) was also observed in the Nb ss layer. The formation mechanism of the carbide will be discussed later. Secondly, both C-rich (point 4) and C-poor (point 3) Nb 5 Si 3 were observed in the compound layers, which can be attributed to the different diffusion rates of the Si and C atoms in Nb. Due to the lighter atomic mass and smaller atomic radius of C compared to Si, the diffusion rate of C should be higher than Si, leading to the longer diffusion distance of the C element. As a result, the C-rich Nb 5 Si 3 is closer to the Nb ss layers. A relatively high oxygen content of 1.3 wt% was detected in the Nb ss layer of the Nb-50Si composite. This was due to the fact that the dipping and stacking of the Nb foils was preformed in the air. Due to the fine Nb and Si powders, it is very difficult to avoid the physisorption of oxygen during the synthesis of the composites. However, interestingly, we observed that the O content (point 5) in the Nb ss layer decreased to 0.3 wt.% in the Nb-40Si-10C composite.   Table 3) and niobium compound layers can be o In the Nb-50Si composite, the average thicknesses of the Nbss layers were decrea 25 µm to 12.7 µm with increasing sintering time, confirm Si element diffusion Nb5Si3 layers (point 2 in Table 3) to the Nbss layers during sintering. Interestingly, addition of C, we observed that the microstructures of the composites sign changed. According to the WDS results shown in Table 3, the Nb2C particle (3-5 present in the Nbss layers and exhibited a morphology different from the carbide i Nb-20Ti-12.5C-Mo-Hf alloys [29]. In addition, a lot of fine carbide (nano-sized) observed in the Nbss layer. The formation mechanism of the carbide will be discus Secondly, both C-rich (point 4) and C-poor (point 3) Nb5Si3 were observed in pound layers, which can be attributed to the different diffusion rates of the Si and in Nb. Due to the lighter atomic mass and smaller atomic radius of C compared diffusion rate of C should be higher than Si, leading to the longer diffusion distan C element. As a result, the C-rich Nb5Si3 is closer to the Nbss layers. A relatively h gen content of 1.3 wt% was detected in the Nbss layer of the Nb-50Si composite. due to the fact that the dipping and stacking of the Nb foils was preformed in the to the fine Nb and Si powders, it is very difficult to avoid the physisorption o during the synthesis of the composites. However, interestingly, we observed th content (point 5) in the Nbss layer decreased to 0.3 wt.% in the Nb-40Si-10C comp   Table 4. Table 3. WDS composition analysis of the micro-areas in Figure 3.

Position Composition (at%) Possible Phase
Nb  The decrease in oxygen content can be explained via thermodynamic calculation. It is assumed that NbO exists in the material, and the following reactions can be assumed: According to the relevant thermodynamic constants in Table 4, the thermodynamic calculation of reaction (4) is as follows:  The decrease in oxygen content can be explained via thermodynamic calculation is assumed that NbO exists in the material, and the following reactions can be assumed NbO + C = Nb + CO According to the relevant thermodynamic constants in Table 4, the thermodynam calculation of reaction (4) is as follows:   As can be seen from Figure 5b, with increasing temperature, the standard Gibbs fr energy of the reaction decreases. When the temperature reaches 1900 K (~1627 °C), t standard Gibbs free energy is negative. As the sintering temperature of the alloy is high As can be seen from Figure 5b, with increasing temperature, the standard Gibbs free energy of the reaction decreases. When the temperature reaches 1900 K (~1627 • C), the standard Gibbs free energy is negative. As the sintering temperature of the alloy is higher than 1750 • C, reaction (4) can proceed smoothly with the sintering process, according to the above thermodynamic calculation results. Figure 6 shows typical SEM micrographs from the fracture surfaces of sintered composites after the TPB tests. It can be seen from Figure 6a that the fracture surface of the Nb-50Si composite basically exhibited an intergranular fracture mode in conjunction with some cleavage fracture features in the Nb 5 Si 3 layer. Meanwhile, transgranular cracking and some ridge-like features were observed in the Nb ss layer. However, more transgranular cracking and some ridge-like features were observed due to the presence of the brittle Nb 2 C phase. This proved that the addition of C changed the fracture mechanisms of the composites. Additionally, regarding the Nb-40Si-10C composite (Figure 6b), partial dimples were observed in the Nb ss layer.
Materials 2023, 16, x FOR PEER REVIEW 7 of 1 than 1750 °C, reaction (4) can proceed smoothly with the sintering process, according t the above thermodynamic calculation results. Figure 6 shows typical SEM micrographs from the fracture surfaces of sintered com posites after the TPB tests. It can be seen from Figure 6a that the fracture surface of th Nb-50Si composite basically exhibited an intergranular fracture mode in conjunction wit some cleavage fracture features in the Nb5Si3 layer. Meanwhile, transgranular crackin and some ridge-like features were observed in the Nbss layer. However, more transgran ular cracking and some ridge-like features were observed due to the presence of the brittl Nb2C phase. This proved that the addition of C changed the fracture mechanisms of th composites. Additionally, regarding the Nb-40Si-10C composite (Figure 6b), partial dim ples were observed in the Nbss layer. In order to investigate the formation mechanisms of the Nb2C phase in the Nbss laye the order of the reactions in the sintering layer was evaluated via thermodynamic calcu lation. As can be seen from Figure 4, three elements, namely Nb, Si, and C, were observe in the sintering layers. Furthermore, the C atoms tended to spread throughout the Nb layers. Hence, what needs to be confirmed is whether the following two reactions oc curred during the sintering process: 15Nb2C + 18Si = 6Nb5Si3 + 15C (6 Nb2C + 4Si = 2NbSi2 + C (7 According to the thermodynamic constants shown in Table 5, thermodynamic reac tions (6) and (7) can be calculated as follows: In order to investigate the formation mechanisms of the Nb 2 C phase in the Nb ss layer, the order of the reactions in the sintering layer was evaluated via thermodynamic calculation. As can be seen from Figure 4, three elements, namely Nb, Si, and C, were observed in the sintering layers. Furthermore, the C atoms tended to spread throughout the Nb ss layers. Hence, what needs to be confirmed is whether the following two reactions occurred during the sintering process: 15Nb 2 C + 18Si = 6Nb 5 Si 3 + 15C Nb 2 C + 4Si = 2NbSi 2 + C According to the thermodynamic constants shown in Table 5, thermodynamic reactions (6) and (7) ∆H Θ 298(7) = −81, 170 J ∆S Θ 298(7) = 6.151 J·K −1 (9) Table 5. Thermodynamic data of the materials in reactions (6) and (7). The molar heat capacity C P of the various substances at 1000-1600 K are shown in Figure 7. Therefore, the standard Gibbs free energy ∆G Θ T in reactions (6) and (7) at different temperatures could be obtained by using the second approximation equation of thermodynamics. It can be seen from Table 5 that, when the temperature was above 1000 K, the standard Gibbs free energy in reactions (6) and (7) decreased with increasing temperature. Furthermore, all of the values of the standard Gibbs free energy shown in Figure 7 are negative. This proves that reactions (6) and (7) could occur during the sintering process. In other words, even though Nb 2 C remained in the compound layers, it can react with Si and enter into C. It was indicated that a part of the carbon in the solid solution of Nb 5 Si 3 and the others diffuses into the Nb ss layers and then reacts with Nb, forming Nb 2 C.  The molar heat capacity of the various substances at 1000-1600 K are shown Figure 7. Therefore, the standard Gibbs free energy ∆ in reactions (6) and (7) at d ferent temperatures could be obtained by using the second approximation equation thermodynamics. It can be seen from Table 5 that, when the temperature was above 10 K, the standard Gibbs free energy in reactions (6) and (7) decreased with increasing te perature. Furthermore, all of the values of the standard Gibbs free energy shown in Figu 7 are negative. This proves that reactions (6) and (7) could occur during the sintering pr cess. In other words, even though Nb2C remained in the compound layers, it can rea with Si and enter into C. It was indicated that a part of the carbon in the solid solution Nb5Si3 and the others diffuses into the Nbss layers and then reacts with Nb, forming Nb2   Figure 8a. Figure 8b presen the region's selection in Figure 8a, which is about 500 nm in diameter. The polycrystalli rings can be observed in Figure 8b, which indicates that there are multiple grains in th region. Therefore, it can be suggested that the grain size of raw Nb foil is below 500 n and that there are many grain boundaries in the raw Nb foil, providing a channel for t diffusion of C atoms. During the sintering process, the C atoms rapidly spread into t Nbss and react with Nb in situ to form micron Nb2C particles in the Nbss layer.   Figure 8a. Figure 8b presents the region's selection in Figure 8a, which is about 500 nm in diameter. The polycrystalline rings can be observed in Figure 8b, which indicates that there are multiple grains in this region. Therefore, it can be suggested that the grain size of raw Nb foil is below 500 nm and that there are many grain boundaries in the raw Nb foil, providing a channel for the diffusion of C atoms. During the sintering process, the C atoms rapidly spread into the Nb ss and react with Nb in situ to form micron Nb 2 C particles in the Nb ss layer.  The appearance of nanosized Nb2C is mainly due to a change in the solid solubility of C in Nbss. According to the Nb-C binary phase diagram [30], it is clear that when the temperature is above 1500 °C the solid solubility of C in Nbss decreases considerably with decreasing temperature Therefore, when the prepared Nb/Nb5Si3 composites are subjected to cooling at a sintering temperature of 1750 °C, a lot of nano-sized Nb2C can be precipitated from Nbss.  Figure 10 shows TEM images and diffraction patterns typical of nanometer Nb2C in the Nb-40Si-10C alloy. Club-shaped nano-sized Nb2C particles can be observed in Figure  10a, the length and width values of which are 100-300 nm and 70-130 nm, respectively Nb2C was also observed in the Nb-40Si-10C alloy, and the zone axes along [21 5 ] are presented in Figure 10b.   Nb2C . The appearance of nano-sized Nb 2 C is mainly due to a change in the solid solubility of C in Nb ss . According to the Nb-C binary phase diagram [30], it is clear that when the temperature is above 1500 • C, the solid solubility of C in Nb ss decreases considerably with decreasing temperature. Therefore, when the prepared Nb/Nb 5 Si 3 composites are subjected to cooling at a sintering temperature of 1750 • C, a lot of nano-sized Nb 2 C can be precipitated from Nb ss .  The appearance of nanosized Nb2C is mainly due to a change in the solid solubility of C in Nbss. According to the Nb-C binary phase diagram [30], it is clear that when the temperature is above 1500 °C the solid solubility of C in Nbss decreases considerably with decreasing temperature Therefore, when the prepared Nb/Nb5Si3 composites are subjected to cooling at a sintering temperature of 1750 °C, a lot of nano-sized Nb2C can be precipitated from Nbss.    Figure 10 shows TEM images and diffraction patterns typical of nanometer Nb 2 C in the Nb-40Si-10C alloy. Club-shaped nano-sized Nb 2 C particles can be observed in Figure 10a, the length and width values of which are 100-300 nm and 70-130 nm, respectively. Nb 2 C was also observed in the Nb-40Si-10C alloy, and the zone axes along [215] are presented in Figure 10b. The lattice constants of each phase can be obtained by analyzing and calculating XRD patterns of the different components of the alloy using the Jade software (MDI Ja 6.0). Table 6 shows the lattice parameters of Nbss and α-Nb5Si3 in the composites. Rega ing the Nb-50Si and Nb-40Si-10C composites, the lattice parameters of Nbss and α-Nb decreased with the addition of C. The assumption that C atoms mainly occupy the sub tutional sites in Nbss and α-Nb5Si3 can be confirmed by the fact that the atomic radius o is smaller than that of Nb and/or Si, thereby forming a replacement solid solution. The average compressive 0.2% flow stress at 1400 °C and fracture toughness at a bient temperature are shown in Table 7. It can be seen that the mechanical properties the Nb/Nb5Si3 composites were significantly enhanced following the addition of C. T could be attributed to the following three reasons: Firstly, the C in Nbss and Nb5Si3 play a key role in solution strengthening and improving high-temperature strength. The d solution of carbon atoms in both the Nbss and Nb5Si3 lattices was predominantly loca at substitutional sites and decreased the lattice parameters, increasing the deformat resistance. As a result, the compressive strength is influenced by the content of strengthened phase, i.e., Nb2C. As mentioned in Table 7, with the addition of C, the v ume fraction of the plastic phase decreased, while that of the strengthened phase creased. Lastly, the precipitated fine carbide played a role in enhancing the compress strength. Allameh et al. [31] reported that, with the addition of TiC particles, some dis cations in the TiC particles were observed, and it was also reported that their interactio played a significant role in strengthening the 44Nb-35Ti-6Al-5Cr-8V-1W-0.5Mo-0.5 (at.%) alloy. Therefore, it can be inferred that nanoscale Nb2C in the Nb/Nb5Si3 composi will produce similar strengthening effects. Table 7. Volume fractions of each phase and mechanical properties of alloys. The lattice constants of each phase can be obtained by analyzing and calculating the XRD patterns of the different components of the alloy using the Jade software (MDI Jade 6.0). Table 6 shows the lattice parameters of Nb ss and α-Nb 5 Si 3 in the composites. Regarding the Nb-50Si and Nb-40Si-10C composites, the lattice parameters of Nb ss and α-Nb 5 Si 3 decreased with the addition of C. The assumption that C atoms mainly occupy the substitutional sites in Nb ss and α-Nb 5 Si 3 can be confirmed by the fact that the atomic radius of C is smaller than that of Nb and/or Si, thereby forming a replacement solid solution. Table 6. Lattice parameters of Nb ss and α-Nb 5 Si 3 in the composites.

Composites a (Nb ss ) a (α-Nb 5 Si 3 ) c (α-Nb 5 Si 3 )
Nb-50Si 3.31519 6.56971 11.89522 Nb-40Si-10C 3.31026 6.53410 11.86043 The average compressive 0.2% flow stress at 1400 • C and fracture toughness at ambient temperature are shown in Table 7. It can be seen that the mechanical properties of the Nb/Nb 5 Si 3 composites were significantly enhanced following the addition of C. This could be attributed to the following three reasons: Firstly, the C in Nb ss and Nb 5 Si 3 played a key role in solution strengthening and improving high-temperature strength. The dissolution of carbon atoms in both the Nb ss and Nb 5 Si 3 lattices was predominantly located at substitutional sites and decreased the lattice parameters, increasing the deformation resistance. As a result, the compressive strength is influenced by the content of the strengthened phase, i.e., Nb 2 C. As mentioned in Table 7, with the addition of C, the volume fraction of the plastic phase decreased, while that of the strengthened phase increased. Lastly, the precipitated fine carbide played a role in enhancing the compressive strength. Allameh et al. [31] reported that, with the addition of TiC particles, some dislocations in the TiC particles were observed, and it was also reported that their interactions played a significant role in strengthening the 44Nb-35Ti-6Al-5Cr-8V-1W-0.5Mo-0.5Hf (at.%) alloy. Therefore, it can be inferred that nanoscale Nb 2 C in the Nb/Nb 5 Si 3 composites will produce similar strengthening effects.
It also can be seen from Table 7 that the fracture toughness of the composites improved with the addition of C. This can be attributed to the following reasons. First, as mentioned in Figure 4, the O content in the Nb ss layer can be reduced or eliminated with the addition of C. This observation also corresponded well to the fracture morphology results shown in Figure 6. It is known that a large amount of energy could be absorbed from the plastic deformation of the Nb ss . When the plasticity of Nb ss increased, more energy could be consumed, resulting in an increase in the fracture toughness of the composites. The ductility of Nb ss can exhibit a strong resistance to crack initiation during the plastic deformation of 44Nb-35Ti-6Al-5Cr-8V-1W-0.5Mo-0.3Hf (at.%), as reported by Sikka and Loria [32]. Second, the fracture toughness can be affected by some physical properties. According to the Ashby model [33], the toughness increment ∆K C can be expressed as Equation (10): where E, V f , σ 0 , and a 0 are the Young's modulus (GPa), volume fraction, yield strength at ambient temperature (MPa), and radius of the Nb ss phase (m), respectively, and C is the material constant representing the degree of constraint imposed upon a ductile particle from the brittle matrix. In the current work, since the Nb ss phase became deformed without interface decohesion (Figure 3), the parameter C is taken to be 1.6 [34]. The volume fraction and average radius of the Nb ss can be obtained from Figure 3. The Young's modulus and Vickers hardness were measured, and the yield strength σ 0 (MPa) of Nb ss phase can be estimated from the Vickers hardness of the Nb ss phase using the following equation [35]: The mechanical and physical properties of the composites are presented in Table 8. Clearly, due to the existence of nano-sized carbide, all of the Young's modulus, Vickers hardness, and yield strength values were increased in the C-doped composites. It has been reported that, the hardness and Young's modulus of Nb 2 C is higher than that of Nb ss [36]. With the addition of C, Nb ss was transformed to Nb 2 C. Based on the rule of mixtures [32], the hardness and Young's modulus of the Nb ss layer would increase, leading to an increase in yield strength, according to Equation (11).

Conclusions
Laminated Nb/Nb 5 Si 3 composites supplemented with C were prepared via spark plasma sintering. Nb ss and γ-Nb 5 Si 3 were found In the Nb-Si-C composites, and micronsized Nb 2 C particles and nano-sized Nb 2 C were observed in the Nb ss layer. The formation of Nb 2 C particles might be attributable to the rapid diffusion of C into the Nb foil during sintering, and the formation of nano-sized Nb 2 C could be attributable to C's solid solubility change in Nb ss . Additionally, with the addition of C, the compressive strength of composites at 1400 • C and the fracture toughness increased from 310 MPa and 11.9 MPa·m 1/2 to 330 MPa and 14.2 MPa·m 1/2 , respectively; the addition of C mainly resulted in solid solution strengthening.