Effect of Additive Ti₃SiC₂ Content on the Mechanical Properties of B₄C–TiB₂ Composites Ceramics Sintered by Spark Plasma Sintering

Abstract

B₄C–TiB₂ composite ceramics with ultra-high fracture toughness were successfully prepared via spark plasma sintering (SPS) at 1900 °C using B₄C and Ti₃SiC₂ as raw materials. The results showed that compared with pure B₄C ceramics sintered by SPS, the hardness of B₄C–TiB₂ composite ceramics was decreased, but the flexural strength and fracture toughness were significantly improved; the fracture toughness especially was greatly improved. When the content of Ti₃SiC₂ was 30 vol.%, the B₄C–TiB₂ composite ceramic had the best comprehensive mechanical properties: hardness, bending strength and fracture toughness were 27.28 GPa, 405.11 MPa and 18.94 MPa·m^1/2, respectively. The fracture mode of the B₄C–TiB₂ composite ceramics was a mixture of transgranular fracture and intergranular fracture. Two main reasons for the ultra-high fracture toughness were the existence of lamellar graphite at the grain boundary, and the formation of a three-dimensional interpenetrating network covering the whole composite.

1. Introduction

Boron carbide is an attractive engineering material with a high melting point, low density, high hardness, high thermal conductivity and a large neutron absorption surface, which makes it a candidate material for wear-resistant parts, cutting tools, light armor products and neutron radiation shielding [1,2]. However, its low sintering property (due to the strong B–C covalent bond and B₂O₃ oxide layer) and poor fracture toughness limits its excellent performance. Spark plasma sintering (SPS) is a kind of electric current-assisted sintering technology, which can enhance the bonding and densification of particles through the combination of mechanical pressure, an electric field and a thermal field [3,4]. SPS adopts the same stamping/die system concept as hot pressing, although the heating methods are fundamentally different. Hot pressing sintering is heated by heater radiation, while the SPS heat source is Joule heat generated by the current of the mold or sample [5,6]. A heating rate of up to 1000 °C/min can be obtained by SPS, and the heating-up time is greatly shortened, which is beneficial to limit the grain growth [7]. In addition, currents can also enhance powder sintering by activating one or more parallel mechanisms, such as surface oxide removal, electromigration and electroplasticity [8].

Reasonable use of additives can stimulate boron carbide densification without any deterioration of mechanical properties. Some additives can react with boron carbide in situ to form nonvolatile second phases, which is helpful for densification and can enhance properties. Ti₃SiC₂ can react with B₄C to form TiB₂ with high hardness and a high melting point, which can be used as an ideal toughening phase for B₄C ceramics [9]. In this study, B₄C–TiB₂ composite ceramics with ultra-high toughness were prepared by the SPS process with different contents of Ti₃SiC₂ as additives, and the influence mechanism of Ti₃SiC₂ content on the microstructure and properties of B₄C–TiB₂ composite ceramics was studied.

2. Experimental Procedure

Commercially available B₄C powders (purity 99.9%, 1 μm, 4.53 g/cm³, Nangong Naiyate Alloy Welding Material Co., Ltd., Nangong, China) and Ti₃SiC₂ powders (purity 99.9%, <74 μm, 4.53 g/cm³, Nanjing Mingchang New Material Co., Ltd., Nanjing, China) were used as raw materials. Ti₃SiC₂–B₄C powders containing 20 vol.%, 25 vol.%, 30 vol.% and 35 vol.% Ti₃SiC₂, respectively, were mixed for 24 h through a small vertical mixer at 80 r/min without adding solvent. Samples were prepared by SPS equipment (HP D 25/4-SD, FCT Systeme GmbH, Frankenblick, Germany) in a vacuum with 35 MPa mechanical pressure at 1900 °C for 5 min. The heating rate was 100 °C/min and the cooling rate was 50 °C/min.

The absolute density of B₄C–TiB₂ composite ceramics was determined using the Archimedes method. Hardness was measured by a Vickers indentation tester (HMV-2TADW E, Shimadzu, Kyoto, Japan) at a 9.81 N load with a holding time of 15 s on the polished surface. Flexural strength was determined by a three-point bending test with a span of 30 mm and a loading speed of 0.5 mm/min, and the specimens used in the test were 3 mm × 4 mm × 35 mm bars. The SENB (Single-Edge Notched Beam) method was used to determine the fracture toughness of the specimens, with dimensions of 2 mm × 4 mm × 20 mm (with 2 mm high notch). The microstructures of the composite ceramics were characterized by X-Ray powder diffraction (XRD, X′ Pert PRO-MPD, Holland Panalytical, Almelo, Netherlands), scanning electron microscope (SEM, S-4800N, Hitachi, Tokyo, Japan), transmission electron microscope (TEM, JEM-2100, JEOL, Tokyo, Japan) and energy dispersive spectrometer (EDS, INCA, OXFORD INSTRUMENTS, Oxford, UK).

3. Results and Discussion

The scanning electron microscope (SEM) images and X-ray diffraction (XRD) patterns of the as-received powders of B₄C and Ti₃SiC₂ are shown in Figure 1. The SEM images show that B₄C particles have ladder-like surface undulation, a typical transgranular fracture which appears during the particle crushing process; Ti₃SiC₂ particles have an obvious lamellar structure. It can be seen from the XRD images that the two kinds of powders are relatively pure and almost no oxide exists (the content of oxide is too small to be detected in XRD). Figure 2 shows the phase composition of B₄C–TiB₂ ceramic composites prepared at 1900 °C with a different content of additive Ti₃SiC₂. There is no diffraction peak of Ti₃SiC₂ in any of the XRD images, which indicates that Ti₃SiC₂ had completely reacted with B₄C. When the temperature is above 1200 °C, the following reactions occur [10]:

$${\mathrm{B}}_4\mathrm{C}+{\mathrm{Ti}}_3{\mathrm{SiC}}_2\to 2{\mathrm{TiB}}_2+\mathrm{TiC}+\mathrm{SiC}+\mathrm{C}$$

(1)

$${\mathrm{B}}_4\mathrm{C}+2\mathrm{TiC}\to 2{\mathrm{TiB}}_2+3\mathrm{C}$$

(2)

The reactions (1) and (2) ended at 1600 °C, and based on the above results, the overall reaction in the system can be described as the following reaction [10]:

$$3{\mathrm{B}}_4\mathrm{C}+2{\mathrm{Ti}}_3{\mathrm{SiC}}_2\to 6{\mathrm{TiB}}_2+2\mathrm{SiC}+5\mathrm{C}$$

(3)

TiC appears as an intermediate product in the whole reaction process but does not exist in the final product. According to the XRD test results, the content of each phase is shown in

Table 1. Contents of different phases in B₄C–TiB₂ composite ceramics sintered at different temperatures.

Sample Name	Content of Ti3SiC2 (vol.%)	TiB2 (wt.%)	B4C (wt.%)	SiC (wt.%)	C (Graphite) (wt.%)
BT20	20	9.4	87.4	1.6	1.7
BT25	25	19.7	73.2	3.3	3.8
BT30	30	29.6	61.9	4.2	4.4
BT35	35	40.6	47.7	6.1	5.6

. TiB₂ and B₄C are the main phase composition of the composites, and a small amount of SiC and C exist. With the increase in Ti₃SiC₂ content, the proportion of TiB₂, B₄C and C (graphite) in the composite increases, while the content of B₄C decreases.

Table 2. Properties of ceramics prepared with different contents of additive Ti₃SiC₂.

Simple Name	Content of Ti3SiC2 (vol.%)	Density (g/cm3)	Relative Density (%)	Hardness (GPa)	Flexural Strengh (MPa)	Fracture Toughness (MPa·m1/2)
BT0	0	2.50	99.20	33.50	224.43	5.96
BT20	20	3.12	101.16	31.14	317.55	17.68
BT25	25	3.13	101.51	28.70	383.25	18.37
BT30	30	3.17	101.54	27.28	405.11	18.94
BT35	35	3.17	102.28	26.71	343.95	19.00

shows the properties of the samples prepared with different contents of Ti₃SiC₂. In our experiment, the pure B₄C samples sintered by SPS were broken when they were taken out of the mold, so the pure B₄C sample BT0 prepared by Direct Current Sintering at 1800 °C for 10 min was used as the control group. BT0 was the reference sample without Ti₃SiC₂, and its hardness, bending strength and fracture toughness were 33.5 GPa, 224.43 MPa and 5.96 MPa·m^1/2, respectively. Compared with BT0, all of the samples containing Ti₃SiC₂ had a higher relative density. Figure 3 shows the BSE (Backscattered Electron) images of BT0 and BT30 after polishing. There were some closed pores in BT0, but not in BT30. This is because the B₄C particles have sharp edges and corners, and its hardness (55 GPa) is very high. Thus, they cannot be extruded and deformed under pressure, leaving a non-contact space inside, and form pores. The hardness of Ti₃SiC₂ (4 GPa) was much smaller than that of B₄C; Ti₃SiC₂ can be extruded and deformed without leaving voids between particles under an external load. After reaction sintering, TiB₂, B₄C, SiC and C (existing in the form of graphite) in the composites have different thermal expansion coefficients, which makes the ceramics more compact after cooling.

The hardness of the second phase particles produced by the reaction is lower than that of B₄C, especially the graphite phase, therefore it is inevitable that the hardness of the composite ceramics is lower than that of the pure B₄C ceramics. Compared with BT0, the flexural strength and fracture toughness of BT20–BT35 are improved, and the hardness decreases. Figure 4a,b show the fracture morphology of BT0 and BT30, respectively. It can be seen that the fracture surface of BT0 is flat, which is a typical transgranular fracture morphology. Comparatively, the fracture surface of sample BT30 is rough, which is a typically mixed fracture morphology of transgranular fracture and intergranular fracture. In Figure 4b, the dark gray flat area is the B₄C matrix, and the light gray rough area is TiB₂ particles, which indicates that the fracture modes of the B₄C phase and the TiB₂ phase are transgranular fracture and intergranular fracture, respectively. In addition, there is a dark gray lamellar phase around the TiB₂ grain, which can be inferred ad carbon phase by energy spectrum analysis. Figure 4c–e show the energy spectrum of B₄C, TiB₂ and C, respectively. Due to the mismatch of thermal expansion coefficients between B₄C, TiB₂ and graphite (B₄C: 4.5 × 10^–6 k^–1; TiB₂: 8.1 × 10^–6 k^–1; Graphite: 1 × 10^–6 k^–1 in the parallel direction, 29 × 10^–6 k^–1 in c direction) [11], there will be large residual stress at the interface of the phases, which will induce crack deflection along the grain boundary and extend the crack propagation path to improve the strength and toughness of the material. The nano TiB₂ particles embedded in the B₄C matrix will introduce internal stress, which will strengthen the B₄C matrix by a lattice distortion effect, and can also nail the dislocations and hinder their movement, so as to enhance the strength of the material.

It should be noted that the fracture toughness of the composites was greatly improved by adding more than 20 vol.% of Ti₃SiC₂. The fracture toughness of BT30 was 18.94 MPa·m^1/2, which was more than three times of that of BT0 (5.80 MPa·m^1/2), and more than twice as high as the highest fracture toughness cited in

Table 3. Comparison of the properties of the B₄C–TiB₂ composite ceramics reported in recent years.

Serial No.	Starting Powder	Relative Density (%)	KIC, (MPa·m1/2)	Flexural Strength (MPa)	Ref. (Year)
1	B4C + 5 wt% (Ti3SiC2 + Si)	-	5.61	457.6	[14] (2019)
2	B4C + 30 wt% (TiB2 + Si)	99.6	5.77	531.2	[15] (2018)
3	B4C + 20 mol%TiB2	97.9	3.7	-	[16] (2020)
4	B4C + 15 wt%SiC + 20 mol%TiB2	98.6	4.2	343.8	[17] (2020)
5	B4C + 6.45 vol.%SiC + 7.78 vol.%TiB2	99.62	6.38	632	[12] (2019)
6	B4C + 30 vol.% Ti3SiC2	98.72	8.0	492.3	[9] (2017)
BT30	B4C + 30 vol.% Ti3SiC2	101.54	18.94	405.11	This work

. Wen Q et al. [9] adopted the same ratio of raw materials as ours to sinter the ceramics. They used 0.5 μm B₄C powders and 0.5–10 μm Ti₃SiC₂ powders as raw materials, and their process was 1850 °C hot pressing sintering for 30 min. Compared with their work, the particle size of the raw powders we used had a greater difference in size (B₄C: 1 μm; Ti₃SiC₂: <74 μm), meaning aggregates were formed more easily, which are beneficial for toughness but unfavorable for strength [12]. Besides, our ceramics had a higher relative density and graphite content, which may be due to the evaporation of Si at a higher sintering temperature, and the electric field [13]. In the analysis of XRD data, only graphite matched with the existing peaks, and the diffraction peaks of other existing forms of carbon did not correspond with the peaks in the experimental data. Moreover, the sintering temperature was 1900 °C, at which point the amorphous carbon would also be graphitized. Therefore, we infer that the lamellar phase in Figure 4b is graphite. Graphite phase exists at the grain boundaries of TiB₂ and B₄C, which reduces the bonding strength of the interface and has an adverse effect on the hardness and strength of the composite ceramics. This is the reason why the flexural strength of BT30 is at a low level in Table 3. At the same time, the existence of graphite can limit grain growth. In the cooling process, microcracks are produced under the effect of interfacial stress produced by different thermal expansion coefficients, and the cracks propagate along the interlayer of graphite during fracture, resulting in lamellar pull-out. This lamellar pull-out process greatly prolongs the path of crack propagation, consumes the energy of cracks, and is beneficial to improving the fracture toughness. It can be seen from Figure 4b that there are traces of particle pull-out and lamellar graphite pull-out in the fracture surface of BT30, which is one of the important reasons for the toughening of the composite.

Figure 5 shows the variation of the relative density and hardness of B₄C–TiB₂ composite ceramics with the content of Ti₃SiC₂. Due to the slight evaporation of silicon [13], the density of the composites is slightly higher than the theoretical density. The evaporation capacity of Si increases with the increase in Ti₃SiC₂ content, resulting in an increase in the relative density. The hardness decreases with the increase in Ti₃SiC₂ content, because the proportion of the second phase with lower hardness, especially graphite, increases.

Figure 6 shows the variation of flexural strength and fracture toughness of B₄C–TiB₂ composite ceramics with additive Ti₃SiC₂ content. When the content of Ti₃SiC₂ is in the range of 20 vol.%–30 vol.%, the flexural strength and fracture toughness are positively correlated with the content of Ti₃SiC₂. Comparing Figure 7a–c, it can be seen that the region as shown in Figure 7e becomes larger and more numerous with the increase in Ti₃SiC₂ content. In this region, TiB₂, graphite and B₄C intersect each other to form a three-dimensional interpenetrating network. During the crack propagation process, multiple two-phase interfaces must be bypassed to disperse into more small cracks and a large number of changes in the propagation direction. The pull-out mechanism of graphite also plays an important role in this agglomeration area. However, this area does not exist in isolation. Every small network links to each other, forming a large network structure covering the whole composite. At the same time, the network divides the B₄C concentration area into small parts and surrounds them, so that there is no large area of a continuous B₄C phase in the composites. The large area of a continuous B₄C phase is very unfavorable to the toughness of the composite. With the increase in interlacing degrees of TiB₂, SiC, graphite and B₄C, the cracks need to bypass more two-phase interfaces, change direction more often, and disperse into more small cracks, which can consume a lot of energy of crack propagation and hinder the spreading of cracks. Therefore, the fracture toughness of B₄C–TiB₂ composite ceramics is greatly improved by the overall three-dimensional interpenetrating network structure.

However, when the content of Ti₃SiC₂ is 35 vol.%, as shown in Figure 7d, a large TiB₂–SiC agglomerated area appears in the BT35 (among which the dark gray, medium gray and light gray phases are B₄C, SiC and TiB₂, respectively). There is a bad stress effect in the multiphase mixing region with more SiC. As shown in Figure 7f, many microcracks due to mismatching of the thermal expansion coefficient of SiC and TiB₂ can be found in this region. These microcracks are conducive to the toughening of the material, but result in a significant decrease in the bending strength.

4. Conclusions

Ultra-high toughness B₄C–TiB₂ composite ceramics were prepared by the SPS method at 1900 °C. The content of additive Ti₃SiC₂ has a great influence on the microstructure and mechanical properties. With the increase in Ti₃SiC₂ content, the relative density and fracture toughness of the material increase, while the hardness decreases, and the flexural strength first increases and then decreases. When the content of Ti₃SiC₂ is 30 vol.%, B₄C–TiB₂ composite ceramics have the highest bending strength and the best comprehensive mechanical properties: hardness 27.28 GPa, bending strength 405.11 MPa, and fracture toughness 18.94 MPa·m^1/2. The fracture mode of the material is a mixture of transgranular fracture and intergranular fracture. The existence of the graphite phase has a negative effect on the hardness and flexural strength of B₄C–TiB₂ composite ceramics, but it is beneficial to the fracture toughness. The main reason for obtaining high fracture toughness is the formation of a three-dimensional interpenetrating network covering the whole composite.