Contribution of boundary non-stoichiometry to the lower-temperature plasticity in high-pressure sintered boron carbide

doi:10.21203/rs.3.rs-2496747/v1

Download PDF

Article

Contribution of boundary non-stoichiometry to the lower-temperature plasticity in high-pressure sintered boron carbide

https://doi.org/10.21203/rs.3.rs-2496747/v1

This work is licensed under a CC BY 4.0 License

Journal Publication

published 21 Aug, 2023

Read the published version in Nature Communications →

Version 1

posted

You are reading this latest preprint version

Non-oxide ceramics exhibit significant advantages in high-tech industry, but their applications have been limited by the brittle nature especially at low to moderate temperatures. It is a great challenge to improve the ceramic plasticity while maintaining the high-temperature strength through the classical strategy, which generally includes decreasing grain size to several nanometers or adding ductile binder phase. Herein, the plasticity of fully dense boron carbide (B₄C) has been greatly enhanced due to the boundary non-stoichiometry induced by high-pressure sintering technology. The effect decreased the plastic deformation temperature of B₄C by 200°C compared to that of conventionally-sintered specimens. Promoted grain boundary diffusion is found to enhance grain boundary sliding, which dominate the lower-temperature plasticity. In addition, the as-produced specimen maintained extraordinary strength before the occurrence of plasticity. The study provides a new and efficient strategy by boundary chemical change to facilitate the plasticity of ceramic materials.

Physical sciences/Materials science/Structural materials/Ceramics

Physical sciences/Engineering/Mechanical engineering

lower-temperature plasticity

boundary non-stoichiometry

grain boundary sliding

boron carbide

advanced ceramics

Non-oxide ceramics such as boron carbide (B₄C) have been widely applied in many advanced structures because of their advantages, which include high hardness, low density, high melting point, excellent chemical stability and wear resistance. However, their application scope is still challenged because of their intrinsic brittleness with limited plasticity, especially within the low to moderate temperature range [1–6].

The lower temperature plasticity of ceramics combined with high strength might improve their in-service reliability through metal-like behavior. Two strategies are generally chosen to improve plasticity at relatively low temperatures [7–10]: (1) decreasing grain size to several nanometers to promote grain boundary sliding; (2) accelerating grain boundary diffusion with low melting point sintering additives. However, it is difficult to sinter dense nanoceramics with grains of less than 10 nm. In addition, secondary phases such as glass and oxide impurities [7, 8] in structural ceramics decrease high-temperature strength significantly, thus degrading a key mechanical property. As examples, Y₂O₃-Al₂O₃-MgO was added to nano-grained Si₃N₄ ceramics as a sintering additive, with the as-prepared ceramics exhibiting superplasticity at 1650°C, with a flow stress of only 4 MPa [7]. Nanocrystalline silicon carbide doped with boron and carbon was found to show a superplastic elongation of > 140% at 1800°C with a yield stress of 80 MPa [8]. In Imamura et al.’s studies about high-temperature deformation of 3Y-TZP, an addition of only 1 wt% amorphous silicate was found to decrease the flow stress by 50% at 1500°C [9].

Recently, it has been proposed that a high density of pre-existing defects can assist the occurrence of low-temperature plasticity in oxide ceramics through a flash sintering method, which could also limit grain growth to obtain fine structures. Via this new strategy, it was found that the enhanced plasticity of ZrO₂ and TiO₂ at low temperature, ranging from room temperature to 400°C, was mainly attributed to the high dislocation density, nanoscale stacking faults and point defects formed during the flash sintering process [11–12].

It has been reported that the sintering of ceramics under high pressure with plastic deformation as the dominant sintering mechanism also favors a higher density of defects and further limits grain growth [13–16]. Ji et al. fabricated fully dense B₄C with stacking faults and twins under a pressure of 80 MPa at 1700°C in 5 min by spark plasma sintering (SPS) [13]. Subsequent studies under a higher pressure of 1.5 GPa further facilitated the fabrication of 3YSZ nanoceramics with a grain size of 60 nm [14]. Very recently, ZrB₂ specimens with finer grains were obtained under 15 GPa at 1450°C, with the high dislocation density and fine grains induced by ultra-high pressure contributing to the enhancement of mechanical and oxidation-resistant properties [15]. Nano-twinned diamond sintered under the condition of 20 GPa and 2000°C, exhibiting an average twin thickness of 5 nm, achieved unprecedented hardness and stability [16]. Moreover, high-pressure sintering technology could lead to rough grain boundaries [13], alter crystal structure [16] or even chemical equilibrium at grain boundaries, which result in improvement in the intrinsic properties such as high-temperature yield stress. Therefore, high-pressure sintering has been chosen in the present study as a method of generating better low-temperature plasticity.

B₄C is an interesting material for various applications, including light-weight armour, cutting tools and nuclear reactor components [17–20]. Additionally, as a non-stoichiometric ceramic, B₄C has a highly tunable B:C ratio, which suggests a promising opportunity to ‘customise’ its intrinsic properties by breaking the chemical equilibrium of 4:1[21–23].

In this work, the high-temperature mechanical properties of B₄C ceramics fully densified by high-pressure low-temperature (HPLT) sintering technology were explored. Compared to reference specimens sintered under the conventional low-pressure high-temperature (LPHT) condition, the HPLT-B₄C ceramics exhibited plasticity at much lower temperatures, which was considered to relate to the boundary non-stoichiometry.

Flexural behaviour at elevated temperatures

Figure 1 shows the flexural deformation behavior and the trend from brittle behavior to ductility as the testing temperature increased. The temperature dependence of flexural strength of both HPLT-B₄C and LPHT-B₄C firstly increased with temperature prior to decreasing again as plasticity began to occur. The flexural strength of HPLT-B₄C increased from 608 MPa to 822 MPa within the range of 1100–1400°C and was still higher than 600 MPa before significant plasticity occurred at 1600°C. After plasticity occurred, the strength reduced to 385 MPa at 1700°C. The flexural strength of LPHT-B₄C increased from 580 MPa to 821 MPa when the temperature increased from 1300 to 1600°C. Significant plasticity first occurred at 1800°C and the flexural strength decreased to 190 MPa at 1900°C. At the initial stage, the increase in flexural strength with temperature within a moderate temperature range might be ascribed to the high temperature-activated relaxation of thermal residual stress [24, 25]. However, after the brittle-ductile transition temperature was reached, a decrease in flexural strength occurred due to the decrease in yield stress at high temperatures.

The photographs and microstructures of HPLT-B₄C and LPHT-B₄C after flexural deformation at 1700°C and 1900°C are shown in Figs. 1(c-d) and 1(i-j). At high temperature, both HPLT-B₄C and LPHT-B₄C specimens exhibited a similar plastic morphology to that observed with metals (Figs. 1(c) and 1(i)). The brittle-ductile transition of HPLT-B₄C and LPHT-B₄C occurred at 1600°C and 1800°C respectively. Substantial plastic deformation of HPLT-B₄C and LPHT-B₄C occurred at 1700°C and 1900°C respectively, revealing that the plastic deformation temperature of HPLT-B₄C was 200°C lower than that of LPHT-B₄C. This difference in temperature of ceramic plasticity may be related to relative density, grain size, internal stress, micro-morphology and grain boundary structure [26–28]. Both materials possessed relative densities above 99%. The average grain sizes of HPLT-B₄C and LPHT-B₄C were similar, 2.5 µm and 3.1 µm respectively (Figs. 1(b) and 1(h)). Therefore, the plasticity at lower temperatures was not mainly attributed to the relative density or grain size. Although previous studies have ascribed the variation of high-temperature plasticity to activated relief of thermal residual stresses [24], the testing results of annealed samples in the present study (Fig. S1 and Table S1) indicated that the plastic deformation temperature was independent of residual stresses. Therefore, the decrease in plastic deformation temperature might be related to special micro-morphology and grain boundary structure formed by the high-pressure low-temperature sintering technology.

As for HPLT-B₄C-1700-0.5, named and listed in Table 1, with the increase in the strain during the high-temperature plastic deformation, the rough failure area induced by microcrack propagation was observed in the surface center where under the maximum tensile strain during the test (Fig. 1(d) and 1(g)). The location was determined from the finite element analysis as shown in Fig. S2. Intergranular fracture was detected clearly near the crack tip (Fig. 1(d)) and a high density of pores was evident on the surface. The high-temperature fracture of B₄C was induced by nucleation, propagation, blunting and coalescence of cracks. These pores could favor the micro-void coalescence ahead of the crack tips. Plasticity at this temperature may relieve stress at the crack tips, facilitating the growth and coalescence of cracks [26, 29]. The morphology revealed that intergranular fracture was the dominant failure mode of HPLT-B₄C in the present work, while LPHT-B₄C-1900-0.5 behaved differently in terms of fracture mode and morphology. The fracture occurred both at the grain boundary and inside grains (Fig. 1(j) and 1(m)), indicating that the failure was caused by both transgranular and intergranular fractures. It is important for these different fracture modes of the two B₄C to be further investigated, as the mechanism may directly lead to the lower-temperature plasticity of the HPLT-B₄C.

Despite the different fracture modes of HPLT-B₄C and LPHT-B₄C as stated above, neither of them is similar to metallic materials which generally exhibit elongated grains after deformation. The electron back scattered diffraction (EBSD) results in Fig. S3 and Fig. S4 show that the grains of the B₄C specimens, either with or without deformation, were equiaxed.

Boundary non-stoichiometry of high-pressure sintered B₄C

It has been reported that the intrinsic properties of B₄C at grain boundaries or within the grains are altered by breaking the stoichiometric ratio of 4:1 [22]. Therefore, the stoichiometry and crystal structure of grain lattice and grain boundary in as-sintered B₄C (without high-temperature deformation) are investigated here.

Electron energy loss spectroscopy (EELS) spectra were used as a preliminary method to study the B₄C structure along the grain boundary and within the grains. The EELS spectra in Fig. 2(a)-(b) showed identical structures on grain boundaries and within the grains in LPHT-B₄C. However, this was not the case for HPLT-B₄C. As shown in Figs. 2(g)-2(h), EELS spectra show a transition from multiple, discrete peaks in the range of 193–205 eV in the lattice of HPLT-B₄C, to a singe smooth curve in the grain boundary. This trend in EELS spectra was consistent with that reported by Xie et al. with the increase in B/C ratio beyond 4:1, confirming the chemical element non-stoichiometry at the grain boundary of HPLT-B₄C [22].

To further confirm the unique crystal structure, a spherical aberration corrected transmission electron microscope (ACTEM) study was performed at the grain boundary and lattice area of both B₄C ceramics (Fig. 2(c-f), 2(i-r)). Atomic models of the ideal B₄C crystal structure in the [211] crystallographic direction yielded the best view [30]. The icosahedra and the chains were projected onto each other. The conventional high-temperature sintering method led to a straight equilibrium grain boundary (Fig. 2(d)). LPHT-B₄C ceramics showed consistent chemical bonding and atomic occupation inside the grain and at the grain boundary (Figs. 2(c) and 2(e-f)), exhibiting perfect structures of icosahedral clusters and CBC chains. The grain boundary and inner grains of B₄C ceramics had the same intrinsic structures and properties, thus leading to the coexistence of transgranular and intragranular fractures in the flexure test at 1900°C, as indicated by the fracture mode shown in Fig. 1.

The atomic arrangement of inner lattice of HPLT-B₄C (Fig. 2(i)) also exhibited the same perfect B₄C structure as that of LPHT-B₄C. However, high resolution TEM images (Fig. 2(j)) indicated that the non-stoichiometry rough grain boundary formed during high-pressure sintering. The atomic images (Figs. 2(m)-2(o)) clearly confirmed that the high-pressure and low-temperature sintering technology could induce the vacancy of carbon in the CBC chains in the grain boundary yielding region. Meanwhile, the atomic configuration of boron carbide with higher B concentrations along grain boundary was also observed (Figs. 2(p)-2(r)). Strong atomic arrangement distortion was observed along the grain boundaries, like the Rasim model [21]. Both B concentration and C atom vacancy induced a larger B/C atomic ratio of the crystal structure near grain boundaries than inner grains.

The result illustrated direct evidence that the high-pressure and low-temperature sintering technology could alter the chemical composition at grain boundaries and may affect the grain boundary performance of B₄C ceramics.

Plastic deformation characterization in the tensile surface

Post-deformation TEM analysis with B₄C ceramics after high-temperature flexural tests was performed in the present study. A TEM bright field image of HPLT-B₄C (Fig. 3(a)) shows the formation of high-density defects including large-sized twins, stacking faults and nano-twins. In Fig. 3(b), the angle between matrix planes and twins is 73°. Figure 3(c) shows a scanning transmission electron microscopy (STEM) image collected by a high-angle annular dark-field (HAADF) detector in ACTEM. Based on the arrangement of the icosahedral clusters (which are the unique atomic chains in the rhombohedral structure), it is revealed that the two crystals separated by the interface have a perfect mirroring relationship while the inclination angle is 73°. As shown in Figs. 3(d) and 3(e), in addition to large-sized twins, nano-twins were also observed in HPLT-B₄C-1900-0.5 grains. According to the TEM images in Fig. S5, nano-twins with around 1 ~ 2 nm width had preexisted after high-pressure sintering prior to high-temperature flexural tests.

By contrast, in B₄C ceramics sintered by conventional high-temperature method, with 150 grains carefully investigated, it was found that only large-sized twins without nanoscale defects could be observed (Figs. 3(j)-3(l)), which might also indicate that the nano-twins in deformed HPLT-B₄C initially formed during high-pressure sintering rather than deformation. In addition, as indicated in EBSD results shown in Fig. S4, the density of large-sized twins basically remained constant after high-temperature plastic deformation.

Neither the large-sized twins nor nano-twins directly led to the change in fracture modes and dominated the lower-temperature plasticity, but might still make a positive contribution to the high-temperature plastic deformation.

Here geometric phase analysis was applied as an indirect way to confirm the effect of twinning on the force field. The results based on the (101) lattice reflections of a 73°<211 > twin interface inside the HPLT-B₄C-1900-0.5 after flexural deformation are shown in Figs. 3(f)-4(i), where ε_xx, ε_xy, and ε_yy are horizontal axis strain, shear strain and vertical axis strain, respectively [31, 32]. The local strain fields showed tensile strain (orange color) along the horizontal axis and compressive strain (blue color) with the vertical axis component. There is also significant shear strain. The normal strain tensors were calculated as a symmetric matrix so as to relieve lattice mismatch at the twin plane. The fracture and large strain of B₄C ceramics are associated with the twin interface [22, 31, 33]. The compressive strain exhibited that, due to the improvement in shear resistance, the 73°<211 > twins in B₄C presented a positive influence on strength, thus maintaining the mechanical properties at relatively high temperatures of ~ 1600°C. The Raman results shown in Fig. S6 and Table S2 confirmed that the structure of B₄C ceramics did not become amorphous after high-pressure and high-temperature deformation, demonstrating the stability of the B₄C ceramics at high temperatures [34].

Plastic deformation mechanisms at high temperature

To further investigate the plastic deformation mechanism, typical temperatures of 1700°C and 1950°C for HPLT-B₄C and LPHT-B₄C respectively were selected in the study. The dependence of the diffusion-controlled deformation kinetics on temperature, stress and strain rate can be expressed as follows [35]:

$$\dot{\epsilon }=\frac{ADGb}{T}{\left(\frac{b}{d}\right)}^{p}{\left(\frac{\sigma }{G}\right)}^{n}{e}^{-\frac{Q}{RT}}$$

where $\dot{\epsilon }$ is the strain rate; A is a constant; D is the diffusion coefficient; b is the Burgers vector; T is the temperature; d is the grain size; p is the grain size exponent; $\sigma$ is the stress; G is the shear modulus; Q is the activation energy; R is the gas constant; n is the stress exponent and is used to identify the dominant mechanism.

Based on the constitutive equations, assuming that the temperature is constant, the relation between the stress $\sigma$ and strain rate $\dot{\epsilon }$ is simplified as [36]:

$${\sigma }=\text{B}{\dot{\epsilon }}^{\frac{1}{n}}$$

where B is a constant.

The activation energy Q was determined from the slope of the line shown in Figs. 4(c) and 4(f), which plot ln(σ⁻ⁿT) at the same strain rate as a function of 1/T in the high-temperature deformation process.

The stress exponents n were calculated with the stress-strain curve obtained under various strain rates and a constant temperature. The activation energy Q was determined from the stress-strain curves obtained under different temperatures and a constant strain rate. The results showed that the plastic yield stress increased with the strain rate at a constant temperature.

HPLT-B₄C and LPHT-B₄C exhibited identical stress exponents (n = 2.1), suggesting that the plastic deformation of HPLT-B₄C and LPHT-B₄C was dominated by grain boundary sliding [33]. This is consistent with the observed intergranular fracture more in HPLT-B₄C when significant plasticity occurred. The values of stress exponent in the present work (n = 2.1) were smaller than that from the available literature (n = 3), which was associated with deformation by dislocation motion [37]. The main reason was that the specimens in the current study possessed much smaller grain size, which favored grain boundary sliding with n = 2 rather than power law creep with n = 3. In addition, the heavy twinning introduced during processing of the HPLT-B₄C would suppress dislocation motion [38]. Moreover, the present study mainly focused on the plastic deformation rather than creep conditions, with the strain rate of high-temperature deformation adopted here (within the range of 2×10^− 4~8×10^− 4 s^− 1) being much higher than that in previous studies (10^− 6 s^− 1).

The values of Q for HPLT-B₄C and LPHT-B₄C were respectively 785 kJ/mol and 1126 kJ/mol, suggesting an acceleration of the dominant diffusion mechanism of HPLT-B₄C relative to LPHT-B₄C. Therefore, it indicated that the high-pressure and low-temperature sintering technology promoted grain boundary diffusion-controlled plasticity at a relative low temperature. The similar strengths and fracture modes of the two materials in the brittle fracture regime indicates that the grain boundaries were not significantly weakened by the HPLT treatment.

The activation energy values calculated from flexural deformation in the present study were higher than those in previous studies, which ranged from 385 kJ/mol to 632 kJ/mol [37, 39]. Effects of different sintering technologies, grain sizes measurement parameters and plastic deformation mechanism could be responsible for this increase in activation energy. As noted above, the dominant deformation mechanism, the strain-rate regime and the microstructures are significantly different in this work to the previous studies.

The above investigations revealed that the dominant mechanism of high-temperature plastic deformation was grain boundary sliding, rather than a direct effect of defects in nanoscale or larger scale.

In a pressure-assisted sintered B₄C sample, both asymmetric and symmetric twins were detected (Fig. S5), consistent with the atomic ratio of B to C in the grain lattice (4:1) [40], while the observed non-stoichiometry resulted in a larger B:C ratio at grain boundaries. According to the previous study [22], when the atomic ratio of B to C increased above 4:1, the intrinsic hardness slightly decreased. Therefore, the hardness of HPLT-B₄C along grain boundaries was correspondingly lower than that inside the grains. The yield stress of B₄C ceramics at high temperatures was positively correlated with the hardness because there was no obvious grain deformation or texture during plastic deformation [25, 41]. Therefore, the grain boundary diffusion was promoted. As shown in Fig. 5, the plastic yielding of HPLT-B₄C started at grain boundaries with slightly lower yield stress, leading to intergranular fractures. In contrast, for LPHT-B₄C, the crack needed higher temperature to conquer the relative higher yield stress, which was identical between grain and boundary, and resulted in the mixture of transgranular and intergranular fractures. Ultimately, it was revealed that the unique crystals with boundary non-stoichiometry formed by high-pressure and low-temperature sintering technology greatly contributed to the improvement in plasticity at lower temperatures.

In summary, flexure experiments at high temperatures showed that B₄C sintered under a high pressure had the better low-temperature plasticity. The plastic deformation temperature of the B₄C ceramics sintered with high-pressure sintering technology was 200°C lower than that of B₄C ceramics fabricated by conventional high-temperature sintering, with the plastic deformation mainly dominated by grain boundary sliding. Boundary non-stoichiometry in HPLT-B₄C decreased the activation energy and thus favored lower-temperature plastic deformation. By regulating the state of the grain boundaries, the high-pressure sintering technology could stimulate the lower-temperature plasticity of ceramics while maintaining high strength before the plasticity occurred. The new technology proposed in the present study could enable non-oxide ceramics to meet the engineering application requirements in a broader range of temperatures.

Sintering of B ₄ C. As-received B₄C powder (97% purity, Mudanjiang Diamond Boron Carbide Co., Ltd., China) with a median particle size of 2.5 µm was used as the starting material. B₄C ceramics were sintered by SPS under two different conditions: (i) 80 MPa and 1800°C for HPLT-B₄C, and (ii) 20 MPa and 2100°C for LPHT-B₄C. For both conditions, 22g raw powder was first poured into a cylindrical graphite die with an inner diameter of 50 mm, prior to the following sintering process in a SPS apparatus (HPD 60, FCT, Germany). and the temperature increased gradually from room temperature to the desired temperature with a heating rate of 100°C/min, and then maintained for 10 mins. The sintering pressure was applied during the entire heating process. After the completion of the above heating process, the pressure was immediately released to 10 MPa at the beginning of the cooling process. In order to minimise the effects of residual stresses, the temperature first decreased to 1600°C at a constant cooling rate of 10°C/min, followed by natural cooling to room temperature.

High-temperature flexural tests. Test bars (2.2 mm × 3.8 mm × 25 mm) for three-point flexural measurements at elevated temperatures were cut from the sintered pellets using diamond saws. The high-temperature three-point flexural strengths were measured by a ceramic test system with New Equipment Machinery Systems (NEMS) in a high temperature vacuum furnace (MTS, AGS-X, SHIMADZU, Japan) with a span of 16 mm according to literatures and ASTM C1211 standard [42–44]. The stress and strain were calculated from the loading and displacement data using Eqs. (3) and (4) [42]. The strain rate was estimated from the crosshead speed according to Eq. (5). The test conditions of HPLT-B₄C and LPHT-B₄C are listed in Table 1.

$$\sigma =\frac{5FL}{4a{c}^{2}}$$

$$\epsilon =\frac{8cl}{{L}^{2}}$$

$$\dot{\epsilon }=\frac{8cv}{{L}^{2}}$$

where F is the applied load; L is the outer span (16 mm); a and c are the width and thickness of the test bar; l is the crosshead displacement; v is the crosshead speed. A minimum of three specimens were measured at each test temperature.

Table 1

Summary of Sintering and Test Conditions of B₄C Ceramics
Samples	Sintering Temperature (°C)	Sintering Pressure (MPa)	Test Temperature (°C)	Crosshead Speed (mm/min)
HPLT-B₄C-1100-0.5	1800	80	1100	0.5
HPLT-B₄C-1300-0.5	1800	80	1300	0.5
HPLT-B₄C-1400-0.5	1800	80	1400	0.5
HPLT-B₄C-1500-0.5	1800	80	1500	0.5
HPLT-B₄C-1600-0.5	1800	80	1600	0.5
HPLT-B₄C-1700-1.0	1800	80	1700	1.0
HPLT-B₄C-1700-0.5	1800	80	1700	0.5
HPLT-B₄C-1700-0.25	1800	80	1700	0.25
HPLT-B₄C-1800-0.5	1800	80	1800	0.5
HPLT-B₄C-1900-0.5	1800	80	1900	0.5
HPLT-B₄C-2000-0.5	1800	80	2000	0.5
LPHT-B₄C-1300-0.5	2100	20	1300	0.5
LPHT-B₄C-1500-0.5	2100	20	1500	0.5
LPHT-B₄C-1600-0.5	2100	20	1600	0.5
LPHT-B₄C-1700-0.5	2100	20	1700	0.5
LPHT-B₄C-1800-0.5	2100	20	1800	0.5
LPHT-B₄C-1900-0.5	2100	20	1800	0.5
LPHT-B₄C-1950-1.0	2100	20	1950	1.0
LPHT-B₄C-1950-0.5	2100	20	1950	0.5
LPHT-B₄C-1950-0.25	2100	20	1950	0.25
LPHT-B₄C-2000-0.5	2100	20	2000	0.5

Microstructural characterization. BSE images (Fig. 1b, 1d, 1h, and 1j) were recorded with a field emission scanning electron microscope (FESEM, Merlin, Zeiss, Germany) and taken combined with the signals from a multi-detector system (Bruker ARGUS ™), which was positioned below the phosphor screen to collect forward scattered electrons. Detailed microstructures (Figs. 3a, 3b, 3d, and 3f-l) were examined using high resolution transmission electron microscopy (HRTEM, Talos-F200S, FEI, USA), with the atomic-scale characterization (Figs. 3c, 3e, 2c-2f, 2i-2l and 2p-2r) performed in a double spherical aberration corrected transmission electron microscope (ACTEM, Titan Cubed Themis G2 300, FEI, USA). The B and C elemental concentrations from grain boundary to within the grain (Figs. 2a, 2b, 2g and 2h) were investigated from the chemical bonding information using electron energy loss spectroscopy (EELS, GIF QuantumER) in the TEM. TEM specimens were prepared from as-processed B₄C by mechanical thinning and subsequent ion milling (EM TXP, Leica, Germany).

Acknowledgements

The authors thank Prof. Richard Todd from University of Oxford for the academic guidance in deformation mechanism and language polish. The authors acknowledge the use of characterisation facilities within the David Cockayne Centre for Electron Microscopy, Department of Materials, University of Oxford. The authors also thank Prof. Danyu Jiang and Prof. Xingang Wang from Shanghai Institute of Ceramics, Chinese Academy of Sciences for the help in high-temperature flexural strength testing. This work was financially supported by the National Natural Science Foundation of China (92163208, 51902233, 52022072, 51832003), the National Key Research and Development Plan of China (2021YFB3701400). The ACTEM work was performed at the Nanostructure Research Center, which is supported by the Fundamental Research Funds for the Central Universities (2021III016GX).

Author contributions

W.J. and Z.Y.F. proposed the idea and designed the project. Z.Y.F. supervised the research. H.Y.X. and W.J. carried out most experiments, analyzed data and wrote the manuscript. J.W.J. and J.L.L. helped the SEM-TKD characterization. J.S.W. and R.H.Y. carried out TEM experiments. W.M.W. and J.Y.Z. developed the theoretical model. B.T.T. and J.Z. performed crystal structure simulations and finite element analysis. All authors discussed the results and revised the manuscript.

Competing interests

The authors declare no competing interests.

Wakai, F., Kodama, Y., Sakaguchi, S. A superplastic covalent crystal composite. Nature 344, 421–423 (1990).
Kim, BN., Hiraga, K., Morita, K. et al. A high-strain-rate superplastic ceramic. Nature 413, 288–291 (2001).
Luo, W., Yan, S., Zhou, J., Ceramic-based dielectric metamaterials. Interdiscip. Mater. 1, 11–27 (2022).
Waku, Y., Nakagawa, N., Wakamoto, T. et al. A ductile ceramic eutectic composite with high strength at 1,873 K. Nature 389, 49–52 (1997).
Lai, A., Du, Z., Gan, C. L. & Schuh, C. A. Shape memory and superelastic ceramics at small scales. Science 341, 1505–1508 (2013).
Jansen, M., High Performance Non-Oxide Ceramics II. Springer Berlin Heidelberg (2002).
Wananuruksawong, R., Shinoda, Y., Akatsu, T., Wakai, F. Evolution of microstructure and intergranular glass chemistry in plastically deformed nanocrystalline Si₃N₄ Ceramics. J. Am. Ceram. Soc. 98, 178–185 (2015).
Shinoda, Y., Nagano, T., Gu, H., Superplasticity of Silicon Carbide, J. Am. Ceram. Soc. 82, 2916–2918 (1999).
Imamura, P. H., Evans, N.D., Sakuma, T., Mecaryney, M.L., High temperature tensile deformation of glass-doped 3Y-TZP, J. Am. Ceram. Soc. 83, 3095–3099 (2000).
Neuman, E. W., Hilmas, G. E., Fahrenholtz, W. G. Mechanical behavior of zirconium diboride–silicon carbide–boron carbide ceramics up to 2200°C. J. Eur. Ceram. Soc. 35, 463–476 (2015).
Li, J., Cho, J., Ding, J., Charalambous, H., Xue, S., Wang, H. Phuah, X. L., Jian, J., Wang, X., Ophus, C., Tsakalakos, T., García, R. E., Mukherjee, A. K., Bernstein, N., Hellberg, C. S., Wang, H., Zhang, X. Nanoscale stacking fault-assisted room temperature plasticity in flash-sintered TiO₂. Sci. Adv. 5, eaaw5519 (2019).
Cho, J., Li, Q., Wang, H. et al. High temperature deformability of ductile flash-sintered ceramics via in-situ compression. Nat Commun 9, 2063 (2018).
Ji, W., Rehman, S., Wang, W. et al. Sintering boron carbide ceramics without grain growth by plastic deformation as the dominant densification mechanism. Sci Rep 5, 15827 (2015).
Ji, W., Xu, H. Y., Wang, W. M., Fu, Z. Y. Sintering dense nanocrystalline 3YSZ ceramics without grain growth by plastic deformation as dominating mechanism. Ceram. Int. 45: 9363–9367(2019).
Xu, H. Y., Ji, W., Guo, W. M., Li, Y. L., Zou, J., Wang, W. M., Fu, Z. Y. Enhanced Mechanical Properties and Oxidation Resistance of Zirconium Diboride Ceramics via Grain-Refining and Dislocation Regulation. Adv. Sci. 9, 2104532 (2022).
Huang, Q., Yu, D., Xu, B. et al. Nanotwinned diamond with unprecedented hardness and stability. Nature 510, 250–253 (2014).
Zhang, W. A review of tribological properties for boron carbide ceramics. Prog. Mater. Sci. 116,100718(2021).
Thévenot, F. Boron carbide – a comprehensive review. J. Eur. Ceram. Soc. 6, 205–25 (1990).
Domnich, V., Reynaud, S., Haber, R.A., Chhowalla, M. Boron carbide: structure, properties, and stability under stress. J. Am. Ceram. Soc. 94, 3605–28 (2011).
Du, X., Zhang, Z., Wang, Y. Hot-pressing knetics and densification mechanisms of boron carbide. J. Am. Ceram. Soc. 98, 1400–1406 (2015).
Rasim. K, Ramlau, R., Leithe-Jasper, A. Local atomic arrangements and band structure of boron carbide. Appl. Organomet. Chem. 130, 6238–6243 (2018).
Xie, K. Y., Domnich, V., Farbaniec, L. Microstructural characterization of boron-rich boron carbide. Acta Mater. 136, 202–214 (2017).
Cheng, C., Reddy, K. M., Hirata, A. Structure and mechanical properties of boron-rich boron carbides. J. Eur. Ceram. Soc. 37 4514–4523 (2017)
Neuman, E.W., Hilmas, G.E., Fahrenholtz W.G. Strength of zirconium diboride to 2300°C. J. Am. Ceram. Soc. 96, 47–50 (2013).
Belon, R., Antou, G., Pradeilles, N. Mechanical behaviour at high temperature of spark plasma sintered boron carbide ceramics. Ceram. Int. 43, 6631–6635 (2017).
Neuman, E. W., Hilmas, G. E., Fahrenholtz, W. G. Ultra-high temperature mechanical properties of a zirconium diboride–zirconium carbide ceramic. J. Am. Ceram. Soc. 99, 597–603 (2016).
Wachtman, J. Mechanical Properties of Ceramics. John Wiley & Sons, Inc. New York City, New York (1996).
Guo, D. et al. Grain boundary sliding and amorphization are responsible for the reverse Hall-Petch relation in superhard nanocrystalline boron carbide. Phys. Rev. Lett. 121, 145504 (2018).
Evans, A. G., Rana, A. High Temperature Failure Mechanisms in Ceramics. Acta Metall., 28, 129–41 (1980).
Fujita, T., Guan, P., Madhav Reddy, K. Asymmetric twins in rhombohedral boron carbide. Appl. Phys. Lett. 104, 021907 (2014).
Li, J. J., Chen, J. C., Wang, H., Chen, N., Wang, Z. C., Guo, L., Deepak, F. L. In situ atomic-scale study of particle‐mediated nucleation and growth in amorphous bismuth to nanocrystal phase transformation. Adv. Sci. 5, 1700992 (2018).
Hÿtch, M., Putaux, JL. & Pénisson, JM. Measurement of the displacement field of dislocations to 0.03 Å by electron microscopy. Nature 423, 270–273 (2003).
Madhav Reddy, K., Guo, J., Shinoda, Y. et al. Enhanced mechanical properties of nanocrystalline boron carbide by nanoporosity and interface phases. Nat Commun 3, 1052 (2012).
Chen, M. W., McCauley, J. W. & Hemker, K. J. Shock-induced localized amorphization in boron carbide. Science. 299, 1563–1566 (2003).
Ruano, O. A., Wadsworth, J., Sherby O. D. Deformation of fine-grained alumina by grain boundary sliding accommodated by slip. Acta Mater. 51, 3617–3634 (2003).
Gifkins R. C. Grain-boundary sliding and its accommodation during creep and superplasticity. Metall. Trans. A 7, 1225–1232 (1976).
Moshtaghioun, B.M., García, D.G., Rodríguez, A.D., Padture, N.P. High-temperature creep deformation of coarse-grained boron carbide ceramics. J. Eur. Ceram. Soc. 35, 1423–1429 (2015).
Guo, H., Zhang. J., Mao, X. Strengthening mechanism of twin lamellas in transparent AlON ceramics. J. Eur. Ceram. Soc. 38, 3235–3239 (2018).
Abzianidze, T.G., Eristavi, A.M., Shalamberidze, S.O. Strength and creep in boron carbide (B₄C) and aluminum dodecaboride (α-AlB₁₂). J. Solid State Chem. 154 191–193 (2000).
Goddard, W. A. I., Hemker, K. J., Xie, K. Y. Atomic-Level Understanding of "Asymmetric Twins" in Boron Carbide. Phys. Rev. Lett. 115, 175501 (2015).
Li, X., Wei, Y., Lu, L. et al. Dislocation nucleation governed softening and maximum strength in nano-twinned metals. Nature 464, 877–880 (2010).
Hollenberg, G. W., Terwilliger, G. R., Gordon, R. S. Calculation of Stresses and Strains in Four-Point Bending Creep Tests. J Am. Ceram. Soc. 54, 196–199 (1971).
Smith, C. J., Ross, M. A., Nicholas, D. L. Ultra-high temperature deformation in TaC and HfC. J. Euro. Ceram. Soc. 38, 5319–5332 (2018).
ASTM C1211-18, Standard Test Method for Flexural Strength of Advanced Ceramics at Elevated Temperatures, ASTM International, West Conshohocken, PA (2018).

There is NO Competing Interest.

2.SupplementaryInformation.docx

Download PDF

Journal Publication

published 21 Aug, 2023

Read the published version in Nature Communications →

Version 1

posted

You are reading this latest preprint version

Contribution of boundary non-stoichiometry to the lower-temperature plasticity in high-pressure sintered boron carbide

Status:

Journal Publication

Version 1

Abstract

Figures

Introduction

Results

Flexural behaviour at elevated temperatures

Boundary non-stoichiometry of high-pressure sintered B₄C

Plastic deformation characterization in the tensile surface

Discussion

Plastic deformation mechanisms at high temperature

Methods

Declarations

References

Additional Declarations

Supplementary Files

Status:

Journal Publication

Version 1

Contribution of boundary non-stoichiometry to the lower-temperature plasticity in high-pressure sintered boron carbide

Status:

Journal Publication

Version 1

Abstract

Figures

Introduction

Results

Flexural behaviour at elevated temperatures

Boundary non-stoichiometry of high-pressure sintered B4C

Plastic deformation characterization in the tensile surface

Discussion

Plastic deformation mechanisms at high temperature

Methods

Declarations

References

Additional Declarations

Supplementary Files

Status:

Journal Publication

Version 1

Boundary non-stoichiometry of high-pressure sintered B₄C