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Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
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Short communication
Harmonized toughening and strengthening in pressurelessly reactivesintered Ta0.8Hf0.2C-SiC composite
Buhao Zhanga,d, Jie Yina,b, , Yihua Huanga, Jian Chena, Xuejian Liua, Zhengren Huanga,c,
State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050,
Bureau of Major R&D Programs, Chinese Academy of Sciences, Beijing 100864, China
Ningbo Institute of Materials Technology and Engineering, Chinese Academy of Sciences, Ningbo 315201, China
University of Chinese Academy of Sciences, Beijing 100049, China
Ta0.8Hf0.2C-SiC composite
Pressureless in-situ densification
Phase transformation
Ta0.8Hf0.2C-27 vol%SiC (99.0% in relative density) composite was toughened and strengthened via pressurelessly in-situ reactive sintering process. HfC and β-SiC particles were formed after reaction of HfSi2 and carbon
black at 1650 °C. Ta0.8Hf0.2C was obtained from solid solutioning of HfC and commercial TaC. The β-α phase
transformation of SiC proceeded below 2200 °C. High aspect ratio, platelet-like α-SiC grains formed and interconnected as interlocking structures. Toughness and flexural strength values of 5.4 ± 1.2 MPa m1/2 and
443 ± 22 MPa were measured respectively. The toughening mechanisms by highly directional growth of discontinuous α-SiC grains were crack branching, bridging and deflection behaviors.
1. Introduction
Ta1-xHfxC, belonging to the family of ultra-high temperature ceramics (UHTCs), has attracted considerable attention due to high melting
point, superior thermal and chemical stabilities [1]. The solid solution
formation not only contributes to densification by decreasing the diffusion activation energy across grain boundaries [2], but may also
improve mechanical properties than their monolithic carbide counterparts: relationship between hardness and x values of Ta1-xHfxC (x = 0,
0.2, 0.5 and 0.8) was reported to vary as an open-down quadratic
function, which could be attributed to solid solution strengthening effect [3]. Notably, Ta0.8Hf0.2C was proved to have the highest melting
point 3942 ± 82 °C among the Ta1-xHfxC solid solutions [4].
Ta0.8Hf0.2C is the candidate material for thermal protection systems
(TPS) as sharp leading edges, nose caps and flight control components
of aerospace vehicles that can be exposed to extreme environments with
temperatures exceeding 2000 °C by dissociated air operate at hypersonic speeds.
Owing to its extremely strong covalent bonding, Ta0.8Hf0.2C is difficult to densify. Advanced techniques including hot pressing, spark
plasma sintering [5] and hot isostatic pressing [6] have been attempted
to enhance the sintering driving force. Pressureless sintering (PLS) is
beneficial to fabricate cost-effective and near-net-shape ceramic
components. Researches on the field-less densification techniques of
Ta0.8Hf0.2C were rather scarce. Quite recently, we reported pressureless
densification of Ta0.8Hf0.2C-based composites with fine mechanical/
thermal properties [7].
The critical application of Ta0.8Hf0.2C-based composites as thermal
protection systems (TPS) is restricted by the poor thermal shock resistance, including thermal shock fracture and damage resistance. The
maximum thermal shock fracture parameter (R) is one of crucial importance to impede the fracture inside bulk ceramics [8]. Thermal
shock damage resistance is considered on the basis of toughness (KIC),
with the pre-existing cracks after fracture initiation, when the maximum stress intensity factor Kmax reaches the material toughness KIC
Therefore, UHTCs need to improve their damage tolerance via increasing both strength and toughness for TPS application [9]. To meet
such demand, incorporating a hard secondary phase to strengthen, together with further extrinsic toughening [10], such as crack deflection
and bridging is an effective strategy. Investigations are needed for the
optimization of Ta0.8Hf0.2C-SiC mechanical performance, typically insitu strengthening and toughening.
SiC, introduced in-situ inside the (Ta,Hf)C bulk, were tailored to be
platelet-like with high aspect ratio of 15.6 after soaking 2 h at 2200 °C.
Homogeneous HfC and SiC formed after reaction between HfSi2 and
Corresponding authors at: State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of
Sciences, Shanghai 200050, China.
E-mail addresses: (B. Zhang), (J. Yin), (Z. Huang).
Received 24 May 2018; Received in revised form 16 August 2018; Accepted 17 August 2018
0955-2219/ © 2018 Published by Elsevier Ltd.
Please cite this article as: Zhang, B., Journal of the European Ceramic Society (2018),
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
B. Zhang et al.
Fig. 1. (a) Thermodynamics analysis of the reactions for TaC-HfSi2-carbon black system; XRD patterns of (b) HfSi2 and carbon black reaction and (c) HfSi2 and carbon
black reaction in TaC matrix from [33°, 43°] at different temperatures.
hardness was measured by the indentation technique (Model 300,
Tukon, Canton, MA, USA) using a load of 3 kg and dwell time of 10 s.
The indentation fracture resistance (KIC) was calculated on the basis of
the equation as reported by Evans and Charles [12].
carbon black at 1650 °C. Ta0.8Hf0.2C was generated via inter-diffusion
among raw TaC and in-situ HfC. The phase and microstructure evolution of SiC proceeded during the high-temperature densification progress. The in-situ toughening mechanism was investigated.
KIC = 0.16(c/a)−1.5∙H∙a1/2
2. Experimental procedure
Where KIC is fracture toughness, a is the half average length of the diagonal of the Vickers indentations (μm), c is the average length of the
cracks obtained in the tips of the Vickers indentations (μm), and H is the
Vickers hardness.
TaC and HfSi2 powders (Haotian Nano Technology Co., Ltd,
Shanghai, China) had a purity of 99.9% and an average particle size of
1 μm. The median size of carbon black (Anyang Delong Chemical Co.,
Ltd., Henan, China) was ∼200 nm. Phenolic resin (P. R.) (Shanghai
QiNan Adhesive Material Factory, Shanghai, China) with a carbon yield
of 50 wt% was used as binder. The oxygen contents of TaC and HfSi2
were analyzed using infrared method after fusion under an inert gas
atmosphere (TC-600C, LECO Instrument LD, MI, USA). TaC, HfSi2 and
carbon contents were precisely calculated for obtaining Ta0.8Hf0.2C27 vol% SiC in stoichiometric and volume ratio based on the reaction
below [11]:
HfSi 2 + 3C = HfC + 2SiC ΔG = −167385−7.456T
3. Results and discussion
To obtain desired properties of (Ta,Hf)C-SiC composite, the chemical reactions in TaC-HfSi2-C system were investigated. Oxide impurities were coated on the surfaces of raw powders, eg: Ta2O5 on TaC
(oxygen contents: 0.59 wt%), and HfO2 and SiO2 on HfSi2 (oxygen
contents: 1.73 wt%). The complexity of the chemical reactions in TaCHfSi2-C system increases due to the surface oxide impurities. Carbon
was an effective reducing agent. Following chemical reactions should
be considered [13,14]:
Where ΔG : Gibbs’ free energy (J/mol), and T: Temperature (K). The
starting powders were blended with 4 wt% phenolic resin in ethanol,
and ball-milled for 4 h using a planetary mill (QM-3SP4, Nanjing NanDa
Instrument plant, Nanjing, China) with WC balls (5 mm in diameter)
and Teflon-coated tanks (Φ 100 mm × 120 mm) at a speed of 400 rpm.
Wear debris contamination was ≤0.2 wt % after milling. The slurries
were dried, crushed, sieved and pressed uniaxially in a steel die at
45 MPa (6 mm × 8 mm × 48 mm), followed by cold isostatic pressing
at 280 MPa for 300 s. The green compacts were pressurelessly sintered
at 1650 °C and 2200 °C for 0, 1 and 2 h in argon (purity 99.999% with
O2 ≤ 1.5 ppm), at a heating rate of 10 °C min−1 in a graphite resistance
furnace (Zhuzhou Norbert High Temperature Instrument Ltd. Co.,
Phase analysis was performed by X-ray diffraction (XRD; Ultima IV
diffractometer, Rigaku, Tokyo, Japan) with Cu Kα radiation
(λ = 1.5406 Å). The microstructure and crack introduced by indentation in hardness test on polished surface were analyzed by scanning
electron microscopy (SEM; Magellan 400, FEI, Hillsboro, American)
equipped with EBSD (INCA SERIES, Oxford Instrument, UK). The
electron backscatter diffraction pattern (EBSP) was acquired at an angle
of 60° and an acceleration voltage of 15 kV, and Aztec software was
used to clarify the phase. In transmission electron microscopy (TEM;
Tecnai G2 F20, FEI Co., Hillsboro, USA) observation, high-resolution
transmission electron microscopy (HRTEM) was conducted. The bulk
density was measured by the Archimedes method. The theoretical
density was determined based on the rule of mixtures (densities of
Ta0.8Hf0.2C and SiC are 14.05 and 3.21 g cm−3 respectively).
Three-point bending strength (3 mm × 4 mm × 36 mm) was measured by a universal tester (Instron-1195, Instron, Canton, MA, USA)
using a 30 mm span and a cross-head speed of 0.5 mm min−1. Vickers
Ta2O5 + 7C → 2TaC+5CO(g)
HfO2 + 3C → HfC + 2CO(g)
SiO2 + 3C → SiC + 2CO(g)
2SiO2 + SiC → 3SiO(g) + CO(g)
The Gibbs free energy (ΔG) of each reaction was calculated by
thermodynamic simulation software (HSC Chemistry 6.1), and they
were shown in Fig.1a. To complete the reactions above, adequate
amount of carbon black was added to react with oxide impurities. The
thermal decomposition of P. R. occurred before densification initiation,
during which all P. R. will be removed from the green body at 1000 °C.
0.9 wt% of pyrolytic carbon from P. R. removal was added into the total
carbon contents precisely for obtaining Ta0.8Hf0.2C-27 vol% SiC. Reactions (2)–(4) occurs since 1600 °C, while Reaction (5) starts at above
1900 °C as shown in Fig. 1a. Besides, since SiO2 impurity could not be
eliminated during the sintering of monolithic SiC without the presence
of carbon additive, Reactions (2)–(4) were considered to occur for removing the surface impurities.
Reaction (1) was taken out and investigated individually. The XRD
patterns of HfSi2-carbon black sintered at 900 °C (P. R. removal) and
1650 °C (reaction (1) initiation) were shown in Fig.1b. Peaks of mixed
HfSi2 and carbon black were consistent with the HfSi2 (PDF#38-1373)
at 900 °C. Well-defined HfC (PDF#39-1491) and β-SiC (PDF#29-1129,
3C-SiC) peaks were detected after the sintering at and above 1650 °C
together with the complete disappearance of HfSi2 peaks, which indicated that reaction (1) was completed. After the temperature reached
2200 °C, existence of α-SiC phases in Fig.1b (PDF#29-1127, PDF#492
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
B. Zhang et al.
Fig. 2. EBSD phase-map of Ta0.8Hf0.2C-SiC composite; EBSP of Ta0.8Hf0.2C, 4H-SiC and 6H-SiC particles in the composite.
Fig. 3(e), new α-SiC nuclei were assumed to grow inside the β-SiC
grain, resulting in β/α hybrid SiC grains at above 1650 °C. Strain at the
β/α interface might influence the orientation growth of final elongated
α-SiC grains during the temperature-increasing stage [19].
The anisotropic α-SiC grains were randomly oriented and uniformly
distributed in the Ta0.8Hf0.2C matrix. The relative density of
Ta0.8Hf0.2C-27 vol%SiC increased from 87.2% to 99.0% after soaking at
2200 °C, as shown in Table 1. Soaking at 2200 °C led to limited grain
growth of Ta0.8Hf0.2C matrix from 11.2 μm to 13.1 μm. Interestingly,
the aspect ratio of the elongated α-SiC grains increased remarkably
from 6.0 to 15.2 after 2 h’s dwelling at 2200 °C. Highly directional
growth (platelet-like) of α-SiC grains was generated. α-SiC grains interconnected with each other as interlocking structures in Fig. 3d. This
was speculated to improve the damage tolerance significantly [20]. The
thermal expansion coefficient mismatch between the Ta0.8Hf0.2C and
SiC [19] during sintering was presumably to be responsible: localized
stresses generated at the grain boundary could guide the directional
grain growth of the hexagonal α-SiC during the dwelling at 2200 °C.
HRTEM result of the sample sintered at 2200 °C for 2 h revealed the
Ta0.8Hf0.2C and α-SiC grain boundary was clean, without any evidence
of inter-diffusion between Ta0.8Hf0.2C and α-SiC in Fig. 3f. Additionally,
grain coarsening of the matrix phase (Ta0.8Hf0.2C) proceeded mainly
during the 2200 °C soaking period based our previous research [7]. As a
result, localized stresses at the grain boundary of Ta0.8Hf0.2C and α-SiC
may be further accumulated and therefore lead to evolution of the highaspect-ratio grains [7,21] and detailed study is currently under way.
The mechanical properties of Ta0.8Hf0.2C-based composites sintered
at 2200 °C for 2 h were shown in Table 2. The Vickers hardness of
Ta0.8Hf0.2C-27 vol%SiC composite was 16.0 ± 0.1 GPa, inheriting well
from the monolithic Ta0.8Hf0.2C ceramic. The flexural strength of the
composite increased to 443 ± 22 MPa, 23.5% higher than the dense
Ta0.8Hf0.2C ceramic.
Toughness improvement in the composite was good. Toughness
values of up to 5.4 ± 1.2 MPa m1/2 was measured, 45.9% higher than
the SiC free sample, 3.7 ± 0.1 MPa m1/2. Crack path was marked by
the Vickers indentation technique as shown in Fig. 4. Platelet-like α-SiC
grains obviously deflected the crack propagation pathway into more
1428, 4 H/6H-SiC) implied a β-α SiC phase transformation.
(Ta,Hf)C solid solution here was formed via mutual diffusion between commercial cubic TaC and in-situ generated cubic HfC phase
(PDF#39-1491 by XRD verification). Limited carbon vacancy [15]
could stabilize in the cubic carbide lattices under carbon-deficient
condition, however, the carbon-rich environment during sintering
(both carbon black and phenolic resin were included in the raw recipe)
was presumed to contribute to the removal of carbon vacancy in our
(Ta,Hf)C solid solution lattice.
The lattice parameter of 2200 °C sintered (Ta,Hf)C was calculated
from the XRD to be 4.4858 ± 0.0003 Å, which followed the Vegard’s
Law (4.4899 Å) well. A Ta0.8Hf0.2C-based composite was formed with
the in-situ introduction of α-SiC at 2200 °C. (Ta,Hf)C solid solution
formation was investigated thoroughly in our previous work [7]. Ta
atoms can ‘dissolve’ more efficiently into the HfC lattice on account of
higher diffusion activation energy of HfC [5]. The main (Ta,Hf)C peak
shifted from (34.63°, 1650 °C) to (34.70°, 2200 °C) in Fig. 1c.
EBSD phase map and EBSP results of polished surface of Ta0.8Hf0.2C27 vol% SiC sintered at 2200 °C were exhibited in Fig. 2. EBSP results
showed the presences of Ta0.8Hf0.2C, (Cubic, 225 Space Group, m-3 m
Laue group) and α-SiC (4 H/6H-SiC, Hexagonal, 186 Space Group, 6/
mmm Laue group). The distribution of α-SiC within the Ta0.8Hf0.2C
matrix was homogeneous and no β-SiC was observed according to the
EBSD map. This confirmed that the phase transformation of β→α SiC
was completed after sintering at 2200 °C. By applying the internal
standard method for quantitative analysis of XRD at 2200 °C [16], the
composition of Ta0.8Hf0.2C and α-SiC phases could be quantified: the
volume fraction of α-SiC was about 27.04 vol%. It was consistent with
the calculated value based on Reaction 1.
The microstructural evolution of Ta0.8Hf0.2C-27 vol%SiC upon
densification was revealed in Fig. 3(a–d). The in-situ β-SiC distributed
in the matrix based on the XRD pattern at 1650 °C. The morphology of
SiC transformed from equiaxed (β-SiC, 1650 °C) to elongated (α-SiC,
2200 °C) grains. This microstructure evolution can be ascribed to the
phase transformation and growth behavior [17]. It was reported that
the β-α phase transformation of SiC could lead to the in-situ growth of
elongated α-SiC grains [18]. As the schematic diagram showed in
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
B. Zhang et al.
Fig. 3. Backscattered electron images (BSE) of samples sintered at (a)1650 °C and 2200 °C for (b)0 h, (c)1 h and (d) 2 h soaking; (e) schematic diagram showing β-α
phase transformation of SiC in Ta0.8Hf0.2C matrix; (f) HRTEM of grain boundary between Ta0.8Hf0.2C and α-SiC grains.
smaller crack-opening displacement (U1) than the previous one (U0)
[22] in Fig. 4: adjacent to a crack deflection, crack-tip of the branching
was observed. When the crack-tip encountered the matrix grains during
its propagation course, stress intensity was reduced heavily. Such energy dissipation demonstrated consumed fracture energy and enhanced
toughness. Highly-homogenous microstructure and highly-elongated
(or platelet-like) α-SiC grains contributed to the higher toughness in the
composites: typically, the formation of a tough interfacial interlocking
microstructure between α-SiC grains and the matrix led to the improvement of toughness [23]. Besides, residual thermal stress field,
originated from the thermal expansion mismatch [24] between SiC and
Ta0.8Hf0.2C, might contribute to the crack bridging and deflection behaviors.
Toughening of our composite was caused by crack-shielding effect,
indicating toughness increased as cracks grew, or so-called R-curve
behavior [25]. This toughening does little contribute to strength increase. However, a balance between the toughness and strength improvements was achieved by in-situ introduction of α-SiC into the
Ta0.8Hf0.2C matrix in our work. The strategy of the strong Ta0.8Hf0.2CSiC bonding and interlocking structure were responsible for the simultaneous improvement of strength and toughness.
Table 1
Information of Ta0.8Hf0.2C-27 vol% SiC sintered at 2200 °C for different soaking
Time (h)
Length (μm)
Width (μm)
Table 2
Mechanical properties of Ta0.8Hf0.2C based composites sintered at 2200 °C for
2 h.
Fraction of SiC
KIC/MPa m1/2
27 vol%
15.9 ± 0.2
16.0 ± 0.1
3.7 ± 0.1
5.4 ± 1.2
359 ± 6
443 ± 22
tortuous and dissipated fracture energy, and therefore toughness was
enhanced. Crack branching generated at the strain debonded area of the
SiC grains and dissipated fracture energy. The width of propagating
cracks became narrower after they passed across SiC particles with
Fig. 4. BSE image of indentation crack propagation on polished surface in Ta0.8Hf0.2C-27 vol% SiC sample.
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
B. Zhang et al.
4. Conclusions
Ta0.8Hf0.2C-27 vol%SiC composite was pressurless reactive-sintered
with a relative density of 99.0% after soaking 2 h at 2200 °C using
commercially available TaC, HfSi2 and carbon black.
Well-documented cubic phase HfC (PDF#39-1491) and β-SiC
(PDF#29-1129, 3C-SiC) peaks were detected after reaction between
HfSi2 and carbon at above 1650 °C. Ta0.8Hf0.2C was formed via mutual
diffusion between raw TaC and in-situ HfC with a lattice parameter of
4.4858 ± 0.0003 Å after sintering at 2200 °C. In-situ β-α SiC phase
transformation occurred during sintering. Strain at the β/α interphase
was found to influence the oriented growth of final elongated α-SiC
The Vickers hardness of Ta0.8Hf0.2C-27 vol%SiC composite was
16.0 ± 0.1 GPa. The flexural strength of the composite increased to
443 ± 22 MPa, which was 23.5% higher than the dense Ta0.8Hf0.2C
ceramic. Toughness values of up to 5.4 ± 1.2 MPa m1/2 was measured,
which was 45.9% higher than its SiC free counterpart. Formation of a
tough interfacial interlocking microstructure between highly-elongated
α-SiC grains and the solid solution matrix led to the improvement of
Financial support from National Natural Science Foundation of
China (No. 51602325), (No. 51572276), Youth Innovation Promotion
Association (CAS), Science Foundation for Youth Scholar of State Key
Laboratory of High Performance Ceramics and Superfine
Microstructures, Shanghai Institute of Ceramics Chinese Academy of
Sciences (SKL201602), Scientific and Technological Innovation Project
of Shanghai Institute of Ceramics are gratefully acknowledged.
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