Journal Pre-proof
Effects of heat treatment on mechanical properties of 3D
Si3 N4f /BN/Si3 N4 composites by PIP
Jie Zhou, Fang Ye, Laifei Cheng, Mingxing Li, Xuefeng Cui, Zhiqiang
Fu, Litong Zhang, Nan Chai
PII:
S0955-2219(21)00513-6
DOI:
https://doi.org/10.1016/j.jeurceramsoc.2021.07.043
Reference:
JECS 14219
To appear in:
Journal of the European Ceramic Society
Received Date:
26 May 2021
Revised Date:
16 July 2021
Accepted Date:
22 July 2021
Please cite this article as: Zhou J, Ye F, Cheng L, Li M, Cui X, Fu Z, Zhang L, Chai N, Effects
of heat treatment on mechanical properties of 3D Si3 N4f /BN/Si3 N4 composites by PIP,
Journal of the European Ceramic Society (2021),
doi: https://doi.org/10.1016/j.jeurceramsoc.2021.07.043
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© 2020 Published by Elsevier.
Effects of heat treatment on mechanical properties of 3D Si3N4f/BN/Si3N4
composites by PIP
Jie Zhoua, Fang Yea1*, Laifei Chenga2*, Mingxing Lia, Xuefeng Cuia, Zhiqiang Fua,
Litong Zhanga, Nan Chaib
a
Science and Technology on Thermostructural Composite Materials Laboratory,
Northwestern Polytechnical University, Xi’an 710072, China
b
Institut für Materialwissenschaft, Technische Universität Darmstadt, Otto-Berndt-
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of
Straße 3, D-64287 Darmstadt, Germany
1
Corresponding author: yefang511@nwpu.edu.cn (F. Ye).

Corresponding author: chenglf@nwpu.edu.cn (L. Cheng).
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ABSTRACT
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The fabrication of three-dimensional silicon nitride (Si3N4) fiber-reinforced
silicon nitride matrix (3D Si3N4f/BN/Si3N4) composites with a boron nitride (BN)
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interphase through precursor infiltration and pyrolysis (PIP) process was reported.
Heat treatment at 1000-1200 °C was used to analyze the thermal stability of the
Si3N4f/BN/Si3N4 composites. It was found after heat treatment the flexural strength
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and fracture toughness change with a pattern that decrease first and then increase,
which are 191±13 MPa and 5.8±0.5 MPa·m1/2 respectively for as-fabricated
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composites, and reach the minimum values of 138±6 MPa and 3.9±0.4 MPa·m1/2
respectively for composites annealed at 1100 °C. The influence mechanisms of the
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heat treatment on the Si3N4f/BN/Si3N4 composites include: (Ⅰ) matrix shrinkage by
further ceramization that causes defects such as pores and cracks in composites, and
(Ⅱ) prestress relaxation, thermal residual stress (TRS) redistribution and a better
wetting at the fiber/matrix (F/M) surface that increase the interfacial bonding strength
(IBS). Thus, heat treatment affects the mechanical properties of composites by
changing the properties of the matrix and IBS, where the load transfer efficiency onto
the fibers is fluctuating by the microstructural evolution of matrix and gradually
increasing IBS.
Keywords: Si3N4f/BN/Si3N4 composites; Microstructural evolution; Mechanical
properties; Heat treatment; Precursor infiltration and pyrolysis.
1. Introduction
With the development of weapons and equipment, the precision guidance field of
aircrafts has put forward more stringent requirements for multi-functional integrated
radomes materials [1-3]. Continuous silicon nitride (Si3N4) fiber-reinforced silicon
nitride matrix (Si3N4f/BN/Si3N4) composites with a boron nitride (BN) protective
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interphase not only possess a series of excellent properties such as low density, high
temperature resistance, oxidation resistance, and moderate dielectric properties similar
to Si3N4 ceramics, but also overcome the brittleness of ceramic, which are candidate
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of radomes materials with excellent comprehensive performance [4-7]. Currently,
there are few public research reports about the Si3N4f/BN/Si3N4 composites. In our
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previous work [8, 9], Si3N4f/BN/Si3N4 composites were fabricated by chemical vapor
infiltration (CVI) and precursor infiltration and pyrolysis (PIP) processes. According
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to our research results, the PIP process can not only shorten the preparation time, but
also control the interface bonding state and greatly improve the mechanical properties
of the Si3N4f/BN/Si3N4 composites.
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Considering their potential high temperature applications, the thermal stability is
a key issue for Si3N4f/BN/Si3N4 composites prepared by PIP [10-12]. For any given
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polymer-derived ceramics (PDCs), which are obtained from the conversion of organic
polymer to inorganic compound, properties are dominated by their underlying
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microstructure tailored by controlling the processing parameters, bringing an
amorphous, nanocrystalline or crystalline structure [13, 14]. When the application
temperature of composites is higher than their preparation temperature (i.e. the
pyrolysis temperature of PDCs matrix), the microstructure and composition of PDCs
matrix will probably change owing to the decomposition and crystallization of
amorphous phase, and the grain growth of nanocrystalline. Thereout, the density and
porosity of composites will change subsequently. More importantly, the bonding state
between fibers and matrix, mainly determining the mechanical behavior of
composites, will change due to the variation of stress state (including stress level and
stress distribution, etc) in composites, which is caused by the matrix performance
(such as the coefficient of thermal expansion (CTE) and elastic modulus) evolution
and different wettability at the fiber/matrix (F/M) surface during the further
ceramization process.
To comprehensively investigate the thermal stability of composites, a heat
treatment process, for which the heat treatment temperature (HTT) is usually higher
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than the preparation temperature, can be performed on the composites [15, 16], and
the microstructural evolution of composites at heat treatment temperature is followed
and its influence mechanism on the mechanical properties of composites is presented.
Based on this, in this paper, the PIP Si3N4f/BN/Si3N4 composites, as well as the
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polymer-derived Si3N4 bulk ceramics, used to supplement the study on the change in
the microstructure and properties of the matrix, were respectively prepared, and the
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effects of HTT (1000-1200 °C) on the microstructure and mechanical behavior of
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Si3N4f/BN/Si3N4 composites were investigated.
2. Experimental Procedure
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2.1. Preparation of Si3N4f/BN/Si3N4 composites
Continuous Si3N4 fibers have a density of 2.35 g∙cm−3, an average filament
diameter of 12.2 μm and an average filament tensile strength of 1.3±0.2 GPa, which
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were provided by Xiamen University in China. The Si3N4 fibers were braided into the
fabrics with a three-dimensional (3D) four-directional structure and a fiber volume
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fraction of 38% by Shaanxi textile research institute. BN interphase with a thickness
of about 500 nm was firstly deposited on the surface of Si3N4 fibers via CVI process
from the precursors of boron trichloride (BCl3) and ammonia (NH3) at 650 °C before
fabricating Si3N4 matrix. The researches indicate that the Si3N4 fibers [17] and the BN
interphase [15] used and prepared in this work are stable in composition and
microstructure below 1300 °C, on which the effect of heat treatment can be ignored.
Polysilazane (PSZ, Institute of Chemistry Chinese Academy of Sciences, China) was
used as the polymer precursor of Si3N4 matrix, and the fabrics were infiltrated with
PSZ under vacuum for 1 h, then cross-linked at 300 °C for 2 h and pyrolyzed at
900 °C for 2 h in nitrogen atmosphere, respectively. After 12 cycles of PIP, the
relatively dense Si3N4f/BN/Si3N4 composites were obtained. Polymer-derived Si3N4
bulk ceramics were also fabricated under the same polymer-derived condition as PIP
Si3N4 matrix. The Si3N4f/BN/Si3N4 composites and Si3N4 bulk ceramics were heat
treated at the temperature of 1000, 1100, and 1200 °C in vacuum for 2 h, respectively,
and the unannealed samples were named as RT (room temperature) for comparative
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analysis.
2.2. Microstructure characterizations
The morphology and microstructure of composites were observed by focused
ion/electron double beam electron microscope (FIB, Helios G4 CX, FEI, America).
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The phase composition and degree of crystallization of composites were analyzed by
X-ray diffractometer (XRD, X’ Pert Pro, Philips, Netherlands) using Cu Kα (λ = 1.54
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Å) radiation. Double Cs corrector transmission electron microscope (TEM, Themis Z,
FEI, America) was employed to precisely characterize the microstructure of Si3N4
2.3. Performance tests
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matrix with different HTTs.
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The density and open porosity of Si3N4f/BN/Si3N4 composites were measured by
Archimedes-method. Three-point bending test was utilized to measure the flexural
strength and elastic modulus of composites with sample size of 40×5×3 mm3 and span
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of 30 mm at a loading rate of 0.5 mm/min by a universal testing machine (SANS
CMT 4304, Sans Materials Testing Co. Ltd., Shenzhen, China). Fracture toughness
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was measured by using single edge notched beam (SENB) method with a sample
dimension of 40×3×5 mm3 over span of 30 mm at a loading rate of 0.05 mm/min,
where the notch was machined by a diamond-coated miniature blade, and the notch
width and length are 0.2 mm and 2.5 mm, respectively.
The elastic modulus and hardness of fibers and matrix of composites were
measured by nano-indentation test using nano-indentation apparatus (TI980, Hysitron,
America) with a triangular diamond micro-indenter with a radius of curvature of 100
nm and a maximum load of 5 mN.
The engineering CTE and physical CTE of Si3N4f/BN/Si3N4 composites and
Si3N4 bulk ceramics with and without heat treatment were measured by a thermal
dilatometer (DIL 402C, Netzsch, Germany) under argon atmosphere and at a heating
rate of 5 °C/min, where the specimens were machined into dimensions of 20×3×3
mm3.
3. Results and discussion
before and after heat treatment
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3.1. Mechanical properties and fracture behavior of Si3N4f/BN/Si3N4 composites
The fundamental properties of Si3N4f/BN/Si3N4 composites are listed in Table 1
as a function of HTT, where the elastic modulus and matrix cracking stress were also
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obtained from flexural stress-displacement curves [18].
Table 1 Fundamental properties of Si3N4f/BN/Si3N4 composites
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with different HTTs.
Bulk density Open porosity Flexural strength Fracture toughness Elastic modulus Matrix cracking
(g/cm3)
(%)
(MPa)
(MPa·m1/2)
(GPa)
stress (MPa)
RT
2.00
8.4
191±13
5.8±0.5
46±2
85.8±4.7
1000 2.00
8.4
182±5
5.4±1.1
76±4
112.1±3.6
1100 1.97
10.7
138±6
3.9±0.4
61±3
108.6±3.9
1200 1.94
11.1
155±5
4.3±0.1
68±2
102.7±3.2
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HTT
(°C)
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It is obvious that as-fabricated composites possess the highest density and lowest
open porosity shown in Table 1, for which the multiple infiltration-pyrolysis cycles at
900 °C result in a larger bulk density of matrix through filling residual pores caused
by the second-to-last pyrolysis. The density and open porosity of composites with
heat treatment at 1000 °C are almost unchanged, but when the HTT rises further, the
density and open porosity of the composites drops and increases obviously,
respectively. The results indicate that the thermal stability of PIP Si3N4f/BN/Si3N4
composites is insufficient above its preparation temperature, where porosity will
increase caused by microstructural evolution of composites [19]. Density and porosity
have significant effect on multiple properties of ceramics, including flexural strength,
modulus, and fracture toughness, which further affect mechanical properties of
composites. Compared with the as-fabricated composites, both the flexural strength
and fracture toughness decrease of annealed composites, which is mainly due to the
increase of defects such as pores and cracks caused by the microstructural evolution
of composites. However, the increase of elastic modulus and matrix cracking stress
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after heat treatment indicate that heat treatment will have a positive effect on the
elastic modulus and matrix cracking stress of the composites.
Fig. 1 shows the typical stress-displacement curves of Si3N4f/BN/Si3N4
composites with different HTTs, indicating that all the mechanical behavior exhibits
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pseudo-ductile fracture mode after bending test and fibers possess effective
toughening effect. Compared with the as-fabricated composites, for the annealed
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composites, the slope change of curves is smaller after the matrix cracking,
corresponding to the inflection point from the linear segment to the nonlinear segment
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of the curve, and the curves drop more sharply after the maximum stress, which
indicates that the interface debonding effect is weakened, causing the ability of
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interface to transfer load and the co-bearing capacity of fibers improve consistent with
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the analysis of stress state, i.e. higher interfacial bonding strength (IBS).
Fig. 1. Typical stress-displacement curves of Si3N4f/BN/Si3N4 composites with
different HTTs after bending test.
Fig. 2 shows the SEM images of fracture surfaces of the Si3N4f/BN/Si3N4
composites with different HTTs. For as-fabricated composites, a large number of
monofilament fibers are pulled out on the fractured section and the over pullout of the
fibers (>300 μm) indicates the low IBS. Thus, the stress cannot be transferred to the
fibers efficiently, and the fibers are pulled out from the matrix at low applied stress
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levels, where the bearing capacity of fibers cannot be fully exploited. After heat
treatment, the pull-out length of fibers is significantly shorter (<100 μm), meaning
that IBS has been increased to a certain extent. It is worth noting that the fracture
morphologies show that the fibers are mainly pulled out in clusters, where the length
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of the fractured fibers is relatively even and plenty of rough matrix without debonding
(a)
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(b)
(d)
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(c)
(e)
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exists between the filaments after heat treatment at 1200 °C.
(f)
(h)
(g)
Fig. 2. SEM images of fractured section of the Si3N4f/BN/Si3N4 composites with
different HTTs after bending test:
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(a, b) RT, (c, d) 1000 °C, (e, f) 1100 °C and (g, h) 1200 °C.
3.2. Microstructural evolution characteristics of Si3N4f/BN/Si3N4 composites
For the continuous fibers reinforced ceramic matrix composites (CFCMCs),
ceramic matrix, one of the basic components, can not only form the composites and
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protect the internal fibers, but more importantly, it can transfer the applied load to the
fibers to achieve the strengthening and toughening effect of fibers, which is closely
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related to the mechanical properties of the composites. As mentioned in the
experimental part, Si3N4 fibers and BN interphase are stable in composition
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microstructure under the studied HTTs in this work. Thus, the mechanical behavior at
service temperature of PIP Si3N4f/BN/Si3N4 composites analyzed in section 3.1
mainly depends on the thermal stability of polymer-derived Si3N4 matrix and its
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microstructural evolution should be paid attention during the heat treatment process.
Fig. 3 shows the XRD patterns of Si3N4f/BN/Si3N4 composites with different
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HTTs, where just only several broad and weak peaks exist. It is hard to express the
changes in the phase composition and crystallinity of the composites before and after
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heat treatment through XRD patterns. However, there is a phenomenon showing that
the composites have a tendency to crystallize after heat treatment at 1200 °C, where a
small sharp peak appears at 2θ = 22.9°, corresponding to the (1 1 0) plane of α-Si3N4
[20].
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Fig. 3. XRD patterns of Si3N4f/BN/Si3N4 composites with different HTTs.
The microstructure of Si3N4 matrix with different HTTs was observed in detail
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by high-resolution TEM given in Fig. 4. The TEM results indicate that the as-
fabricated Si3N4 matrix is amorphous and possesses a uniform structure with fine
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particles, which can be maintained after heat treatment at 1000 °C. It is clear that the
microstructure of the matrix becomes more uniform with finer particles with the
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increase of HTT. Meanwhile, after heat treatment at 1100 °C, Si3N4 matrix slightly
crystallizes, which is not obvious enough and difficult to be found. The interplanar
spacing of α-Si3N4 is 0.34. When the HTT increases from 1100 °C to 1200 °C, the
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size of α-Si3N4 nano-crystallites increases from 1-2 nm to 3-5 nm and the number also
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(a)
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increases accordingly.
(b)
(c)
(d)
(a) RT, (b) 1000 °C, (c) 1100 °C and (d) 1200 °C.
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Fig. 4. High-resolution TEM images of Si3N4 matrix with different HTTs:
The above TEM analysis shows that the microstructure of Si3N4 matrix has
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changed after heat treatment, which will affect the density and porosity of composites.
Fig. 5 shows the micro-area feature of Si3N4 matrix and the F/M zone in
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Si3N4f/BN/Si3N4 composites with different HTTs. It is concluded that due to
shrinkage and gas release during pyrolysis, PDCs are inherently porous but the
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densification of composites can be achieved by repeated cycles of PIP (Figs. 5(a and
b)). However, heat treatment promotes inorganic and crystalline conversion at the
expense of porosity, and more pores with larger size are generated in the matrix with
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the increase of HTT (Figs. 5(c~h)). Based on the conversion from polymer to ceramic
[21, 22], the generation of pores mainly lies in the microstructure rearrangement of
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the matrix by intensifying the degree of inorganic matrix and the further release of gas
at high temperature, leading to matrix shrinkage and more pores remaining in the
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matrix as listed in Table 1 [23]. Combined with the analysis about Fig. 4, the
microstructure rearrangement and the more uniform microstructure of matrix may
cause a better wetting at the F/M surface particularly at 1200 °C, which leads to a
stronger F/M bonding.
(b)
(c)
(d)
(e)
(f)
(g)
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(a)
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(h)
Fig. 5. SEM images of matrix and F/M zone of the Si3N4f/BN/Si3N4 composites with
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different HTTs: (a, b) RT, (c, d) 1000 °C, (e, f) 1100 °C and (g, h) 1200 °C.
The nano-indentation test was carried out to obtain the in-situ properties of fibers
and matrix in the composites before and after heat treatment. Fig. 6 plots the elastic
modulus and hardness of the Si3N4 fibers and Si3N4 matrix in the Si3N4f/BN/Si3N4
composites with different HTTs measured by nano-indentation test. The elastic
modulus and hardness of the fibers with different HTTs are increased from 151.2 GPa
and 12.5 GPa to 166.3 GPa and 13.7 GPa, respectively. Combined with our previous
research [8], it can be obtained that the structure of Si3N4 fibers is almost stable at this
temperature interval of 1000-1200 °C and polymer-derived Si3N4 fibers will become
denser with heat treatment, which increases its elastic modulus and hardness. In
comparison, the increasing HTT from RT to 1200 °C results in a change of first
decreasing and then increasing of the elastic modulus and hardness of the Si3N4
matrix, which is due to the poor thermal stability of PDCs. Affected by the content
and strength of Si-N bond, as the degree of ceramization of the matrix intensifies, its
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modulus and hardness gradually increase [24, 25]. However, when PDCs prepared at
low temperature are heated to higher temperature, their average elastic modulus and
hardness will decrease due to the increase in porosity [26, 27]. As Figs. (4 and 5)
show, the matrix shrinkage that accompanies the increase in solid density at higher
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temperatures and bring increase in porosity. With the increase of HTT, the increase of
porosity and solid density are simultaneous, and affect the mechanical properties of
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the matrix together, where the dominant factor is the increase in porosity before
1100 °C while the one at 1200 °C is the increase of the degree of ceramization. The
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changes in both elastic modulus and hardness will inevitably affect the mechanical
properties of the composites, especially the increase of elastic modulus of composites
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after heat treatment listed in Table 1.
(b)
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(a)
Fig. 6. (a) Elastic modulus and (b) hardness of the Si3N4 fibers and Si3N4 matrix in the
Si3N4f/BN/Si3N4 composites with different HTTs.
3.3. Stress state in Si3N4f/BN/Si3N4 composites before and after heat treatment
In addition to the porosity and the properties of the fibers and matrix, the F/M
interface bonding characteristics is also a main factor affecting the mechanical
properties of CFCMCs. Generally, the physical and the chemical compatibility
between the fibers and matrix determines the F/M interfacial bonding strength
together [28], where the former is related to the fiber surface roughness and the
degree of thermal mismatch between the fibers and matrix, while the latter is related
to the chemical reaction between the fibers and matrix [29]. Based on our previous
research [9], the F/M interfacial bonding strength is relatively weak in the as-
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fabricated Si3N4f/BN/Si3N4 composites, because the lack of interface reaction between
the BN interphase and Si3N4 matrix and the action of tensile stress on BN interphase
caused by the prestress from matrix shrinkage in the PIP process. However, the
microstructural evolution of the matrix at high temperature will change both the
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prestress state and CTE of matrix, affecting the degree of thermal mismatch between
the fibers and matrix, where the variation of both prestress and thermal residual stress
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(TRS) will lead to the different F/M interface bonding state, and subsequently
different mechanical behavior of composites.
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The engineering CTE and physical CTE of the as-fabricated Si3N4f/BN/Si3N4
composites and PIP Si3N4 matrix are shown in Fig. 7. It can be seen from the Fig. 7
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that the CTE increases uniformly before the preparation temperature (900 °C). For the
Si3N4 matrix, when the temperature exceeds 900 °C, the CTE begins to decrease and
the physical CTE drops to 0 at around 970 °C. Subsequently, the matrix no longer
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expands and begins to shrink, where the effect of atom thermal motion is weaker than
that of matrix shrinkage. However, the physical CTE begins to increase at around
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1143 °C indicating that the effect of shrinkage begins to decrease. All changes of the
matrix CTE are related to the further ceramization and atom thermal motion of matrix.
Interestingly, although the matrix is shrinking, the engineering CTE and physical
CTE of the composites gradually increase and exceed the maximum CTE of the fibers
(1.5×10-6 K-1 [30]) and matrix above 1000 °C (Fig. 7), which is related to the release
of prestress. In order to clarify this phenomenon, the stress state and thermal
expansion behavior of the composites in different temperature sections are illustrated
in Fig. 8. During the infiltration process, the precursor evenly wraps the fibers due to
surface tension, and the precursor gradually becomes inorganic with the temperature
increases. Due to the restraint of the fibers, the matrix is not only porous, but also
cracks perpendicular to the axial direction of the fibers, which generates prestress, i.e.
axial compressive stress (τ) and radial tensile stress (F), on the fibers (Fig. 8(a)). As
the temperature increases, the matrix shrinks further, where the radial stress causes the
debonding energy of the F/M interface decrease and the axial stress makes elastic
strain energy of the fibers increase. When the temperature reaches about 1000 °C, the
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elastic strain energy of the fiber is greater than the debonding energy of the F/M
interface. At this time, the fibers quickly expand free of the bondage of the matrix and
fiber/matrix (F/M) zone returns to a state of energy balance with small compressive
stress (Fig. 8(b)). With the release of fibers’ elastic strain energy, the matrix shrinkage
(b)
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(a)
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Fig. 7. (a) Engineering CTE and (b) physical CTE of the PIP Si3N4f/BN/Si3N4
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composites and PIP Si3N4 matrix without heat treatment.
(b)
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Fig. 8. The schematic diagram of characteristics of the F/M zone in different
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temperature sections: (a) around 900 °C, (b) around 1100 °C.
According to the analysis of the thermal expansion behavior of the composites,
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the as-fabricated composites have already had a prestress distribution at 900 °C due to
the conversion of the precursor to ceramic, which can be released at higher
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temperature with the microstructure rearrangement of the matrix. Under the co-effect
of heat and stress during the heat treatment process, the local creep of the polymerderived Si3N4 matrix [31], where the denser microstructure of the matrix (Fig. 4) and
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pores (Fig. 5) are obtained, relaxes the prestress and causes a better wetting at the F/M
surface, which is helpful to increase the IBS.
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In addition, microstructural evolution also affects TRS by the changes in CTE.
The calculation formula for TRS of fibers and matrix is as follows [32]:
E
σr =σmr =σfr =(αm -αf )ΔT Emm
Ef
σma =-(αm -αf )ΔT
(2)
+1
λEm Vf
λVf (
Ef
-1)+1
Em
(1)
σfa =(αm -αf )ΔT
λEf Vf
λVf (
Ef
-1)+1
Em
(3)
where σ and α stand for the TRS and CTE, while the subscript r, a, m and f represent
the radial direction, axial direction, matrix and fibers respectively; here, αm is the
average value in ΔT; ΔT equals to T-T0, where T0 is preparation temperature or HTT
of composites, and T is the testing temperature; λ is the preform woven coefficient
and here λ = 0.766 for the 3D preform in this work. The CTE of Si3N4 matrix with
heat treatment were also measured (supporting information), which were used to
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calculate TRS. The calculated TRS at the F/M interface of the Si3N4f/BNSi3N4
composites are listed in Table 2, where the negative value indicates a pressure and
positive value means a tension.
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Table 2 The calculated TRS at the F/M interface in the Si3N4f/BN/Si3N4 composites
with different HTTs.
αm (×10-6 K-1) αf (×10-6 K-1) σr (MPa)
σma (MPa)
σfa (MPa)
RT
2.88
-75.6
33.5
-84.7
1000
2.51
-60.3
24.3
-69.4
-64.1
24.1
-75.2
-98.7
43.3
-110.9
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HTT (°C)
1.5 [30, 33]
2.51
1200
2.73
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1100
Axial tress affects matrix cracking, and radial tress is related to interface
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bonding, both of which make an important influence on mechanical properties of
composites. As listed in Table 2, with the increase of HTT, all the calculated TRS
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show a pattern that decrease first and then increase. For as-fabricated composites,
corresponding stress value should be -75.6+σ'r , 33.5+σ'ma and -84.7+σ'fa respectively,
where σ'r , σ'ma and σ'fa are the prestress, i.e. radial tensile stress, axial tensile stress
and axial compressive stress pre-generated by the matrix shrinkage. It is obtained that
with the increase of radial compressive stress listed in Table 2, the radial fretting also
leads to a stronger F/M bonding.
Based on the above analysis about stress state, the increase of IBS, caused by
prestress relaxation, TRS redistribution, and a better wetting at the F/M surface, and
the properties of matrix discussed in Section 3.2 co-explain the fracture behavior of
the Si3N4f/BNSi3N4 composites shown in Fig. 2
3.4. Influence mechanism of heat treatment on mechanical properties
Si3N4f/BN/Si3N4 composites
The factors that can affect the mechanical properties of the composites are
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analyzed above. Due to the microstructural evolution of matrix, porous matrix and the
higher IBS, co-determined by prestress relaxation, TRS redistribution, and a better
wetting at the F/M surface, will cause different mechanical properties of
Si3N4f/BN/Si3N4 composites and then the influence mechanism of heat treatment on
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mechanical properties of the composites are discussed.
As listed in Table 1, both the flexural strength and fracture toughness decrease of
re
annealed composites, which is decidedly influenced by the defects such as pores and
cracks by microstructural evolution of matrix. However, after heat treated at 1200 °C,
lP
both the flexural strength and fracture toughness have a certain increase, which is due
to the effective load-bearing capacity of fibers by high-modulus matrix and increasing
na
IBS.
Through the above analysis and the fracture morphologies shown in Fig. 2, IBS
ur
(τi) can be roughly obtained according to the classical shear-lag model predictions:
σ𝑑
lc = 2τf
i
(4)
Jo
where σf and d are the in-situ strength and diameter of the fibers, lc is the critical fiber
pull-out length. Through our previous research, it can be considered that the in-situ
strength of the fibers is almost unchanged before and after heat treatment [8].
Therefore, according to the predicted result based on formula (4), heat treatment
significantly increases IBS, at least 3 times that of as-fabricated composites (<27
MPa). However, that may not be a bad thing. A moderate bonding strength between
fibers and matrix is a required condition for the bearing capacity of fibers in
composites, which is often underestimated although it is essential. A relatively strong
interface bonding can effectively realize the load transfer function during the loading
process. Since the effective load transfer is maintained up to failure, the fracture curve
no longer exhibits the plateau-like feature as the stress-displacement curve of asfabricated composites, where failure occurs at a higher stress level with limited fiber
pull-out length [34]. Based on the above analysis, it can be considered that the IBS is
moderate after heat treatment, and load transfer at F/M interface is better. However,
according to the data in Table 1, both the flexural strength and fracture toughness of
of defects such as pores and cracks during heat treatment.
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of
composites decrease after heat treatment, which can be attributable to the generation
It is worth noting that compared to the as-fabricated composites, both the elastic
modulus and matrix cracking stress increase of annealed composites (Table 1), which
modulus of the composites can be expressed as,
-p
is related to the properties of fibers and matrix and IBS. The theoretical elastic
re
Ec =Ef λVf +Em Vm
(5)
lP
where Ec are the elastic modulus of composites. Em, Ef, Vm, and Vf refer to the elastic
modulus and the volume fraction of the matrix and fibers, respectively. λ is the
preform woven coefficient and here λ = 0.766 for the 3D preform in this work. In this
na
case, the predicted minimum elastic modulus of the completely-cracked as-fabricated
Si3N4f/BN/Si3N4 composites should be equal to 44 GPa when Ef = 151.2 GPa, Em ≈ 0,
ur
Vf = 0.38, which is close to the value at RT shown in Table 1. Heat treatment further
affects the elastic modulus of Si3N4f/BN/Si3N4 composites by affecting the IBS and
Jo
the matrix microstructure [35], which is not considered by the formula (5). For asfabricated composites, its interface bonding is relatively weak and load transfer ability
of fibers is inefficient, so its elastic modulus is lowest [36]. The remarkably increased
slope indicates that the elastic modulus of the composites has increased after heat
treatment shown in Fig. 1, which is attributed to the increase of IBS by the prestress
relaxation, TRS redistribution, and a better wetting at the fiber/matrix surface.
However,
the increase in the porosity of the matrix leads to the decrease of the elastic modulus
of composites. The elastic modulus decreases due to porous matrix after heat
treatment at 1100 °C, which increases by increased matrix modulus and higher IBS to
offset the effect of porosity after heat treatment at 1200 °C.
Matrix cracking stress were also obtained from the end of linear elastic
deformation process of Si3N4f/BN/Si3N4 composites, co-determined by factors such as
the fiber volume fraction, fiber architecture, IBS, TRS, density, matrix strength, etc.
As-fabricated composites possess the lowest matrix cracking stress attributed to the
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of
extremely weak IBS by prestress, and the influence of prestress on interface bonding
is greatly reduced by microstructure rearrangement after heat treatment. Overall
porosity is well-known to reduce the mechanical properties of ceramics [37, 38], and
the increasing axial tensile stress of matrix (Table 2) also reduces the matrix cracking
-p
stress. In this case, the fracture behavior after heat treatment at 1200 °C shown in
Figs. 2(g and h) can be explained. This can be understood as the IBS increases caused
re
by prestress relaxation, TRS redistribution, and a better wetting at the F/M surface,
while the fracture energy of the porous matrix decreases, resulting in a higher crack
lP
propagation energy at the F/M interface than in the porous matrix. Thus, compared
with the desired F/M region, the cracks can simultaneously propagate and deflect in
na
the porous matrix near the interface.
After heat treatment, the Si3N4f/BN/Si3N4 composites with moderate IBS possess
lower mechanical properties than as-fabricated composites, which is attributed to the
ur
composites damage by the formation of pores [39, 40]. With the increase of HTT, it is
clear that the flexural strength and fracture toughness of composites show a pattern
Jo
that decrease first and then increase. Although the mechanical properties do not
decrease much in the case of heat treatment at 1000 °C, the microstructure (Fig. 5)
and elastic modulus (Fig. 6) of the matrix still indicate that composites have damaged.
Both the lowest value of flexural strength and fracture toughness appears at 1100 °C.
With heat treatment at 1100 °C, the matrix becomes porous via further ceramization,
and its average modulus decreases, where its ability to transfer load is greatly
compromised and the reinforcing effect of the fibers cannot be effectively exerted.
However, when the HTT raises to 1200 °C, both the average modulus of the matrix
and IBS increase, which together determine the improvement of the mechanical
properties of the composites. The improvement of fracture toughness mainly depends
on the strengthening and toughening effect of fibers, although porous high-modulus
matrix also contributes to it [41].
The above analysis indicates that the heat treatment has a vital influence on the
mechanical properties of the Si3N4f/BN/Si3N4 composites without other postprocessing. On the one hand, due to the matrix shrinkage, the density decreases and
ro
of
the defects such as pores and cracks increase, thereby decreasing the flexural strength
and fracture toughness of the composites decisively. On the other hand, the increase
in the IBS and the variation on the in-situ properties of fibers and matrix via heat
treatment are beneficial to improve the mechanical properties of composites. There
-p
are the same findings existing in silicon carbide fiber-reinforced silicon carbide
matrix (SiCf/PyC/SiC) composites with a pyrolytic carbon (PyC) interphase fabricated
re
by PIP process [42, 43], where the composites possess better mechanical properties
fabricated at higher temperature or heat treated at appropriate high temperature than
lP
as-fabricated composites. Combining the above analysis, it provides a research idea
for preparing high-performance composites by PIP process, which is by adjusting the
na
suitable PIP temperature and post-processing to obtain a moderate IBS and highly
dense composites. This deserves further investigations for Si3N4f/BN/Si3N4
ur
composites.
Jo
4. Conclusions
3D Si3N4f/BN/Si3N4 composites were fabricated via PIP process, and then heat
treated at various temperatures (1000-1200 °C) to study the thermal stability. The
effects of heat treatment on the matrix microstructure, stress state at F/M zone, IBS
and mechanical properties of Si3N4f/BN/Si3N4 composites influenced by the former
three factors were carefully investigated. The main conclusions are as follows:
(1) Heat treatment has a vital influence on the mechanical properties of
Si3N4f/BN/Si3N4 composites. As the increase of HTT, the flexural strength and
fracture toughness of the Si3N4f/BN/Si3N4 composites show a pattern of
decreasing first and then increasing, while the elastic modulus and matrix
cracking stress of composites show a pattern of increasing first and then
decreasing, which is determined by the porous matrix, in-situ properties of fibers
and matrix, and IBS together.
(2) As the increase of HTT, the microstructure of the Si3N4 matrix becomes more
uniform with finer particles, and the matrix gradually precipitates nano-
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of
crystallites above 1100 °C via further ceramization at the expense of porosity,
which increases the defects such as pores and cracks in the composites and
decisively decreases the flexural strength and fracture toughness compared with
the as-fabricated composites.
-p
(3) As the increase of HTT, the elastic modulus of fibers gradually increases and that
of matrix shows a tendency of decreasing first and then increasing affected by
re
the increase of porosity and solid density of matrix together.
(4) Under the co-effect of heat and stress during the heat treatment process, the
lP
prestress relaxation, TRS redistribution, and a better wetting at the fiber/matrix
surface by the matrix shrinkage co-increase IBS and bring an effective load
na
transfer, which dominates the increase of elastic modulus and matrix cracking
stress of annealed composites compared with as-fabricated composites.
(5) The IBS and properties of fibers and matrix determine the performance of
ur
Si3N4f/BN/Si3N4 composites by PIP. Thus, the suitable PIP temperature and heat
treatment to obtain a moderate IBS and highly dense composites are deserved
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further investigations to obtain high-performance Si3N4f/BN/Si3N4 composites.
Declaration of Interest Statement
The authors have no affiliation with any organization with a direct or indirect
financial interest in the subject matter discussed in the manuscript.
5. Acknowledgement
This work was supported by the National Natural Science Foundation of China
(Grant No. 51632007, 52072304), the 111 Project of China (B08040), and National
Science and Technology Major Project (Grant: 2017-VⅠ-0007-0077). We would like
to thank the Analytical & Testing Center of Northwestern Polytechnical University
for the kind assistance with electron microscopic characterization in this work.
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