2.12 Properties and Characteristics of SiC and SiC/SiC
Composites
J. Lamon
CNRS/National Institute of Applied Science, Villeurbanne, France

ß 2012 Elsevier Ltd. All rights reserved.

2.12.1

Introduction

324

2.12.2
2.12.2.1
2.12.2.1.1
2.12.2.1.2
2.12.2.1.3
2.12.2.1.4
2.12.2.1.5
2.12.2.1.6
2.12.2.1.7
2.12.2.2
2.12.2.2.1
2.12.2.2.2
2.12.2.2.3
2.12.3
2.12.3.1
2.12.3.2
2.12.3.3
2.12.3.4

2.12.4
2.12.5
2.12.6
2.12.6.1
2.12.6.2
2.12.6.3
2.12.6.4
2.12.6.5
2.12.6.6
2.12.6.7
2.12.6.8
2.12.6.9
2.12.7
References

b-SiC Properties23
Mechanical Properties
Elastic modulus23
Poisson’s ratio23
Shear modulus23
Hardness23
Fracture toughness23
Fracture strength
Thermal creep23
Thermal Properties23
Thermal conductivity
Specific heat
Thermal expansion
SiC/SiC Composite
Fibrous Preform

Coating of Fibers
Infiltration of the SiC Matrix: The CVI Process
Infiltration of the SiC Matrix: The NITE Process
Properties of CVI SiC/SiC
Properties of NITE-SiC/SiC
Mechanical Behavior of CVI SiC/SiC
Tensile Stress–Strain Behavior
Damage Mechanisms
Ultimate Failure
Reliability
Interface Properties: Influence on the Mechanical Behavior
Fracture Toughness
Fatigue and High-Temperature Behavior
Thermal Shock
Creep Behavior
Concluding Remarks

325
325
325
325
325
325
325
326
326
326
326
326
327

327
327
327
327
328
328
330
330
330
331
333
333
334
335
336
336
336
337
337

Abbreviations
C/C
C/SiC
CMC
CVD
CVI

Carbon matrix composite reinforced by
carbon fibers
SiC matrix composite reinforced by

carbon fibers
Ceramic matrix composite
Chemical vapor deposition
Chemical vapor infiltration

LPS
MI
NITE
PIP
PyC
RS
SENB

Liquid phase sintering
Melt infiltration
Nanopowder infiltration and transient
eutectic-phase
polymer impregnation and pyrolysis
Pyrocarbon
Reaction sintering
Single edge notch bending

323


324

Properties and Characteristics of SiC and SiC/SiC Composites

SEP

SiC/SiC

Socie´te´ Europe´enne de Propulsion
SiC matrix composite reinforced by SiC
fibers

2.12.1 Introduction
Silicon carbide is composed of tetrahedra of carbon
and silicon atoms with strong bonds in the crystal
lattice. This produces a very hard and strong ceramic
with outstanding characteristics such as high thermal
conductivity, low thermal expansion, and exceptional
resistance to thermal shock and to corrosion in
aggressive environments at high temperatures. However, this implies a few inadequate characteristics for
structural applications, such as low fracture toughness, high sensitivity to the presence of microstructural flaws, brittle behavior, and lack of reliability.
Reinforcing with continuous SiC-based fibers allows
these weaknesses to be overcome. The composite
SiC/SiC that is obtained is damage tolerant, tough,
and strong, and it can be insensitive to flaws and
notches. The concept of composite material is very
powerful. Composites can be tailored to suit enduse applications through the sound selection and
arrangement of the constituents. Ceramic matrix composites (CMCs) reinforced with continuous ceramic
or carbon fibers are of interest in thermostructural
applications.1–4 They are lightweight and damage tolerant and exhibit a much greater resistance to high
temperatures and aggressive environments than metals
or other conventional engineering materials.
CMCs can be fabricated by different processing
techniques, using either liquid or gaseous precursors.
The chemical vapor infiltration (CVI) method can
produce excellent SiC/SiC composites with a highly

crystalline structure and excellent mechanical properties.5 The quality of the material obtained by the
polymer impregnation and pyrolysis (PIP) method is
insufficient. A novel processing technique (nanopowder infiltration and transient eutectic-phase processing, NITE) was claimed to achieve good material
quality.5–7
The SiC/SiC composites prepared using the
CVI method and reinforced with the latest nearstoichiometric SiC fibers (such as Hi-Nicalon type
S and Tyranno-SA3 fibers) appear to be promising
candidates for nuclear applications7–12 because of
their high crystallinity, high purity, near stoichiometry and radiation resistance of the b-phase of SiC,
as well as excellent resistance at high temperatures to

fracture, creep, corrosion, and thermal shock. Studies
on the b-phase properties suggest that CVI SiC/SiC
composites have the potential for excellent radiation
stability.3 CVI SiC/SiC is also considered for applications as structural materials in fusion power
reactors because of low neutron-induced activation
characteristics coupled with excellent mechanical
properties at high temperature.10–12
The CVI technique has been studied since the
1960s.13–19 It derives directly from chemical vapor
deposition (CVD).13–15 In very simple terms, the
SiC-based matrix is deposited from gaseous reactants
on to a heated substrate of fibrous preforms (SiC).15
CVI is a slow process, and the obtained composite
materials possess some residual porosity and density
gradients. Despite these drawbacks, the CVI process
presents a few advantages: (1) the strength of reinforcing fibers is not affected during the manufacture of
the composite; (2) the nature of the deposited material can be changed easily, simply by introducing the
appropriate gaseous precursors into the infiltration
chamber; (3) a large number of components; and

(4) large, complex shapes can be produced in a
near-net shape.
Development of CVI SiC/SiC composites began in
the 1980s when SEP (Socie´te´ Europe´enne de Propulsion), Amercorm, Refractory Composites, and others
began to develop equipment and processes for producing CVI components for aerospace, defense, and other
applications. The development of CVI SiC/SiC composites has been inspired by the poor oxidation resistance of their predecessor CVI C/C composites. CVI
SiC/SiC components have been produced and tested.
SNECMA (formerly SEP) is at the forefront of this
technology and has demonstrated satisfactory component performance in engine and flight tests.
The mechanical properties of SiC/SiC composites depend on the fiber–matrix interface. Pyrocarbon (PyC) has proved to be an efficient interphase
to control fiber–matrix interactions and composite
mechanical behavior.20 But PyC is sensitive to oxidation at temperatures above 450  C. A few versions of
high-temperature-resistant CVI SiC/SiC composites
have been produced. In order to protect the PyC
interphase against oxidation, multilayered interphases and matrices have been developed.3,21 Multilayered matrices contain phases that produce sealants
at high temperatures, preventing oxygen from reaching the interphase.22 This composite is referred to as
CVI SiC/Si–B–C. Oxidation-resistant interphases
such as BN or multilayered materials can also
be coated on the fibers. An ‘oxygen getter’ can be


Properties and Characteristics of SiC and SiC/SiC Composites

added to the matrix to scavenge oxygen that might
ingress into the matrix (enhanced CVI SiC/SiC).
The mechanical behavior of CMCs displays several typical features that differentiate them from the
other composites (such as polymer matrix composites, metal matrix composites, etc.) and from homogeneous (monolithic) materials. These features are
due to heterogeneous and multiscale composite
microstructure and the respective properties of the
constituents (interphases, fiber, and matrix). The

main characteristics of CVD SiC, CVI SiC/SiC,
and NITE-SiC/SiC are reviewed in this chapter.
Features of mechanical behavior of SiC/SiC are discussed with respect to microstructure, on the basis of
the large amount of work done on CVI SiC/SiC.

E ¼ E0 expðÀCVp Þ

325

½1Š

E0 ¼ 460 GPa for CVD SiC (polycrystalline, highpurity, very dense, and pore-free SiC material) and
C ¼ 3.57.
No significant difference was obtained between
the elastic moduli for a- and b-polycrystalline SiC
or among those of hot-pressed, sintered, and CVD
materials.
The elastic modulus at elevated temperatures has
been empirically expressed as:
E ¼ E0 À BT expðÀT0 =T Þ

½2Š

À1

with E0 ¼ 460 GPa, B ¼ 0.04 GPa K , and T0 ¼ 962 K.
2.12.2.1.2 Poisson’s ratio23

2.12.2 b-SiC Properties23
Silicon carbide has a myriad polytypes depending on

the varied stacking of closed atomic planes.23 Only
CVD SiC material is inherently highly crystalline,
pure, and stoichiometric, which is critical to irradiation stability. Much emphasis is placed on CVD SiC
in this chapter, as it corresponds very closely to the
matrix of CVI SiC/SiC. The reader will find further
details on the SiC structure–property relationships in
the excellent comprehensive review by Snead and
colleagues.23 Here the main data from Snead’s paper
are summarized.
Only the 3C–SiC crystal, known as b-SiC, has
the sequence showing cubic symmetry out of the
infinite number of variations. All the other polytypes
which show noncubic symmetry are classified as
a-SiC. a-SiC is formed above 2373 K and b-SiC at
1273–1873 K.
Various fabrication techniques, such as sintering,
direct conversion, gas-phase reaction, and polymer
pyrolysis, are currently used for the synthesis of SiC.
The CVD technique is one of the most familiar gasphase reaction methods for the synthesis of highly
crystalline, stoichiometric, high-purity b-SiC.
2.12.2.1

Mechanical Properties

The Poisson ratio of CVD SiC with excess residual
silicon yields the lowest value ($0.13). The highest
value of 0.21 was typically obtained for pure CVD
SiC. The temperature dependence is very minor.
2.12.2.1.3 Shear modulus23


The shear modulus at room temperature of 191 GPa for
CVD SiC has been determined by the four-point bending technique. This value was also derived from the
elastic modulus and Poisson’s ratio (n), using the conventional formula for isotropic solids: G ¼ E/2(1 þ n).
The temperature dependence of shear modulus can be
estimated from E by applying this formula.
2.12.2.1.4 Hardness23

There appears to be no significant difference between
Vicker’s and Knoop hardness: H $ 20.7–24.5 GPa has
been reported for CVD b-SiC. By contrast, slightly
higher values were obtained by nanoindentation.
Nanoindentation is known to yield local values
which depend on microstructural features. The aforementioned exponential function of porosity for elastic
modulus can be extended to the hardness evaluation:
HV ¼ 27:7 expðÀ5:4Vp Þ

½3Š

where HV is the Vicker hardness.
Currently, there is no high-temperature data
reported for high-purity CVD SiC.

2.12.2.1.1 Elastic modulus23

Generally, a dense and high-purity SiC material, for
example, CVD SiC, exhibits the highest elastic modulus; however, the elastic modulus decreases with
increasing porosity or impurity concentration. The
elastic modulus at room temperature is conventionally
expressed as an exponential function of porosity (Vp):


2.12.2.1.5 Fracture toughness23

Values between 2.4 and 5.1 MPa√m have been
measured for CVD b-SiC, depending on the test
technique employed and grain size. Fracture toughness of CVD SiC increases slightly at elevated temperatures. It does not exceed 6 MPa√m.


326

Properties and Characteristics of SiC and SiC/SiC Composites

ec ¼ Ap ðs=GÞn ðt =tÞp

2.12.2.1.6 Fracture strength

As is usual with brittle ceramics, fracture data exhibit
a significant scatter, as flaws that have a random
distribution induce fracture. An important consequence is that the fracture stress is not an intrinsic
characteristic. It is, instead, a statistical variable,
which depends on several factors including the test
method, the size of test specimens, and the number of
test specimens.24 Therefore, a universal reference
value of fracture strength cannot be recommended.
It is widely accepted that the Weibull model satisfactorily describes the statistical distribution of failure strengths:
 ð

m
½4Š
P ¼ 1 À exp À ðs=s0 Þ dV =V0
where P is the probability of failure, s is the stress,

s0 is the scale factor, m is the Weibull modulus, V is
the volume of specimen, and V0 is a reference volume
(1 m3 is generally used); m reflects the scatter in data,
and s0 is related to the mean value of the strength.
The strength data for a given geometry and stress
state can be determined using eqn [4]. However, m,
s0, and V0 must be available. It is important to note
that the estimate of s0 depends on V0.24 It will be
substantially different if V0 ¼ 1 m3 or 1 mm3. This
dependence is ignored in most publications, even in
the work by Snead and coworkers23 in which a number
of s0 values are reported. When V0 is not given, the
estimate of s0 is meaningless. The strength cannot be
determined safely. Unfortunately, reliable s0 values
(characteristic strength in a few papers) cannot be
recommended here until the authors have completed
their papers. The values of Weibull modulus of CVD
SiC at room temperature reported in Snead et al.23 span
a large range, from 2 to 12. The following values were
measured using tensile tests on CVI SiC/SiC minicomposites: m ¼ 6.1, s0 ¼ 10.5 MPa (V0 ¼ 1 m3).25,26
2.12.2.1.7 Thermal creep23

Primary and secondary creep deformations have been
reported in the literature for CVD SiC (high-purity and
polycrystalline b-SiC). Creep in SiC is highly dependent on the crystallographic orientation. The loading
orientation of 45 from the CVD growth axis is the
direction in which the most prominent creep strain is
observed. A review of creep behaviors of stoichiometric
CVD SiC has been provided by Davis and Carter.27
Primary creep of CVD SiC occurs immediately

upon loading and tends to saturate with time. The
primary creep strain generally obeys the following
relationship:

½5Š

where Ap , p, and t are creep parameters, and t is
the time elapsed. n ¼ 1.63, Ap ¼ 29, p ¼ 0.081, and
t ¼ 0.0095 s for the temperature of 1923 K. These parameters are for the loading orientation of 45 from
the CVD growth axis. In severe conditions, primary
creep strain in the CVD SiC can reach as high as 1%.
Steady-state creep rates for polycrystalline materials have been measured only above $1673 K, when
the stress axis is 45 inclined from the deposition
direction; temperatures as high as 2023 K are
required when the stress axis is parallel to the deposition direction. The strain rate is given by a powerlaw creep equation:
de=dt ¼ As ðs=GÞn expðÀQ =kb T Þ

½6Š

where As ¼ 2.0 Â 103, n ¼ 2.3, Q ¼ 174 kJ molÀ1 (activation energy), s is the applied stress, G is the shear
modulus, and kb is the Boltzmann constant.
2.12.2.2

Thermal Properties23

2.12.2.2.1 Thermal conductivity

It is reasonable to assume that the single-crystal form
of SiC, compared to the other varieties, exhibits the
highest thermal conductivity. However, high-purity

and dense polycrystalline CVD SiC exhibits practically the same conductivity as the single-crystal
material. It is worth noting that the impurity content
of the very high thermal conductivity CVD SiC materials is negligibly small, and this material has near
theoretical density ($3.21 g cmÀ3). The curve-fitting
to the single-crystal SiC data above 300 K yields
an upper limit of the thermal conductivity of SiC
(in W mÀ1 KÀ1):
Kp ¼ ðÀ0:0003 þ 1:05  10À5 T ÞÀ1

½7Š

2.12.2.2.2 Specific heat

The temperature dependence of the specific heat
can be treated in two temperature regions: a rapid
increase at low temperatures (below 200 K), and a
gradual increase at higher temperatures. No systematic difference can be distinguished between the structural types. The specific heat, Cp (in J kgÀ1 K), over the
temperature range 200–2400 K can be approximately
expressed as
Cp ¼ 925:65 þ 0:3772T À 7:9259 Â 10À5 T 2
À 3:1946 Â 107 =T 2

½8Š


Properties and Characteristics of SiC and SiC/SiC Composites

The specific heat of SiC at room temperature is taken
as 671 Æ 47 J kgÀ1 K.
2.12.2.2.3 Thermal expansion


The coefficient of thermal expansion for b-SiC has
been reported over a wide temperature range. The
average value in the interval from room temperature
to 1700 K is a ¼ 4.4 Â 10À6 KÀ1.
At higher temperatures
5 Â 10À6 KÀ1
At lower temperatures
a ¼ 2.08 þ 4.51 Â 10À3T.

(T > 1273 K),

(550 < T < 1273 K),

It is worth addressing the processing method first
because this information is useful for a better understanding of the structure of SiC/SiC. The manufacture of long fiber-reinforced composites requires
three main steps14,15,28,29:
1. preparation of fibrous preform,
2. fiber coating, which provides an interface material
(interphase), and
3. infiltration of the matrix.
Fibrous Preform

The preforms of SiC/SiC composites are made of
refractory SiC-based continuous fibers. The latest
near-stoichiometric SiC fibers (such as Hi-Nicalon
type S and Tyranno-SA3 fibers) are the most appropriate for those CVI SiC/SiC foreseen for nuclear
applications. These fibers exhibit high strength, high
stiffness, low density, and high thermal and chemical
stability to withstand long exposures at high temperatures.30 Finally, the fiber diameter must be small

(<20 mm) so that the fibers can be woven easily.
The fiber preforms may consist of
1. A simple stack of unidirectional fiber layers or
fabrics (1D or 2D preforms).
2. A multidirectional fiber architecture (3D preforms).
Weaving in four or five directions can also be used.
The 2D layers are stacked and kept together using a
tool or using fibers in the orthogonal direction (3D
preforms).
2.12.3.2

a multilayer ((PyC/SiC)n or (BN/SiC)n sequences).
PyC-based interphases have been the subject of
extensive studies and have been shown to be the
most appropriate with respect to controlling crack
deflection and mechanical properties. With the CVI
process, the gas precursor is CH4 for carbon, and
BCl3 and NH3 for boron nitride. Multilayered interphases may be deposited via pulsed CVI.

a¼

2.12.3 SiC/SiC Composite

2.12.3.1

327

Coating of Fibers

An interface material is deposited on the fibers. This

interphase acts as a deflection layer for the matrix
cracks. It consists essentially of PyC, boron nitride, or

2.12.3.3 Infiltration of the SiC Matrix:
The CVI Process
The basic chemistry of making a coating and a matrix
by CVI is the same as that of depositing a ceramic on a
substrate by CVD.13–15 The reactions consist of cracking a hydrocarbon for deposition of carbon and cracking of methylchlorosilane for deposition of SiC. In the
I-CVI process (isobaric isothermal CVI) the preform
is kept in a uniformly heated chamber. Temperature
and pressure are relatively low (<1200  C, <0.5 atm).
A few alternative CVI techniques have been proposed to increase the infiltration rate.15,28,29 These
techniques require more complicated CVI chambers
and are not appropriate to the production of large or
complex shapes or a large number of pieces.
The forced CVI (F-CVI) technique was proposed in
the mid-1980s.29 The precursor gas is forced through
the bottom surface of the preform under a pressure P1,
and the exhaust gases are pumped from the opposite
face under a pressure P2 < P1. The fibrous preform is
heated from the top surface and sides, and cooled from
the bottom (cold) surface. The densification times are
significantly shorter when compared to I-CVI (10–24 h
for a SiC matrix, a few hours for carbon), and the
conversion efficiency of the precursor is relatively
high. However, the technique is not appropriate for
complex shapes. Only one preform per run can be
processed, and complex graphite fixtures are required
to generate the temperature and pressure gradients.
In order to overcome the aforementioned limitations of the F-CVI technique, alternative techniques

using thermal gradients or pressure gradients have
been examined for many years.15 In the thermal
gradient process, the core of the fibrous preform is
heated in a cold-wall reactor. The heat loss by radiation is favorable to get a lower temperature in the
external surface. The densification front advances
progressively from the internal hot zone toward the
cold side of the preform. In the P-CVI process, the
source gases are introduced during short pulses.15
The P-CVI process is appropriate for the deposition
of thin films or multilayers.


328

Properties and Characteristics of SiC and SiC/SiC Composites

2.12.3.4 Infiltration of the SiC Matrix:
The NITE Process
Reaction sintering (RS), liquid phase sintering (LPS),
PIP, melt infiltration (MI), and their hybrid processes
are alternative options. PIP requires development of
a near-stoichiometric polymer precursor. The other
methods have issues in phase and uniformity control.
The NITE process is based on LPS,5,7,30 which
has been improved owing to the progress in reinforcing fibers and availability of fine nano-SiC powders.
A slurry of b-SiC nanopowders and additives is infiltrated into SiC fabrics and dried for making prepreg
sheets. After the layup of the sheets, hot pressing is
applied to make NITE-SiC/SiC. Small amounts of
sintering aids (Al2O3, Y2O3, SiO2), high temperatures
(1750–1800  C), and pressures ranging from 15 to

20 MPa are required for matrix densification. The
NITE process was claimed to present great advantages such as flexibility in the shape and size of the
components.7 The successful development of NITE
is due to appropriate fiber protection and the emergence of advanced SiC fibers such as Tyranno-SA3.

Table 1
fabrics

2.12.4 Properties of CVI SiC/SiC
Table 1 is a complete list of the mechanical and
thermophysical properties of first generation 2D
CVI SiC/SiC composites reinforced with SiC Nicalon fibers of first generation.2,31 An average strain-tofailure of 0.3% and a tensile strength of 200 MPa
have been reported. Higher strengths and strainsto-failure appear in Tables 2 and 3, which give the
available properties measured on other generations
of SiC/SiC composites reinforced with advanced
Hi-Nicalon or Hi-Nicalon type S fibers.3,32,33 The
behavior of stronger Nicalon-reinforced SiC/SiC
is discussed in a subsequent section. It can be noted
that the strain-to-failure can reach 1%, and the tensile strength can exceed 300 MPa. As discussed in
a subsequent section, a high strain-to-failure can be
obtained when the performances of the reinforcing
tows and the load transfers during loading have not
been impaired as a result of the processing conditions. Ideally, the strain-to-failure should coincide
with that of reinforcing tows, that is, about 0.8%.

Mechanical and thermophysical properties of 2D SiC/SiC composites reinforced with 0/90 balanced Nicalon™

Property

Fiber content (%)

Specific gravity
Porosity (%)
Tensile strength (MPa)
Strain-to-failure (%)
Young’s modulus (GPa)
Poisson’s ratio
n12
n13
Flexural strength (MPa)
In-plane compressive strength (MPa)
Thru-the-thickness compressive strength (MPa)
Interlaminar shear strength (MPa)
In-plane thermal diffusivity (10À5 m2 sÀ1)
Thru-the-thickness thermal diffusivity (10À5 m2 sÀ1)
In-plane coefficient of thermal expansion (10À6 KÀ1)
Thru-the-thickness coefficient of thermal expansion (10À6 KÀ1)
Fracture toughness (MPa√m)
Specific heat (J kgÀ1 KÀ1)
Total emissivity
In-plane thermal conductivity (W mÀ1 KÀ1)
Thru-the-thickness thermal conductivity (W mÀ1 KÀ1)

Temperature
23  C

1000  C

1400  C

40

2.5
10
200
0.3
230

40
2.5
10
200
0.4
200

40
2.5
10
150
0.5
170

0.5
0.18
300
580
420
40
12
6
3
1.7

30
620
0.8
19.0
9.5

400
480
380
35
5
2
3
3.4
30
1200
0.8
15.2
5.7

280
300
250
25
5
2

30
0.8


Source: Choury, J. J. Thermostructural composite materials in aeronautics and space applications. In Proceedings of GIFAS Aeronautical
and Space Conference, Bangalore, Delhi, India, Feb 1989; pp 1–18; Lacombe, A.; Rouge`s, J. M. In AIAA’90, Space Program and
Technologies Conference’90, Huntsville, AL, Sept 1990; The American Institute of Aeronautics and Astronautics: Washington, DC, 1990;
AIAA-90-3837.


Properties and Characteristics of SiC and SiC/SiC Composites

Table 2
Mechanical properties of a CVI SiC/Si–B–C
composite with a self healing matrix and a multilayer reinforcement of Hi-Nicalon™ fibers, and 2D CVI-enhanced
SiC/SiC composite reinforced with 0/90 five harness satin
fabrics of Hi-Nicalon™ fibers
Property

CVI SiC/Si–B–C
Fiber type

Temperature
Room
temperature

1200  C

Hi-Nicalon™
fibers
Plain weave
2.3
13
315


Hi-Nicalon™
fibers
Plain weave

Reinforcement
Density
Porosity (%)
Tensile strength
(MPa)
Strain-to0.5
failure (%)
Young’s modulus
220
(GPa)
Interlaminar shear
Strength (MPa)
31
Flexural strength
699
(MPa)
2D CVI-enhanced SiC/SiC composite
Fiber type
Hi-Nicalon™
Fiber content (%)
35
Reinforcement
0/90 five harness
satin
Density

2.2
Porosity (%)
10
Tensile strength
324
(MPa)
Strain-to-failure
0.74
(%)
Young’s modulus
207
(GPa)

329

Table 3
Room-temperature properties of 2D melt infiltrated CVI SiC/SiC and 2D CVI SiC/SiC composites reinforced with Hi-Nicalon type S fibers
2D melt infiltrated CVI SiC/SiC
Fiber type
Fiber content (%)
Density
Tensile strength (MPa)
Strain-to-failure (%)
Young’s modulus (GPa)
2D CVI SiC/SiC
Fiber type
Fiber content (%)
Density
Tensile strength (MPa)
Strain-to-failure (%)

Young’s modulus (GPa)
45 off-axis tensile strength (MPa)
45 off-axis strain-to-failure (%)

Hi-Nicalon type S™
35
2.2
341–412
0.60
232–262
Hi-Nicalon type S™
35
2.25
305
0.60
214
167
0.66

Source: Morscher, G.; Pujar, V. Int. J. Appl. Ceram. Technol.
2009, 6, 151–163.

23
620
Hi-Nicalon™
35
0/90 five
harness satin
2.2
10

259
0.50
212

Source: Bouillon, E.; Habarou, G.; Spriet, P.; et al.
Characterization and nozzle test experience of a self sealing
ceramic matrix composite for gas turbine applications. In
Proceedings of IGTI/ASME TURBO EXPO Land, Sea and Air 2002,
Amsterdam, The Netherlands, June 3–6, 2002; Power Systems
Composites Datasheet.

The strain-to-failure is an interesting characteristic
for CMCs for several reasons. First of all, it is not
sensitive to scale effects, so that it may be regarded as
an intrinsic property and so various CMCs can be
compared easily. Then, it reflects the degree of
damage tolerance, whereas the strength reflects the
load-carrying capacity. These characteristics need to
be differentiated, as most components are usually
subjected to strain-controlled loading conditions.
A fracture toughness of 30 MPa√m was measured
using conventional techniques designed for monolithic materials. It can be regarded as a high value
when compared to monolithic SiC. However, it is

worth pointing out that it represents the fracture
toughness of an equivalent homogeneous material.
As discussed in a subsequent section, critical stress
intensity factor (KIC) is not an intrinsic property, and
it is not an appropriate concept for long fiberreinforced composites. Furthermore, besides the
resistance to crack propagation, damage tolerance is

an important property for CMCs. It cannot be characterized by fracture toughness. This situation is new
when compared to homogeneous materials. Anyway,
the fracture toughness KIC may be regarded as an
index to compare materials. It cannot be used for
design purposes for the aforementioned reasons.
Table 1 shows that CVI SiC/SiC retains its properties at high temperatures. These properties can be
enhanced by using advanced fibers. Durability will be
addressed in a subsequent section.
Properties vary according to factors, including
preform architecture, fiber type, matrix properties,
fiber–matrix bond strength, loading conditions, etc.
For instance, high tensile strengths (up to 400 MPa)
were obtained with Hi-Nicalon™ SiC fibers,34 or
with Nicalon fibers and rather strong interfaces.35
Further details on microstructure versus properties
are discussed in subsequent sections. The mechanical
behavior of 2D CVI SiC/SiC composites exhibits
features that are related to composite microstructure.
Thus, it deserves special attention because it differs
significantly from that of the more conventional
homogeneous materials. A clear understanding will
be beneficial to a sound use of CVI SiC/SiC.


Properties and Characteristics of SiC and SiC/SiC Composites

Tables 3 and 4 show that the ultimate strength
and Young’s modulus tend to decrease under off-axis
tensile conditions.36 It is worth pointing out that the
strain-to-failure is an invariant. It is interesting to note

that in 2D CVI SiC/SiC, the directions of the principal stresses coincide with those of the fiber tows.

2.12.5 Properties of NITE-SiC/SiC
The matrix of NITE-SiC/SiC comprises polycrystalline SiC and a small amount of isolated oxides.
The microstructure is highly crystalline and highly
dense. Table 5 lists the typical available properties of NITE-SiC/SiC.7 Thermal conductivity
($30 W mÀ1 KÀ1) is quite high when compared to
CVI SiC/SiC (below 15 W mÀ1 KÀ1) reinforced with
either Nicalon (Table 1) or Hi-Nicalon fibers.37 The
high proportional stress limit is claimed to be an interesting feature.7 However, it is worth pointing out that
it reflects a high load-carrying capacity. By contrast,
Table 4
Off-axis properties of a first generation of 2D
CVI SiC/SiC reinforced with Nicalon fibers
Property

0

20

45

Strain-to-failure (%)
Proportional limit (MPa)
Saturation stress (MPa)
Tensile strength (MPa)
Young’s modulus (GPa)

0.3
80

150
190
220

0.3
70
145
170
210

0.3
70
145
170
210

Source: Aubard, X.; Lamon, J.; Allix, O. J. Am. Ceram. Soc. 1994,
77, 2118–2126.

Table 5
Room-temperature properties of NITE-SiC
composites

the low strain-to-failure indicates a limited damage
tolerance. The strain-to-failure does not increase after
aging at high temperatures up to 1500  C. This trend is
consistent with the strong fiber–matrix interactions
induced by the surface roughness of Tyranno-SA3
fibers.38 A comprehensive database on properties of
NITE-SiC/SiC is not available. NITE-SiC/SiC has

been reported to retain ultimate strength and a proportional stress limit after exposure at temperatures up to
1300  C.6

2.12.6 Mechanical Behavior of CVI
SiC/SiC
2.12.6.1

400

300
(b)
200

100

0

0.2

0.4

0.6

0.8

1

1.2

Longitudinal tensile strain (%)


Figure 1 Typical tensile stress–strain behaviors measured
on 2D SiC/SiC composites possessing PyC-based
interphases and fabricated from untreated or treated
Nicalon (ceramic grade) fibers: (a) strong fiber/coating
interfaces and (b) weak fiber/coating interfaces.

200
TyrannoSA3
UD
53
3.11
0.6
358
408
0.13
358
32

TyrannoSA3
Cross plied
51
3.06
3.8
148
167
0.08
288

Kohyama, A. In Ceramic Matrix Composites; Krenkel, W., Ed.;

Wiley-VCH: Weinheim, Germany, 2008; Chapter 15, pp 353–384,
reproduced with permission.

150
Stress (MPa)

Reinforcement
Fiber content (%)
Density
Porosity (%)
Proportional limit (MPa)
Tensile strength (MPa)
Strain-to-failure (%)
Young’s modulus (GPa)
Thermal conductivity
(W mÀ1 KÀ1)

(a)

0

Property
Fiber type

Tensile Stress–Strain Behavior

Figures 1 and 2 summarize the typical stress–strain
behavior of 2D CVI SiC/SiC composites. The behavior is initially linear under strains below 0.03%. Then,

Longitudinal tensile stress (MPa)


330

100

50

0

0

0.1

0.2

0.3
0.4
Strain (%)

0.5

Figure 2 Typical tensile stress–strain behaviors
measured on 2 different test specimens (2D SiC/SiC
reinforced with Hi-Nicalon S fibers).

0.6


Properties and Characteristics of SiC and SiC/SiC Composites


the nonlinear deformations result essentially from
transverse cracking in the matrix (the cracks are perpendicular to fibers oriented in the loading direction).
Saturation of matrix damage is indicated by the end of
the curved domain marked by a point of inflection.
Then the ultimate portion of the curve reflects the
deformation of fibers. Fiber failures may initiate prior
to ultimate fracture. Such mechanical behavior is
essentially damage-sensitive.
A damage-sensitive stress–strain behavior is obtained
when the initial contribution of the matrix to load
carrying is significant. The elastic modulus of the
matrix (Em) is not negligible when compared to that
of the fiber (Ef). Its contribution to the modulus of the
composite (Ec) is illustrated by the mixtures law,
which provides satisfactory trends for continuous
fiber-reinforced composites:

1.2

½9Š

where Vm is the volume fraction of matrix and Vf is
the volume fraction of fibers oriented in the loading
direction in a 2D woven composite.
In 2D CVI SiC/SiC composites, Em (%410 GPa)
> Ef (200–380 GPa), Vm $ Vf the initial contribution
of the matrix to Ec is significant. Then, as it decreases
when the matrix cracks, the behavior becomes controlled by the tows. The 2D SiC/SiC composites
exhibit an elastic damageable behavior (Figure 3).
This means that the response of the damaged material is elastic as indicated by the linear portion of the

curves on reloading. Figure 4 shows the dependence
of the elastic modulus on damage.

Damage Mechanisms

The basic damage phenomena in unidirectional composites under on-axis tensile loads involve multiple
microcracks or cracks that form in the matrix perpendicular to fiber direction and that are arrested by the
fibers by deflection in the fiber–matrix interface. In the
composites reinforced with fabrics of fiber bundles,
matrix damage is influenced by a multilength scale
structure.39 Furthermore, 2D CVI SiC/SiC is a heterogeneous medium because of the presence of fibers,
large pores (referred to as macropores) located between
the plies or at yarn intersections within the plies, and a
uniform layer of matrix over the fiber preform (referred
to as the intertow matrix) (Figure 5). Much smaller

1.0
0.8
E / E0

Ec ¼ Em Vm þ Ef Vf

2.12.6.2

331

0.6
0.4
A
0.2


F
D

0

0

0.2

0.4

0.6

EfVf
G 2.E
0
0.8

1.0

Strain (%)
Figure 4 Relative elastic modulus versus applied strain
during tensile tests on various 2D woven SiC/SiC composites
reinforced with treated fibers: (A) Nicalon/(PyC20/SiC50)10/
SiC, (D) Nicalon/PyC100/SiC, (F) Hi-Nicalon/PyC100/SiC,
(G) Hi-Nicalon/(PyC20/SiC50)10/SiC.

350
Longitudinal tow


300
Macropore

Transversal tow

Stress (MPa)

250
200
150
100
8 ϫ 30
16 ϫ 120

50
0

0

0.2

0.4
0.6
Strain (%)

0.8

1


Layer
0.5 mm

Figure 3 Stress–strain curves in tension of 2D SiC/SiC
reinforced with treated Nicalon fibers. The open and filled
symbols represent ultimate failure data point obtained with
the specimens of volumes V1 and V2, respectively.

Figure 5 Micrograph showing the microstructure of a
2D CVI SiC/SiC composite.


332

Properties and Characteristics of SiC and SiC/SiC Composites

Longitudinal strain = 0.06%

Longitudinal strain = 0.2%

Longitudinal strain = 0.6%

Longitudinal strain = 0.4%

Longitudinal strain = 0.8%

Figure 6 Schematic diagram showing matrix cracking in a 2D SiC/SiC composite during a tensile test.

pores are also present within the tows. Under on-axis
tension, damage in 2D CVI SiC/SiC occurs essentially

in the formation of matrix cracks perpendicular to
longitudinal fiber axis and their deflection either
by the tows (first and second steps) or by the fibers
within the tows (third step). These steps (Figure 6)
correspond to deformation increments:
Step 1: cracks initiate at macropores where stress
concentrations exist (deformations between
0.025% and 0.12%);
Step 2: cracks form in the transverse yarns and in
the interply matrix (deformations between 0.12%
and 0.2%);
Step 3: transverse microcracks initiate in the longitudinal tows (deformations larger than 0.2%).
These microcracks are confined within the longitudinal tows. They do not propagate in the
rest of the composite. The matrix in the longitudinal tows experiences a fragmentation process and the crack spacing decreases as the load
increases.
As mentioned earlier, the directions of principal
stresses are dictated by fiber orientation rather than
by the loading direction. Thus, under on-axis conditions, all the matrix cracks are perpendicular to
the loading direction. Then, under off-axis tension,
matrix cracks that are located in the tows are

perpendicular to fiber direction, whereas those
located between the tows are perpendicular to the
load direction. On-axis loading conditions are discussed later.
The resulting Young’s modulus decrease illustrates
the importance of damage in the mechanical behavior
(Figure 4). The major modulus loss (70%) is caused
by both the first families of cracks located on the
outside of the longitudinal tows (deformations
<0.2%). By contrast, the microcracks within the longitudinal tows are responsible for only a 10% loss. The

substantial modulus drop reflects important changes
in load sharing: the load gets carried essentially by
the matrix-coated longitudinal tows (tow reloading).
During microcracking in the longitudinal tows, load
sharing is affected further, and the load becomes
carried essentially by the filaments (fiber reloading).
The elastic modulus reaches a minimum described by
the following equation (Figure 4):
Emin ¼ 1=2Ef Vf

½10Š

where Vf is the volume fraction of fibers.
Equation [10] implies that the matrix contribution
is negligible. At this stage, matrix damage and
debonding are complete (saturation). The load is
carried by fibers only. The mechanical behavior is
controlled by the fiber tows oriented in the direction
of loading.


Properties and Characteristics of SiC and SiC/SiC Composites

Ultimate Failure

Ultimate failure generally occurs after saturation
of matrix cracking. The fibers break when the
applied load is close to the maximum. Matrix damage
and ultimate failure thus appear to be successive
phenomena.

The ultimate failure of a tow of parallel fibers
involves two steps:
 a first step of stable failure and
 a second step of unstable failure.
During the first step, the fibers fail individually as the
load increases. In the absence of fiber interactions,
the load is carried by the surviving fibers only (equal
load sharing). Fiber interactions cause tow weakening. The ultimate failure of a tow (second step)
occurs when the surviving fibers cannot tolerate the
load increment resulting from a fiber failure. At this
stage, a critical number of fibers have been broken.
The ultimate failure of a longitudinal tow coated with
matrix also involves a two-step mechanism and global
load sharing when a fiber fails. In the presence of
multiple cracks across the matrix and associated
interface cracks, the load-carrying capacity of the
matrix is tremendously reduced or eliminated. The
matrix-coated tows behave like dry tows subject to
the typical stress field generated by the presence of
matrix cracks. The ultimate failure of a matrix-coated
tow occurs when a critical number of fibers have
failed. This mechanism operates in the tows within
textile CVI SiC/SiC composites. The ultimate failure
of the composite is caused by the failure of a critical
number of broken tows (!1) depending on the stress
state: $1 under an axial tension, >1 in bending.
It is worth pointing out that the failure mechanism
of CVI SiC/SiC composites differs from that observed
in polymer matrix impregnated tows, where local load
sharing prevails when a fiber fails. In these composites,

the fibers fail first. Therefore, the uncracked matrix is
able to transfer the loads.
2.12.6.4

Reliability

The ultimate failure of CVI SiC/SiC composites is
highly influenced by stochastic features. As fibers
are brittle ceramics, they are sensitive to the presence
of flaws (stress concentrators) that are distributed
randomly. As a consequence, the strength data exhibit
significant scatter, as illustrated by Figure 7.39,40 The
figure shows that the magnitude of the strength and
scatter decrease from single fibers to tows, then to
infiltrated tows, and finally to woven composites.

0.012
Composites (2D)
0.01
Density (MPa–1)

2.12.6.3

333

0.008
0.006

Composites (1D)


0.004
Tows
0.002

Fibers

0
0

500

1000
1500
Stress (MPa)

2000

2500

Figure 7 Strength density functions for SiC fibers (NLM
202), SiC fiber tows, SiC/SiC (1D) minicomposites, and 2D
SiC/SiC composites.

As a result of the previously mentioned two-step
failure mechanism, the ultimate failure of an entity is
dictated by the lowest extreme of the strength distribution pertinent to its constituent: that is, tows versus
filaments, infiltrated tows versus fibers, and 2D composites versus infiltrated tows. The lowest strength
extremes correspond respectively to the critical number of individual fiber breaks (%17% for the SiC
Nicalon™ fibers and for the SiC Hi-Nicalon™ fibers)
and to the critical number of tow failures (!1). The

gap between tows and SiC infiltrated tows results
from the method of strength determination: the critical
number of individual fiber breaks was taken into
account for tow strength determination, whereas the
strength of infiltrated tows and composites was underestimated because the total cross sectional area of the
specimens was used.
The flaw populations are truncated during the
successive damage steps, which leads to a homogeneous ultimate population of flaws.40 This process of
progressive elimination of flaws governs the trends in
the ultimate failure. The tensile stress–strain curves
obtained on a batch of several CVI SiC/SiC test
specimens coincide quite well (Figure 5), whereas
the strength data exhibit a certain scatter (Figure 5).
This scatter is limited (Figure 8). Dependence of
composite strength on the stressed volume is not
significant (Figure 8). Furthermore, dependence on
the loading conditions is not so large (Figure 9): for
instance, the flexural strength is 1.15 times as large as
the tensile strength40,41 when measured on specimens
having comparable sizes (Figure 9).
The Weibull model is not appropriate to describe
the volume dependence of strength data,40 as the
weakest link concept is violated. However, the


334

Properties and Characteristics of SiC and SiC/SiC Composites

1


Probability

0.8
0.6

0.4
0.2

0
200

250

300

350

Ultimate stress (MPa)
Figure 8 Scale effects in 2D woven SiC/SiC composites.
Influence of specimen dimensions on ultimate failure in
tension: () 8 Â 30 mm2 and (○) 160 Â 120 mm2.
Reproduced from Bansal, N. P., Ed. Handbook of Ceramics
and Glasses; Kluwer Academic: New York, 2005, with
permission from Springer.

1
Tensile tests

Failure probability


0.8
Three-point
bending
tests

0.6

0.4
Four-point
bending
tests

0.2

0
260

280

300
320
340
Maximum stress (MPa)

360

380

Figure 9 Strength distributions for 2D woven SiC/SiC

composites tested under various loading conditions:
tension, three-point bending and four-point bending.
Reproduced from Bansal, N. P., Ed. Handbook of Ceramics
and Glasses; Kluwer Academic: New York, 2005, with
permission from Springer.

Weibull modulus (m) can be extracted from the statistical distribution of the strength data: m is in the
range of 20–29. This value provides an evaluation of
the scatter in strength data. It reflects a small scatter.
2.12.6.5 Interface Properties: Influence on
the Mechanical Behavior
The fiber–matrix interfacial domain is a critical part
of composites because load transfers from the matrix

to the fiber and vice versa occur through the interface.
Most authors support the concept of weak interfaces to
increase fracture toughness. They assign toughening to
crack-bridging and fiber pullout. Weak interfaces are
detrimental to the composite strength. A high strength
requires that the matrix carry a part of the load. This is
obtained with strong interfaces, which implies that
the deflection cracks at interfaces are short and/or
that significant sliding friction takes place. These latter
requirements, to be met for strong composites, are
therefore incompatible with the former ones for
tough composites, if toughening is based solely upon
crack-bridging and fiber pullout.
Fiber–matrix interfaces exert a profound influence on the mechanical behavior and lifetime of
composites. Efforts have been directed toward optimization of interface properties. Fiber–matrix interfaces in CVI SiC/SiC composites consist of a thin
coating layer (<1-mm thick) of one or several materials deposited on the fiber (interphase). CVI SiC/SiC

composites with rather strong interfaces have been
obtained using fibers that have been treated in order
to increase the fiber/coating bond.35,42 The concept
of strong interfaces has been established on CVI
SiC/SiC composites with PyC and multilayered
(PyC/SiC)n fiber coatings. Less interesting results
have been achieved with BN interphases.43 Table 6
gives the various values of the interfacial shear stresses measured using various methods on CVI SiC/SiC
composites with PyC-based fiber coatings: the interfacial shear stresses range between 10 and 20 MPa for
the weak interfaces, whereas they are larger than
100–300 MPa for the strong interfaces.43–47
In the presence of a weak bond between the fibers
and the matrix or coating, single, long interface
cracks are created during matrix cracking (adhesive
failure type, Figure 10). The associated interface
shear stresses are low, and load transfers through
the interface crack are poor. The matrix is subjected
to low stresses and the volume of matrix that may
experience further cracking is reduced by the presence of the long interface cracks. The matrix
crack density is small. The crack spacing at saturation
as well as the pull-out length tends to be long
(>100 mm). Toughening results essentially from sliding friction along the cracked interfaces. However, as
a result of matrix unloading due to long interface
cracks, the fibers carry most of the load, which
reduces the composite strength. The corresponding
tensile stress–strain curve exhibits a narrow curved
domain limited by a stress at matrix saturation which
is distinctive of the ultimate strength (Figure 1).



Properties and Characteristics of SiC and SiC/SiC Composites

335

Table 6
Interfacial shear stresses (MPa) measured using various methods on 2D SiC/SiC composites with PyC-based
fiber coatings and reinforced with either as-received or treated fibers
SiC/C/SiC SiC/
(C/SiC)n/SiC

Interphase

2D woven
Microcomposites
Minicomposites
2D woven

PyC (0.1)
PyC (0.1)
PyC (0.1)
PyC (0.5)
(PyC/SiC)2
(PyC/SiC)4

12

2D woven

PyC (0.1)
PyC (0.5)

(PyC/SiC)2
(PyC/SiC)4

203

Crack
spacing

Crack
spacing

Tensile tests
(hysteresis
loops)

Untreated fibers
0.7
3
21–115
4
2
9
Treated fibers
140
190
370
150
90

Tensile tests

(curved
domain)

Push-out tests
(curved domain)

Push-out
tests
(plateau)

14–16
31
28

12–10
19.3
12.5

8

4–20
40–80

165–273
100–105
133
90

Source: Rebillat, F.; Lamon, J.; Guette, A. Acta Mater. 2000, 48, 4609–4618; Lamon, J.; Rebillat, F.; Evans, A. G. J. Am. Ceram. Soc.
1995, 78, 401–405; Lissart, N.; Lamon, J. Acta Mater. 1997, 45, 1025; Rebillat, F.; Lamon, J.; Naslain, R.; Lara-Curzio, E.; Ferber, M. K.;

Besmann, T. J. Am. Ceram. Soc. 1998, 81, 965; Rebillat, F.; Lamon, J.; Naslain, R.; Lara-Curzio, E.; Ferber, M. K.; Besmann, T. J. Am.
Ceram. Soc. 1998, 81, 2315–2326.

Debond crack

ld

Matrix

Coating

Fiber

(a)
Debond crack

t

ld

t
Matrix

Coating

In the presence of stronger fiber/coating bonds,
the matrix cracks are deflected within the coating
(cohesive failure type, Figure 10) into short and
branched multiple cracks. Short interphase cracks as
well as improved load transfers allow further cracking

of the matrix via a scale effect, leading to a higher
density of matrix cracks (which are slightly opened).
Sliding friction within the coating as well as multiple
cracking of the matrix increases energy absorption,
leading to toughening. Short interphase cracks and
improved load transfers reduce the load carried by
the fibers, leading to strengthening. The associated
tensile stress–strain curve exhibits a wide, curved
domain and the stress at matrix cracking saturation
is close to the composite strength (Figure 1).
The interphase is ineffective when fiber surface
is too rough, although deflection of matrix cracks
occurs. Because of strong fiber–matrix interactions
in the interface cracks, premature fracture of composite occurs under small strains close to the strain
at proportional limit. This phenomenon is observed
on CVI SiC/SiC reinforced with Tyranno-SA3
fibers.38
2.12.6.6

Fracture Toughness

Fiber

(b)
Figure 10 Schematic diagram showing crack
deflection when the fiber coating/interface is strong (a)
or weak (b).

The CVI SiC/SiC composites develop a network of
matrix cracks under load. The density of matrix

cracks is enhanced by rather strong interfaces: the
crack spacing may be as small as 10–20 mm whereas it
is at least 10 times larger in the presence of rather


336

Properties and Characteristics of SiC and SiC/SiC Composites

weak interfaces. Matrix cracking is an alternative
mechanism of energy dissipation.
A process zone of diffused matrix microcracks is
generated at the notch tip or at the tip of a preexisting
main macroscopic crack. Extension of this crack
results from the random failures of fiber bundles
located within the process zone.35 Due to the presence of a more or less large process zone associated
with a jagged crack, a crack length cannot be defined
and conventional concepts of fracture mechanics
are not appropriate (stress intensity factor) or cannot
be easily determined (strain energy release rate,
J-integral). Although the validity of the stress intensity factor concept to measure fracture toughness
is questionable, this is an interesting characteristic
for comparing CVI SiC/SiC composites to other
materials. Fracture toughness values on the order of
30 MPa√m have been measured on single edge notch
bending (SENB) test specimens.2,31 Strain energy
release rates ranging from 3 to 8 KJ mÀ2 have been
determined on CVI SiC/SiC composites, respectively, with weak or strong interfaces.35 The corresponding values of the J-integral ranges from
11 KJ mÀ2 (weak interfaces) to 29 KJ mÀ2 (strong
interfaces).35 These values are quite high. The aforementioned stress intensity factors are maintained up

to at least 1400  C.2
2.12.6.7 Fatigue and High-Temperature
Behavior
During cyclic fatigue at room temperature, matrix
damage is determined by the maximum stress. It is
created during the first cycles. Fatigue resistance is
governed by the damage of fibers and the fiber–
matrix bonds. Two different fatigue behaviors have
been observed: after 1000 cycles, either the elastic
modulus remains constant and the specimen is running out, or it decreases until the specimen ultimately
fails.48 The modulus degradation reflects either wear
at cracked interfaces48 or growth of interface cracks.
Under stresses smaller than 100 MPa, ultimate failures are generally not observed after 106 cycles under
tension–tension fatigue.
At high temperatures, additional phenomena activated by environment (oxidation, creep, or slow crack
growth) may operate and cause the extension of initial
stress-induced matrix damage and interface cracks as
well as the weakening of fibers by degradation or
reloading. Reloading of fibers involves changes in
load sharing. If the strength of tows is exceeded by
the applied stresses according to the mechanism

described previously, ultimate failure occurs. The rupture of tows dictates the lifetime.
The matrix cracks created upon loading become
the pathways for the ingress of oxygen into the material. The PyC interphase is consumed, which causes
fiber reloading. Creep of the SiC matrix (at very high
temperatures, above 1200  C) makes the stresses on
the fiber to increase, which enhances matrix creep
and further fiber reloading. Creep of fibers (at temperatures above 1200  C) causes matrix reloading and
possible matrix and interface cracking or crack propagation, leading to fiber reloading by decrease of load

carried by the matrix.
Finally, the slow crack growth in fibers (at temperatures below 1100  C) is activated by oxidation
of carbon grain boundaries, leading to delayed
failure.49,50 This phenomenon, which was observed at
intermediate temperatures between 500 and 900  C,
was first referred to as the ‘pest phenomenon’ by a few
authors. The SiC/SiC were claimed to be susceptible
to degradation by oxidation embrittlement.
In order to protect the PyC interphase against
oxidation, multiple coating concepts have been explored
and multilayered interphases and matrices have been
developed.21 Such multilayered matrices contain
phases that produce sealants at high temperatures,
causing healing of the cracks and preventing oxygen
from reaching the cracks and the interphases.2,22,51
Lifetime is also improved with oxidation-resistant
interphases such as BN or multilayers.52
2.12.6.8

Thermal Shock

CVI SiC/SiC composites have been tested under
thermal shock with excellent results.2,53 CVI SiC/
SiC generally had good strength retention after thermal shock cycles involving heating up to the desired
temperature and then cooling down in water at 20  C.
2.12.6.9

Creep Behavior

CVI SiC/SiC and CVI SiC/Si–B–C composites

exhibit primary creep only, even during long tests
(Figures 11 and 12).54
Creep of CMCs involves local stress transfers
depending on the respective creep rates of the fiber
and the matrix. Such stress transfers may lead to
fiber failures or matrix cracking and debonding, and
sliding at the interfaces. When the matrix is elastic
and creep-resistant, fiber creep induces stress transfers from the fibers onto the matrix, which may cause
matrix cracking. This creep-induced matrix damage
has been observed on CVI SiC/SiC composites.55–57


Properties and Characteristics of SiC and SiC/SiC Composites

60 MPa

de/dt (% s–1)

10-4

100 MPa
150 MPa

-0.8

450 MPa (fiber)

10-5

10-6

0.1

1

10

100

Time (h)

Figure 11 Creep rate curves for a damage strain
e0 = 0.8% and for various applied constant stresses for
the SiC/Si–B–C composite, and under 450 MPa at 1200  C
in argon for a Nicalon NL 202 fiber.

-4

10

de/dt (% s–1)

0.22%

-0.8

-5

10

0.14%


-6

10

0.1

1

10
Time (h)

Figure 12 Creep rate curves for the SiC/SiC composite
under a constant stress of 150 MPa (e0 ¼ 0.14% and
e0 ¼ 0.22%) at 1200  C.

In CVI SiC/SiC composites, the SiC matrix is far
more creep-resistant than the SiC fibers, which creep
at 1100  C.55,58,59
The creep behavior of CVI SiC/SiC composites
with a multilayered matrix (SiC/Si–B–C) is caused
by the creep of the Nicalon SiC fibers, whatever
the extent of initial damage created upon loading
(Figure 11). The Si–B–C matrix is less creepresistant and stiffer than the SiC matrix.

properties of CVD SiC and the benefits of reinforcement by SiC-based continuous fibers. In particular,
CVI SiC/SiC exhibits damage tolerance, limited sensitivity to flaws and notch, high load-carrying capacity, and improved reliability. As opposed to earlier
generations of SiC/SiC that were reinforced with
Nicalon fibers, the database for SiC/SiC reinforced
with the advanced near-stoichiometric fibers is

incomplete. The properties of Nicalon-reinforced
SiC/SiC should provide a useful baseline. The main
features of the mechanical behavior of SiC/SiC
composites have been described, and the relationships
between the microstructure and properties have
been discussed. They have been established on CVI
SiC/SiC composites reinforced with Nicalon fibers.
They should be reproduced on those CVI SiC/
SiC composites reinforced with near-stoichiometric
Hi-Nicalon S fibers. The inherent surface roughness
of Tyranno-SA3 fibers is an issue.
Precise knowledge of the mechanisms that govern
the mechanical behavior is useful for proper use or
design with SiC/SiC. Damage and ultimate fracture
of CVI SiC/SiC involve load transfer from matrix to
fibers at various length scales defined by the 2D woven
structure and the tow microstructure. At high temperatures, additional load transfers are driven by the local
stress relaxation induced by temperature and/or environment. The ultimate fracture and delayed failure
are dictated by the tows. Scale effects and scatter in
strength data are limited when compared to monolithic
ceramics. Fiber–matrix interfaces and interphases
play a significant role in damage tolerance and loadcarrying capacity. Interfaces resistant to crack extension
are beneficial to composite performances. However,
this scheme is invalid when the fiber surface is rough.
Properties of CVI SiC/SiC can be tailored via engineering of the interfaces and the use of advanced fibers.

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