Comprehensive nuclear materials 2 12 properties and characteristics of sic and sic composites

Comprehensive nuclear materials 2 12 properties and characteristics of sic and sic composites Comprehensive nuclear materials 2 12 properties and characteristics of sic and sic composites Comprehensive nuclear materials 2 12 properties and characteristics of sic and sic composites Comprehensive nuclear materials 2 12 properties and characteristics of sic and sic composites Comprehensive nuclear materials 2 12 properties and characteristics of sic and sic composites Comprehensive nuclear materials 2 12 properties and characteristics of sic and sic composites

J Lamon

CNRS/National Institute of Applied Science, Villeurbanne, France

ß 2012 Elsevier Ltd All rights reserved.

2.12.2 b-SiC Properties23

325

2.12.6.5 Interface Properties: Influence on the Mechanical Behavior 334

Abbreviations

C/C Carbon matrix composite reinforced by

carbon fibers

C/SiC SiC matrix composite reinforced by

carbon fibers

CMC Ceramic matrix composite

CVD Chemical vapor deposition

CVI Chemical vapor infiltration

LPS Liquid phase sintering

MI Melt infiltration NITE Nanopowder infiltration and transient

eutectic-phase PIP polymer impregnation and pyrolysis

RS Reaction sintering SENB Single edge notch bending

323

SEP Socie´te´ Europe´enne de Propulsion

SiC/SiC SiC matrix composite reinforced by SiC

fibers

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

How-ever, this implies a few inadequate characteristics for

structural applications, such as low fracture

tough-ness, high sensitivity to the presence of

microstruc-tural 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

end-use applications through the sound selection and

arrangement of the constituents Ceramic matrix

com-posites (CMCs) reinforced with continuous ceramic

or carbon fibers are of interest in thermostructural

applications.1–4They are lightweight and damage

tol-erant 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

prop-erties.5 The quality of the material obtained by the

polymer impregnation and pyrolysis (PIP) method is

insufficient A novel processing technique

(nanopow-der infiltration and transient eutectic-phase

proces-sing, NITE) was claimed to achieve good material

quality.5–7

The SiC/SiC composites prepared using the

CVI method and reinforced with the latest

near-stoichiometric 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

stoichiome-try 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 theb-phase properties suggest that CVI SiC/SiC composites have the potential for excellent radiation stability.3 CVI SiC/SiC is also considered for ap-plications 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 reinfor-cing fibers is not affected during the manufacture of the composite; (2) the nature of the deposited mate-rial 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 Propul-sion), Amercorm, Refractory Composites, and others began to develop equipment and processes for produc-ing CVI components for aerospace, defense, and other applications The development of CVI SiC/SiC com-posites has been inspired by the poor oxidation resis-tance 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 compo-nent performance in engine and flight tests

The mechanical properties of SiC/SiC compo-sites depend on the fiber–matrix interface Pyrocar-bon (PyC) has proved to be an efficient interphase

to control fiber–matrix interactions and composite mechanical behavior.20But PyC is sensitive to oxida-tion 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 inter-phases and matrices have been developed.3,21 Multi-layered matrices contain phases that produce sealants

at high temperatures, preventing oxygen from reach-ing the interphase.22This 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

added to the matrix to scavenge oxygen that might

ingress into the matrix (enhanced CVI SiC/SiC)

The mechanical behavior of CMCs displays

sev-eral typical features that differentiate them from the

other composites (such as polymer matrix

compo-sites, metal matrix compocompo-sites, etc.) and from

homo-geneous (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

dis-cussed with respect to microstructure, on the basis of

the large amount of work done on CVI SiC/SiC

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

irradia-tion 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.23Here 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

gas-phase reaction methods for the synthesis of highly

crystalline, stoichiometric, high-purityb-SiC

2.12.2.1 Mechanical Properties

2.12.2.1.1 Elastic modulus23

Generally, a dense and high-purity SiC material, for

example, CVD SiC, exhibits the highest elastic

modu-lus; 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 (V ):

E¼ E0expðCVpÞ ½1

E0¼ 460 GPa for CVD SiC (polycrystalline, high-purity, very dense, and pore-free SiC material) and

C¼ 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:

E¼ E0 BTexpðT0=TÞ ½2 with E0¼ 460 GPa, B ¼ 0.04 GPa K1, and T0¼ 962 K

2.12.2.1.2 Poisson’s ratio23 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 bend-ing technique This value was also derived from the elastic modulus and Poisson’s ratio (n), using the con-ventional formula for isotropic solids: G¼ 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 afore-mentioned exponential function of porosity for elastic modulus can be extended to the hardness evaluation:

HV¼ 27:7 expð5:4VpÞ ½3 where HVis the Vicker hardness

Currently, there is no high-temperature data reported for high-purity CVD SiC

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 tough-ness of CVD SiC increases slightly at elevated tem-peratures It does not exceed 6 MPa√m

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

conse-quence 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

satis-factorily describes the statistical distribution of

fail-ure strengths:

P ¼ 1  exp 

ð ðs=s0ÞmdV=V0

½4

where P is the probability of failure,s is the stress,

s0is the scale factor, m is the Weibull modulus, V is

the volume of specimen, and V0is a reference volume

(1 m3is generally used); m reflects the scatter in data,

ands0is related to the mean value of the strength

The strength data for a given geometry and stress

state can be determined usingeqn [4] However, m,

s0, and V0must be available It is important to note

that the estimate of s0 depends on V0.24 It will be

substantially different if V0¼ 1 m3

or 1 mm3 This dependence is ignored in most publications, even in

the work by Snead and coworkers23in which a number

of s0values are reported When V0 is not given, the

estimate ofs0is 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.23span

a large range, from 2 to 12 The following values were

measured using tensile tests on CVI SiC/SiC

mini-composites: m¼ 6.1, s0¼ 10.5 MPa (V0¼ 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

polycrystallineb-SiC) Creep in SiC is highly

depen-dent 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:

ec ¼ Apðs=GÞnðt=tÞp ½5 where Ap, p, and t are creep parameters, and t is the time elapsed n¼ 1.63, Ap¼ 29, p ¼ 0.081, and

t ¼ 0.0095 s for the temperature of 1923 K These para-meters 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 mate-rials have been measured only above1673 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 depo-sition direction The strain rate is given by a power-law creep equation:

de=dt ¼ Asðs=GÞnexpðQ =kbTÞ ½6 where As¼ 2.0  103

, n¼ 2.3, Q ¼ 174 kJ mol1 (acti-vation energy),s is the applied stress, G is the shear modulus, and kbis 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 practi-cally the same conductivity as the single-crystal material It is worth noting that the impurity content

of the very high thermal conductivity CVD SiC mate-rials is negligibly small, and this material has near theoretical density (3.21 g cm3) The curve-fitting

to the single-crystal SiC data above 300 K yields

an upper limit of the thermal conductivity of SiC (in W m1K1):

Kp¼ ð0:0003 þ 1:05  105TÞ1 ½7

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 system-atic difference can be distinguished between the struc-tural types The specific heat, Cp(in J kg1K), over the temperature range 200–2400 K can be approximately expressed as

Cp ¼ 925:65 þ 0:3772T  7:9259  105T2

 3:1946  107=T2 ½8

The specific heat of SiC at room temperature is taken

as 671 47 J kg1K

2.12.2.2.3 Thermal expansion

The coefficient of thermal expansion forb-SiC has

been reported over a wide temperature range The

average value in the interval from room temperature

to 1700 K isa ¼ 4.4  106K1

At higher temperatures (T> 1273 K), a ¼

5 106K1

At lower temperatures (550< T < 1273 K),

a ¼ 2.08 þ 4.51  103T

It is worth addressing the processing method first

because this information is useful for a better

under-standing of the structure of SiC/SiC The

manufac-ture 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

2.12.3.1 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

appro-priate 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

tempera-tures.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 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

a multilayer ((PyC/SiC)nor (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 BCl3and NH3for boron nitride Multilayered inter-phases may be deposited via pulsed CVI

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.13–15The reactions consist of ing a hydrocarbon for deposition of carbon and crack-ing 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 pro-posed 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.29The 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 limita-tions 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 radia-tion 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

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

reinfor-cing fibers and availability of fine nano-SiC powders

A slurry ofb-SiC nanopowders and additives is

infil-trated 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

advan-tages such as flexibility in the shape and size of the

components.7The successful development of NITE

is due to appropriate fiber protection and the

emer-gence of advanced SiC fibers such as Tyranno-SA3

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 Nica-lon fibers of first generation.2,31An average strain-to-failure of 0.3% and a tensile strength of 200 MPa have been reported Higher strengths and strains-to-failure appear inTables 2and3, 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 ten-sile 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 condi-tions Ideally, the strain-to-failure should coincide with that of reinforcing tows, that is, about 0.8%

Table 1 Mechanical and thermophysical properties of 2D SiC/SiC composites reinforced with 0/90 balanced Nicalon™ fabrics

Poisson’s ratio

In-plane coefficient of thermal expansion (106K1) 3 3

Thru-the-thickness coefficient of thermal expansion (106K1) 1.7 3.4

Thru-the-thickness thermal conductivity (W m1K1) 9.5 5.7

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;

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

mono-lithic 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 fiber-reinforced composites Furthermore, besides the resistance to crack propagation, damage tolerance is

an important property for CMCs It cannot be char-acterized 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 1shows that CVI SiC/SiC retains its prop-erties at high temperatures These propprop-erties 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

Table 2 Mechanical properties of a CVI SiC/Si–B–C

composite with a self healing matrix and a multilayer

rein-forcement of Hi-Nicalon™ fibers, and 2D CVI-enhanced

SiC/SiC composite reinforced with 0/90 five harness satin

fabrics of Hi-Nicalon™fibers

Room temperature

1200
C

CVI SiC/Si–B–C

Fiber type Hi-Nicalon™

fibers

Hi-Nicalon™ fibers Reinforcement Plain weave Plain weave

Porosity (%) 13

Tensile strength

(MPa)

315

Strain-to-failure (%)

0.5 Young’s modulus

(GPa)

220 Interlaminar shear

Flexural strength

(MPa)

2D CVI-enhanced SiC/SiC composite

Fiber type Hi-Nicalon™ Hi-Nicalon™

Reinforcement 0/90 five harness

satin

0/90 five harness satin

Tensile strength

(MPa)

Strain-to-failure

(%)

Young’s modulus

(GPa)

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.

Table 3 Room-temperature properties of 2D melt infil-trated CVI SiC/SiC and 2D CVI SiC/SiC composites rein-forced with Hi-Nicalon type S fibers

2D melt infiltrated CVI SiC/SiC

Tensile strength (MPa) 341–412 Strain-to-failure (%) 0.60 Young’s modulus (GPa) 232–262 2D CVI SiC/SiC

Tensile strength (MPa) 305 Strain-to-failure (%) 0.60 Young’s modulus (GPa) 214

45
off-axis tensile strength (MPa) 167

45
off-axis strain-to-failure (%) 0.66

Source: Morscher, G.; Pujar, V Int J Appl Ceram Technol.

2009, 6, 151–163.

Tables 3 and 4 show that the ultimate strength

and Young’s modulus tend to decrease under off-axis

tensile conditions.36It 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

princi-pal stresses coincide with those of the fiber tows

The matrix of NITE-SiC/SiC comprises

polycrys-talline SiC and a small amount of isolated oxides

The microstructure is highly crystalline and highly

dense Table 5 lists the typical available

pro-perties of NITE-SiC/SiC.7 Thermal conductivity

(30 W m1K1) is quite high when compared to

CVI SiC/SiC (below 15 W m1K1) reinforced with

either Nicalon (Table 1) or Hi-Nicalon fibers.37The

high proportional stress limit is claimed to be an

inter-esting feature.7However, it is worth pointing out that

it reflects a high load-carrying capacity By contrast,

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 propor-tional stress limit after exposure at temperatures up to

1300
C.6

SiC/SiC

2.12.6.1 Tensile Stress–Strain Behavior Figures 1 and 2 summarize the typical stress–strain behavior of 2D CVI SiC/SiC composites The behav-ior is initially linear under strains below 0.03% Then,

Table 5 Room-temperature properties of NITE-SiC

composites

Property

Tyranno-SA3

Tyranno-SA3

Proportional limit (MPa) 358 148

Tensile strength (MPa) 408 167

Strain-to-failure (%) 0.13 0.08

Thermal conductivity

(W m1K1)

32

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

Wiley-VCH: Weinheim, Germany, 2008; Chapter 15, pp 353–384,

reproduced with permission.

0 0 100 200 300 400

Longitudinal tensile strain (%)

(a)

(b)

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.

0.2 0.1 0 50 100 150 200

0.3 Strain (%)

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

Table 4 Off-axis properties of a first generation of 2D

CVI SiC/SiC reinforced with Nicalon fibers

Saturation stress (MPa) 150 145 145

Tensile strength (MPa) 190 170 170

Source: Aubard, X.; Lamon, J.; Allix, O J Am Ceram Soc 1994,

77, 2118–2126.

the nonlinear deformations result essentially from

transverse cracking in the matrix (the cracks are

per-pendicular 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:

Ec¼ EmVmþ EfVf ½9

where Vmis the volume fraction of matrix and Vfis

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 Ecis significant Then, as it decreases

when the matrix cracks, the behavior becomes

con-trolled by the tows The 2D SiC/SiC composites

exhibit an elastic damageable behavior (Figure 3)

This means that the response of the damaged

mate-rial is elastic as indicated by the linear portion of the

curves on reloading.Figure 4shows the dependence

of the elastic modulus on damage

2.12.6.2 Damage Mechanisms The basic damage phenomena in unidirectional com-posites under on-axis tensile loads involve multiple microcracks or cracks that form in the matrix perpen-dicular 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.39Furthermore, 2D CVI SiC/SiC is a hetero-geneous 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

Strain (%)

0.6

0

0

50

100

150

200

250

300

350

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 V and V , respectively.

0 0 0.2 0.4 0.6 0.8 1.0 1.2

0.2

F

A

D G 0.4

Strain (%)

E0

EfVf

1.0

Figure 4 Relative elastic modulus versus applied strain during tensile tests on various 2D woven SiC/SiC composites reinforced with treated fibers: (A) Nicalon/(PyC 20 /SiC 50 ) 10 / SiC, (D) Nicalon/PyC 100 /SiC, (F) Hi-Nicalon/PyC 100 /SiC, (G) Hi-Nicalon/(PyC 20 /SiC 50 ) 10 /SiC.

Macropore

Longitudinal tow

Transversal tow

Layer 0.5 mm

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

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

lon-gitudinal tows (deformations larger than 0.2%)

These microcracks are confined within the

lon-gitudinal tows They do not propagate in the

rest of the composite The matrix in the

longi-tudinal tows experiences a fragmentation

pro-cess 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

condi-tions, 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 dis-cussed 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 longi-tudinal 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):

Emin¼ 1=2EfVf ½10 where Vfis 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

Longitudinal strain = 0.06% Longitudinal strain = 0.2%

Longitudinal strain = 0.6% Longitudinal strain = 0.8%

Longitudinal strain = 0.4%

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