5.

1 Design Objectives and Design Philosophies, Interphases and Interfaces in

Fiber-Reinforced CMCs
Walter Krenkel and Florian Reichert, Lehrstuhl Keramische Werkstoffe, Bayreuth, Germany
r 2018 Elsevier Ltd. All rights reserved.

5.1.1 General Properties of Ceramic Matrix Composites 1

5.1.2 Classification of CMCs 4
5.1.2.1 Oxide and Non-Oxide CMCs 4
5.1.2.2 Manufacture of CMCs 5
5.1.2.3 Weak Interphase and Weak Matrix Composites 6
5.1.2.3.1 Weak interphase composites 6
5.1.2.3.1.1 Concepts of weak interphases 6
5.1.2.3.2 Weak matrix composites 8
5.1.2.3.3 Mechanical behavior of WMCs and WICs 9
5.1.3 Tailoring the Fiber/Matrix Bondings and the Microstructure of C/C-SiC Composites 10
5.1.4 New Design Philosophies and Industrial Applications 13
5.1.5 Summary 16
References 17

5.1.1 General Properties of Ceramic Matrix Composites

Ceramics exhibit extraordinary good properties especially at high temperatures. The high Young’s modulus, strength, corrosion
and oxidation resistance together with the high hardness allow the application in different fields like friction and wear or medical
and chemical engineering. The biggest drawback, which prevents an increasing spread of monolithic ceramics in new applications,
is the brittle failure behavior, which results in a volume dependent strength and a statistical failure rate. Technical ceramics should
therefore not subjected to tensile stresses. With ceramic matrix composites (CMCs), it is possible to overcome this disadvantage.
CMCs are a relatively young class of composite materials in comparison to other established structural materials or monolithic
ceramics. A fiber reinforcement within a ceramic matrix is applied to prevent brittle fracture and subcritical crack growth. The
different mechanical behaviors of a CMC material as compared to a monolithic ceramic are shown exemplarily in Fig. 1.
Monolithic ceramics show a catastrophic failure behavior with a sudden break in the elastic region. The Young’s modulus of
CMCs is lower and the stress–strain curve shows typically a linear and a nonlinear region. This results in a higher elongation at
break, and in combination with a step-wise failure after reaching the ultimate strength, in a quasi-ductile fracture behavior.
CMC materials allow lightweight designs of components for corrosive and/or high temperature environments. Their low
density and high maximum operating temperature make CMCs interesting materials for high temperature lightweight applications
under tensile load, for example, for thermal protection systems of spacecraft, kiln furnitures, gas turbines, or high performance

Fig. 1 Typical stress–strain behavior of a ceramic matrix composite (CMC) material in comparison to a monolithic ceramic material under tensile
loading.

Comprehensive Composite Materials II, Volume 5 doi:10.1016/B978-0-12-803581-8.09986-0 1

2 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs

friction applications. The densities of oxide (oxide fiber composites (OFC)) and non-oxide composites (e.g., SiC/SiC, C/SiC) are
compared with other structural materials in Fig. 2. The values of CMCs are lying between magnesium and aluminum, which are
well known lightweight materials. The low density accompanied by the high temperature stability results in the highest mass
specific strength of all materials beyond 11001C (Fig. 3). In contrast, carbon-fiber reinforced plastics (CFRP) are the lightest
structural materials at room temperature due to the low density of polymers and the high strength of carbon fibers, resulting in
very high mass specific strength values of about 40 km. Ceramic composites show a considerably lower mass specific strength of
about 10–15 km. However, there is no alternative structural material at the highest temperatures, explaining the growing interest
for this class of structural materials.
In addition to the low density, CMCs keep their properties over a broad range of temperature. CMCs are processed at high
temperatures, commonly above 10001C. The mismatch of the coefficients of thermal expansion (CTE) between fiber and matrix
results in residual stresses after the cooldown to room temperature.1–3 However, these inherent stresses are reduced, when the

Fig. 2 Comparison of the density of different materials for structural applications. Reproduced from Krenkel, W., 2011. Technische Keramik. In:
Henning, F., Moeller, E. (Eds.), Handbuch Leichtbau: Methoden, Werkstoffe, Fertigung. München: Hanser, pp. 393–411.

Fig. 3 Mass specific strength vs. service temperature of different structural materials. Reproduced from Krenkel, W., 2011. Technische Keramik.
In: Henning, F., Moeller, E. (Eds.), Handbuch Leichtbau: Methoden, Werkstoffe, Fertigung. München: Hanser, pp. 393–411.
 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs 3

composite is used at temperatures close to the processing temperatures. In combination with the closure of the microcracks in the
matrix, this effect result in an increase of strength at higher service temperatures. Fig. 4 shows this behavior of a C/C-SiC material
manufactured by the liquid silicon infiltration process (LSI). The bending strength at 1200°C is increased by about 18% and the
Young’s modulus by 40% compared to the strength and modulus at room temperature.4
Designing components with CMCs means also designing with a fiber reinforced material and thus a material with anisotropic
properties. The anisotropy results from the fibers and influences mechanical and thermomechanical properties, for example,
stiffness, strength, coefficient of thermal expansion, or thermal conductivity. By way of example, the dependence on the strength
and stiffness of a LSI-derived short fiber reinforced composite with different fiber orientations is given in Table 1. The flexural
strength and Young´s modulus of unidirectional (UD) aligned short fiber reinforced composites are decreasing with an increasing
angle between fiber orientation and tensile stress vector. The calculated mean values fit to the results of in-plane isotropic short fiber
reinforced CMCs, which demonstrates the influence of the fiber orientation on the mechanical properties. Also, the coefficient of
thermal expansion is heavily dependent on the fiber orientation, which is depicted exemplarily in Fig. 5. Two different CMCs, a C/
C-SiC and an Ox/Ox composite, are compared with isotropic materials, monolithic ZrO2 and two metals (steel and superalloy). The
CTEs of the ceramic materials are significantly lower than the values of the metals, which means that joining of different materials,
like CMC with metal or CMC with ceramic, needs to consider this behavior in order to minimize thermally induced stresses.
Additionally, CMCs with different materials of fibers and matrix exhibit different CTEs in dependence on the fiber orientation. This
is shown with the values of C/C-SiC in fiber direction (8) and perpendicular to the fiber direction (>). The reason for that is the

Fig. 4 Strength and Young’s modulus of 2D-C/C-SiC composites at room temperature, 900 and 12001C. Reproduced from Krenkel, W., 2005.
Carbon fibre reinforced silicon carbide composites (C/SiC, C/C-SiC). In: Bansal, N.P. (Ed.), Handbook of Ceramic Composites. Boston, MA:
Springer US, pp. 117–148.

Table 1 Mechanical properties of liquid silicon infiltration (LSI)-derived short fiber reinforced C/C-SiC with different fiber orientations in
comparison with isotropic SF C/C-SiC (fiber length 6 mm, 3-point-bending test according to DIN EN 658–3)

Fiber orientation of unidirectional (UD) composites angle a¼ Mean value of Isotropic

all UD composite (all
0 15 30 45 60 75 90 composites fibers
degree degree degree degree degree degree degree randomly
distributed)

Flexural strength (MPa) 67 62 48 43 30 28 25 44 46

Young’s modulus (GPa) 45 40 30 25 22 19 16 28 25
Strain to failure (%) 0.17 0.16 0.17  0.17 0.16 0.16 0.16 0.17
4 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs

Fig. 5 Coefficients of thermal expansion (CTE) vs. temperature of selected metals and ceramic materials. Reproduced from Krenkel, W., 2011.
Technische Keramik. In: Henning, F., Moeller, E. (Eds.), Handbuch Leichtbau: Methoden, Werkstoffe, Fertigung. München: Hanser, pp. 393–411.

very low CTE of carbon fibers in comparison to the SiC matrix. The anisotropy in the CTE of the Ox/Ox composite with mullite
fibers (Nextel 720) and a mullite matrix is less pronounced, because fibers as well as matrix consist of the same material.
The design objective of CMCs for structural applications is to achieve a ceramic composite material with a non-brittle failure
behavior. The interphase between the fibers and the matrix plays the key role to achieve a suitable load transfer by weak bondings.
By definition, a composite material with an interphase has two interfaces, one between matrix and interphase and a second one
between interphase and fiber surface. The interphase itself is mostly a thin coating on the fibers, which has to adopt various tasks:

• Protection against corrosive attacks of the fiber, for example, due to combustion atmospheres in engines or diffusing atoms in
high temperature environments.
• Mechanical contact between fiber and matrix, load transfer by shear stresses.
• Crack deflection around the fiber.
A suitable fiber/matrix interface resp. interphase result in a crack bridging by the fibers and a crack branching at the crack tip,
whereby energy is dissipated and the crack tip is blunted. The crack extends around the fibers, resulting in a debonding of the fibers
and in the generation of frictional forces by the fiber pull-out, leading to a quasi-ductile nonlinear material behavior.

5.1.2 Classification of CMCs

CMCs can be classified by several aspects, which are the type of matrix and fibers (oxide or non-oxide CMCs), the processing
method or the fiber/matrix stiffness ratio, differentiating CMCs into weak matrix composites (WMC) and weak interphase
composites (WIC).

5.1.2.1 Oxide and Non-Oxide CMCs

The classification of CMCs according to their main constituents fibers and matrix results in two classes:

• Non-oxide CMC (e.g., C/C, C/C-SiC, C/SiC, SiC/SiC).

• Oxide CMCs (abbreviated Ox/Ox, OCMC or OFC, e.g., Al2O3/Al2O3).

Non-oxide CMCs show a very low creep rate at high temperatures, thus are suitable materials for highly stressed components at
temperatures even exceeding 11001C. Composites with a carbon matrix or fibers suffer from the poor oxidation resistance of
carbon beyond 4001C, thus are in need of oxygen free atmospheres or they are only applicable in limited life structures. The latter
approach is applied, for example, for space or brake applications. SiC/SiC composites are of interest due to their low creep rate and
their excellent mechanical properties at high temperatures, resulting in potential candidates for applications in nuclear power
plants5 and advanced gas turbines.6 In case of severe atmospheres, including oxygen and water contents, SiC/SiC composites are
in need of environmental barrier coatings (EBCs) to prevent an active oxidation, which leads to the formation of gaseous
 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs 5

reaction products, like SiO.7–9 The occurrence of water vapor, for example, in combustion environments, leads to the formation of
Si(OH)4, a gaseous reaction product, which is removed from the CMC surface, leading to a fast corrosion.10 To prevent or to
miminize corrosive effects, EBCs can be applied, for example, made of rare earth silicates.11–13
Ox/Ox composites are inherently stable in air and are oxidation resistant materials up to 11001C. Current limits are the grain
growth in the oxide fibers at high temperatures and the further sintering of the matrix, both resulting in a strength decrease and
embrittlement of the composite.14–17 Additional water vapor in the environment causes a faster failure of Ox/Ox composites due
to the hydroxylation of silicon containing composites and/or environmental assisted subcritical crack growth.18–20 In comparison
to non-oxide CMCs, they show a lower mass specific strength level because of their higher density (about 3 g/cm³) and their lower
mechanical properties.

5.1.2.2 Manufacture of CMCs

An overview of different processing methods for CMCs is given in Fig. 6, which shows different approaches to infiltrate a fiber
preform. Possible preforms comprise continuous fiber tows, short fibers, 2D-fabrics, nonwovens or 3D-fiber architectures. The
infiltration of the fiber preform can be made with a fluid, either by a gas phase (chemical vapor infiltration (CVI)) or by a liquid
phase (slurry, sol, precursor, or molten metals). The CVI process shows a high variability and reproducibility and results in materials
with the highest thermomechanical properties. However, it is a long-lasting process of several days or weeks. The infiltration with a
liquid phase can be distinguished between processes with an in-situ reaction during/after infiltration or without any reaction between
fiber preform and infiltration medium. The slurry infiltration ((SI)-process) features no reaction and is mainly used for Ox/Ox
composites. Nano-scaled powders and multimodal particle size distributions are applied to achieve a good fiber bundle infiltration
and a tailored matrix porosity.21 The infiltration of ceramic precursors like polysilazanes, polycarbosilizanes, or polysiloxanes are
used in the polymer infiltration and pyrolysis process (PIP). Well-established polymer techniques can be utilized to form the green
body and the high diversity of precursors result in a high variability of ceramic matrices.22,23 The disadvantage of this process is the
high porosity of the matrix after pyrolysis, which is the result of the high volumetric shrinkage of the precursor due to the increase of
density from the polymer to the ceramic state accompanied by a certain weight loss (20–30%).24 Several infiltration cycles each
followed by a pyrolysis step are necessary to densify the composite and to achieve high mechanical properties.25
The LSI process is an example for a liquid phase infiltration with in situ-reaction. The main advantages of this process are the
fast processing cycles without any re-infiltration steps, the robustness and the high reproducibility. The matrix consists of silicon
carbide, residual carbon, and free silicon, the latter limits the application temperature to about 13501C. The LSI process is a well-
established method for the series production of C/C-SiC composites, for example, for automotive brakes26 or elevator emergency
brake parts.27 In most cases, silicon is used for the infiltration, resulting in silicon carbide matrices. Instead of silicon, other reactive
melts or alloys can be used, resulting in matrices comprising ultra-high-temperature ceramics (UHTCs)28 or MAX-phases.29

Fig. 6 Processing routes for ceramic matrix composite (CMC) manufacturing.

6 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs

An overview of typical mechanical and thermal properties of SiC-containing CMCs with 2D-carbon fiber reinforcements is given
in Table 2. The CVI process results in composites with the highest strength and strain to failure. All composites show a certain amount
of matrix porosity, with the lowest porosities (less than 5%) for the LSI process. Because of the non-isotropic fiber distribution, the
properties show a high degree of anisotropy. In comparison to monolithic ceramics, the bending strengths of the CMCs are lower.
However, CMCs can be exposed to tensile stresses comparable to ductile materials like metals or other composite materials.

5.1.2.3 Weak Interphase and Weak Matrix Composites

To achieve a damage-tolerant behavior of CMCs, two main approaches can be followed:

• Weak interphase composites (WIC).

• Weak matrix composites (WMC).

In both approaches, a weak component is introduced to enable a selective matrix cracking. There are two theoretical models in
literature dealing with the crack deflection at the interface of two brittle solids. The first one is derived from simulations of Cook
and Gordon.30 It describes a debonding at the interface ahead of a crack and a coalescence of this nucleated debond crack and the
primary crack, which was observed for several material combinations.31–36 According to this model, debonding is possible, when a
crack is generated between material 1 and 2 and material 2 does not fail (Fig. 7(b)). Material 1 can be the matrix or the interphase,
material 2 can be the interphase or the fiber. Thus, a failure prediction can be made, which uses the ratio of the Young’s moduli of
the two materials E2/E1, the critical tensile stress of the undamaged material 2 ðsc2 Þ and the tensile strength of the interface
perpendicular to the original crack ðsci Þ. The Young’s moduli and the tensile strength of the undamaged material 2 ðsc2 Þ, for
example, the fiber tensile strength, can be determined easily and thus an estimation of ðsci Þ for various fiber/matrix and
fiber/interphase systems can be made (Fig. 8).
A different approach follows the theoretical model firstly developed by He and Hutchinson,37 which was subsequently
improved and complemented.38–40 The model describes a crack at the fiber/matrix interface and the three possibilities for crack
progress comprising one-side deflection, two-side deflection or penetration of the fiber without deflection (Fig. 7(a)).
The He–Hutchinson diagram (Fig. 9) displays the crack behavior, with the ratio of the debonding energy Gi and the fiber
fracture energy Gf on the y-axis and the ratio of the Young’s moduli of fiber and matrix plotted on the x-axis. The curve progression
is dependent on the angle between crack and fiber/matrix interface and also on residual stresses and clamping forces, which can
occur due to the mismatch of the CTE between fiber and matrix.38 The diagram is representative for a material without clamping
forces (no mismatch of the CTEs) and an angle of 90 degree between the fiber/matrix interface and the crack. At similar moduli of
fiber (Ef) and matrix (Em), the x-axis value converges to 0. Thus, a very low debond energy (Gi) compared to the fracture energy of
the fiber (Gf) is needed to achieve crack deflection. A fracture energy ratio of 0.25 or lower is necessary to achieve a debonding,
when fiber and matrix possess the same Young’s modulus. If the Young’s modulus of the fiber is significantly higher than the
modulus of elasticity of the matrix, debonding can only be achieved at higher fracture energy ratios. Both material concepts are
used for CMCs and are discussed in more detail in the following.

5.1.2.3.1 Weak interphase composites

The He–Hutchinson diagram shows that composites made of constituents with similar Young’s moduli (e.g., CVI-derived SiC/SiC)
are in need of a sufficiently low fracture energy of the interface compared to the fiber in order to achieve a crack deflection at the
fiber/matrix interface.

5.1.2.3.1.1 Concepts of weak interphases

An overview of different interphase concepts and fiber/matrix transitions in WIC shows Fig. 10. Fig. 10(a) illustrates a monolayer
coating with a debonding at the interface. The crack extends around the reinforcing fibers and allows debonding and crack bridging
by the fibers. Examples are coatings made of amorphous pyrolytic carbon (PyC) or BN.41–46 Other designs provide interphases
with a layered crystal structure that allow an easy cleaving (Fig. 10(b)), for example, hexagonal BN, which is formed in-situ47 or is
deposited by the CVI method above B15001C.48 Other designs provide a fiber coating made of layered perovskites,49 hex-
aluminates,50 phosphates51–53 or anisotropic PyC. The last example is exemplarily shown in Fig. 11(a), displaying a SiC/SiC
composite with a single layer of anisotropic PyC. The cleavage planes of these coatings result in a crack deflection around the fibers.
The same effect can be achieved with multilayer interphases (Fig. 10(c)), which are usually deposited by the CVI process.
An example is given in Fig. 11(b), which illustrates a CVI-derived SiC/SiC composite with a PyC and SiC multilayer interphase. In
general, the multilayer interphase has the advantage, that the different fiber coatings allow a separation of functions. Thus, the
crack deflection by weak boundary layers can be combined with layers, which serve as diffusion barriers and a protection against
corrosion, for example, in oxidative environments.54 Layers can consist of PyC, boron nitride or SiC.55 Especially boron and silicon
containing interphases show self-healing effects in oxidizing atmosphere due to the formation of glass phases,56,57 thus providing
an enhanced corrosion protection of the fibers.
A weak interphase can also be achieved by an adjusted porosity of the interphase (Fig. 10(d)), which obviously reduces the
interphase strength and modulus. It can be manufactured by pore-forming fiber coatings, for example, based on ceramic pre-
cursors, sol–gel systems or CVI processes.58 Examples can be found with oxide-based porous coatings, for example, made of
zirconia,59 rare earth aluminates60 or zircon–carbon.61 Additionally, fugitive coatings are an alternative concept. Fibers are coated
 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs
Table 2 Typical thermal and mechanical properties of SiC-containing ceramic matrix composites (CMCs) with 2D-carbon fiber reinforcements manufactured by different processes

Property Unit Gasphase infiltration processes Liquid infiltration processes

Chemical vapor infiltration (CVI) (isothermal) CVI (p, T-gradient) Polymer infiltration and pyrolysis (PIP) Liquid silicon infiltration (LSI)
C/SiC C/SiC C/SiC C/SiC C/C-SiC

Tensile strength MPa 350 300–320 265 240–270 80–190

Strain to failure % 0.9 0.6–0.9 0.5 0.8–1.1 0.15–0.35
Young’s modulus GPa 90–100 90–100 90 60–80 50–70
Compression strength MPa 580–700 450–550 590 430–450 210–320
Flexural strength MPa 500–700 450–500 500 330–370 160–300
Interlam. shear strength MPa 35 45–48 20 35–37 28–33
Porosity % 10 10–15 8 10–15 2–5
Fiber content Vol.% 45 42–48 52 40–45 55–65
Density g/cm3 2.1 2.1–2.2 1.8 1.9–2.0 1.9–2.0
Coefficient of thermal expansion 10–6K1
ǁ 3a 3 2–5d 3  1–2.5b
> 5a 5 4.06d 4 2.5–7b
Thermal conductivity W/(m*K)
ǁ 14.3–20.6a 14 14b – 17.0–22.6c
> 6.5–5.9a 7 5.3–5.5b – 7.5–10.3c
Specific heat J/(kg*K) 620–1400 – 600 – 690–1550
a
¼ RT–10001C.
b
¼ RT–15001C.
c
¼ 200–16501C.
d
¼ RT–15001C.
ǁ, >, parallel resp. transverse to fiber direction.
Source: Reproduced from Krenkel, W., 2011. Technische Keramik. In: Henning, F., Moeller, E. (Eds.), Handbuch Leichtbau: Methoden, Werkstoffe, Fertigung. München: Hanser, pp. 393–411.

7
8 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs

Fig. 7 Illustration of (a) the He–Hutchinson model of a crack at the fiber/matrix interface with the two possibilities of fiber penetration or
deflection around the fiber and (b) the Cook–Gordon approach of a nucleated crack at the fiber/matrix interface. Reproduced from Pompidou, S.,
Lamon, J., 2007. Analysis of crack deviation in ceramic matrix composites and multilayers based on the Cook and Gordon mechanism.
Composites Science and Technology 67 (10), 2052–2060.

Fig. 8 Regions of non-debonding and debonding according to the model of Cook and Gordon for various fiber/interphase, interphase/matrix, and
fiber/matrix systems. Reproduced from Pompidou, S., Lamon, J., 2007. Analysis of crack deviation in ceramic matrix composites and multilayers
based on the Cook and Gordon mechanism. Composites Science and Technology 67 (10), 2052–2060.

with a material, which can be removed after matrix formation, for example, carbon or molybdenum.62–64 This leaves a gap
between fibers and matrix. A crack can be deflected easily, but the stiffness and the thermal conductivity of the composite is
reduced significantly.

5.1.2.3.2 Weak matrix composites

CMCs having a matrix with a much lower modulus of elasticity compared to the fibers, the ratio (Ef  Em)/(Ef þ Em) of the
He–Hutchinson model approaches the value 1 (Fig. 9). The interface then becomes less important because even with
strong bondings between of fiber and matrix, a crack deflection is possible due to the weakness of the matrix. Examples of
this so-called WMC are materials having a porous or micro-cracked matrix, such as PIP-derived C/C or SiC/SiC materials or
 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs 9

Fig. 9 He–Hutchinson diagram with the relative fracture energy of the fiber and the interface vs. the stiffness ratio of fiber and matrix.
Reproduced from He, M.-Y., Hutchinson, J.W., 1989. Crack deflection at an interface between dissimilar elastic materials. International Journal of
Solids and Structures 25 (9), 1053–1067.

Fig. 10 Different types of fiber/matrix interphases: (a) monolayer, (b) interphase with a layered crystal structure, (c) multilayer, (d) porous
interphase. Reproduced from Naslain, R.R., 1998. The design of the fibre-matrix interfacial zone in ceramic matrix composites. Composites Part A:
Applied Science and Manufacturing 29 (9–10), 1145–1155.

SI-derived Ox/Ox composites. Fig. 12 shows exemplarily the microstructure of an Al2O3/Al2O3 composite with a nanoporous
matrix, manufactured via SI. There is a sharp interface between fibers and matrix and no interphase is necessary to allow a
damage-tolerant fracture behavior. The material fails with distinct delaminations and a subsequent fiber pull-out as shown in
Fig. 13, which displays the stress–strain curves of a 3-point-bending test of WMCs made of Nextel 610 fabrics with an Al2O3–ZrO2
matrix.

5.1.2.3.3 Mechanical behavior of WMCs and WICs

WMC and WIC materials differ in their mechanical behavior due to their different matrix microstructure. For example, an off-axis
load of a 2D-reinforced WIC results in no change of the elastic modulus. Even if the fiber direction is not parallel to the load
direction (e.g., for a 45 degree loading), the modulus of elasticity of the composite material is not reduced since fiber and matrix
have similar moduli of elasticity. The dependence on strength and Young’s moduli from the angle between fiber axis and loading
direction is strongly pronounced in WMCs. The tensile strength and Young’s modulus of samples with an angle of 45 degree
between fiber axis and load direction is greatly reduced, since the strength of the reinforcing fibers cannot be exploited and the
weak matrix has to carry the load. Fig. 14 demonstrates this different mechanical behaviors with two examples. It displays the
stress–strain curves of a tensile test of a CVI-derived SiC/SiC material representing a WIC and a PIP-derived C/C composite
representing a WMC.
As a result, a distinction between WIC and WMC materials is possible by comparison of the Young’s modulus ratio in
0/90 degree and þ 45/  45 degree direction. Different CMCs are summarized in Fig. 15. C/C composites are commonly man-
ufactured via the polymer impregnation route. As a result, the carbon matrix is microporous and micro-cracked, resulting in a
typical WMC behavior. In contrast, SiC fiber composites with strong fiber/matrix bondings (e.g., CVI derived SiC matrix) are
produced with a weak interphase to ensure a damage tolerant failure behavior, resulting in a typical WIC behavior. Ox/Ox
composites belong either to WMC or WIC composites, depending on the manufacturing process.
10 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs

Fig. 11 Representatives of weak interphase composites (WIC): SiC/SiC composites with (a) a single layer interphase made of anisotropic pyrolytic
carbon (PyC) (reproduced from Droillard, C., 1993. Processing and characterization of SiC-matrix composites with multilayered C/SiC interphase,
PhD thesis, Bordeaux) and (b) a multilayered interphase made of PyC and SiC with 10 layers (reproduced from Bertrand, S., 1998. Improvement
of the durability of SiC/SiC composites with multilayered (PyC-SiC)n or (BN-SiC)n Interphases. PhD thesis, Bordeaux).

Fig. 12 Typical micrograph of a weak matrix composites (WMC) material made of alumina fibers and a nanoporous alumina matrix.

An example of a mixture of strong and weak fiber/matrix bondings in one material are C/C-SiC composites, derived from
the LSI process. Si- and SiC-rich matrix regions form a stiff and strong matrix and interface due to the reactive bonding to the
surrounding carbon. However, the weaker bondings between carbon matrix and carbon fibers (C/C-regions) can be assigned to
the WMC classification.

5.1.3 Tailoring the Fiber/Matrix Bondings and the Microstructure of C/C-SiC Composites

Modifying the fiber/matrix bondings allows the tailoring of the microstructure and mechanical properties of LSI-derived C/C-SiC
composites. In this process, a porous C/C body is infiltrated with liquid silicon which reacts with carbon to SiC at the crack
surfaces. The interface in the CFRP composite (green body) can be adapted by a prior pretreatment of the carbon fibers in
nitrogen atmosphere at various temperatures, which leads to a reduction of oxygen functional groups on the fiber surfaces.
This selectively induced defunctionalization lowers the fiber/matrix bondings with increasing fiber pretreatment temperatures.65,66
 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs 11

Fig. 13 Stress–strain curves of Nextel 610/Al2O3–ZrO2 composites (eight layers Nextel DF11) measured via 3-point-bending test according to DIN
EN 658-3.

Fig. 14 Tensile stress–strain curves of 2D fiber reinforced (a) SiC/SiC material representing weak interphase composite (WIC) materials
(reproduced from Camus, G., 2000. Modelling of the mechanical behavior and damage processes of fibrous ceramic matrix composites:
Application to a 2-D SiC/SiC. International Journal of Solids and Structures 37 (6), 919–942) and (b) C/C composite representing weak matrix
composite (WMC) materials (reproduced from Koch, D., Tushtev, K., Grathwohl, G., 2008. Ceramic fiber composites: Experimental analysis and
modeling of mechanical properties. Composites Science and Technology 68 (5),1165–1172) with angles of þ 45/  45 degree and 0/90 degree
between the fiber orientation and loading direction.

Fig. 16 shows C/C resp. the resulting C/C-SiC microstructures with weak and strong fiber/matrix bondings. The results
are a different mechanical behavior and different fracture surfaces (Fig. 17). Composites with fibers of moderate treatment
at 4001C exhibit only a slight fiber bundle pull-out and the material shows no damage-tolerant fracture behavior.
Further reduction of the fiber/matrix bondings by thermal fiber pretreatments at 600 and 8001C increase the fiber pull-out
12 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs

Fig. 15 Classification of ceramic matrix composites (CMCs) into weak matrix composite (WMC) and weak interphase composite (WIC).
Reproduced from Koch, D., Tushtev, K., Grathwohl, G., 2008. Ceramic fiber composites: Experimental analysis and modeling of mechanical
properties. Composites Science and Technology 68 (5), 1165–1172.

Fig. 16 Micrographs of (a) C/C after pyrolysis with high fiber/matrix bondings and (b) resulting C/C-SiC; (c) C/C after pyrolysis with low fiber/
matrix bondings and (d) resulting C/C-SiC.

effect, resulting in higher strengths and strains of the composites. A fiber pretreatment at 16001C further reduces the
fiber/matrix bondings, resulting in cracks at the fiber/matrix interface during the pyrolysis. The cracks are filled with silicon in the
subsequent silicon infiltration step. The C/C-SiC composites become more brittle, since the fibers are partially converted to
SiC and have very strong bonds formed by the SiC. The results are a decreasing strength and elongation at break and an increasing
SiC content.
The reason for the decreasing fiber/matrix-bondings can be found in a removal of oxygen from the fiber surface by the thermal
treatment in oxygen free atmospheres. The oxygen groups, which were applied directly after the fiber production during an
oxidation process, were removed with the release of CO, CO2, and H2O. The investigation of carbon fibers with X-ray photo-
electron spectroscopy (XPS), which is a highly surface sensitive method, displays the reduction of oxygen from fiber surfaces with
increasing temperature (Fig. 18).
This favorable tailoring of the microstructure can be utilized to produce graded C/C-SiC composites, for example, for friction
pads. Fabric plies with high strength and damage tolerance form the core of the pad, an outer friction layer with a high SiC-content
 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs 13

Fig. 17 Fracture surfaces of C/C-SiC composites manufactured via liquid silicon infiltration (LSI) process after 4-point-bending test and different
fiber pretreatments in nitrogen atmosphere: (a) 4001C, (b) 8001C, (c) 16001C, (d) selected corresponding stress–strain curves. Reproduced from
Reichert, F., Langhof, N., Krenkel, W., 2015. Influence of thermal fiber pretreatment on microstructure and mechanical properties of C/C-SiC with
thermoplastic polymer-derived matrices. Advanced Engineering Materials 17 (8), 1119–1126.

Fig. 18 Oxygen concentration on carbon fiber surfaces determined via X-ray photoelectron spectroscopy (XPS). Colors indicate different test series.

reduces wear and ensures a constant and high coefficient of friction. An example of a gradient C/C-SiC composite is shown in
Fig. 19, which is used as a brake pad material, in elevator emergency brakes.27

5.1.4 New Design Philosophies and Industrial Applications

The ability to sustain tensile stresses allows new design approaches and innovative applications with CMCs. Besides aerospace and
military applications, CMCs play a more and more important role as structural material for industrial applications. Fig. 20 shows two
examples, where CMCs replaced successfully their metallic counterparts and their use leads to considerably longer lifetimes.4,26,27,66
14 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs

Fig. 19 Cross section of a liquid silicon infiltration (LSI)-derived C/C-SiC composite with an increasing SiC-content from the core to the surfaces.
The SiC content was adjusted with a thermal fiber pretreatment. Reproduced from Frieß, M., Krenkel, W., Kochendörfer, R., et al., 2005. High-
temperature materials and hot structures. In: Jacob, D., Sachs, G., Wagner, S. (Eds.), Basic Research and Technologies for Two-Stage-to-Orbit
Vehicles. Weinheim: WILEY-VCH Verlag Gmbh, pp. 499–580.

Fig. 20 High performance brake disk made of C/C-SiC (courtesy of Volkswagen AG) and kiln furniture made of Ox/Ox composites (courtesy of
Schunk Kohlenstofftechnik GmbH).

The unique properties of CMC materials offer new design possibilities, in particular, if they are joined with other materials in
so-called hybrid structures. Fig. 21 shows exemplarily a jacketed metallic pipe, which is reinforced by an outer jacket made of CMC.
The intended applications of this hybrid design are hot steam containing pipes in power plants. The metallic pipe ensures the
tightness and corrosion resistance against steam whereas the CMC reinforcement prevents the creep deformation of the metallic
pipe at high temperatures, thus allowing higher operating temperatures and a longer service life in comparison to steel pipes
without reinforcements. An interphase between steel pipe and CMC reinforcement compensates the different CTE between steel
and CMC material.
A wound oxide CMC structure is used, which allows an optimal alignment of the fibers with the load direction. Oxide fibers as
well as oxide matrices are necessary to avoid oxidation in air, which would weaken non-oxide CMCs during the long service time.
Additionally, the maximum processing temperature for the CMC synthesis is limited to about 7501C to prevent any changes in the
microstructure of the steel pipe. Thus, a matrix is required that shows a low viscosity to allow infiltration and offers an unusual low
 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs 15

Fig. 21 Jacketed pipe comprising an inner steel pipe with a ceramic matrix composite (CMC) reinforcement to increase the maximum operating
temperature and to suppress ternary creep of the metal pipe.

Fig. 22 Circumferential stresses in the inner steel and outer ceramic matrix composite (CMC) jacket at four representative radii in dependence on
the coefficients of thermal expansion (CTE) of the CMC material. Reproduced from Spatz, C., 2015. Entwicklung einer faserkeramischen Armierung
für heißdampfführende Stahlrohre zur Vermeidung von Kriechverformung, first ed. PhD Thesis, Göttingen, Cuvillier Verlag.

Fig. 23 Circumferential stresses in the inner steel and outer ceramic matrix composite (CMC) pipe in dependence on the CMC wall thickness.
Reproduced from Spatz, C., 2015. Entwicklung einer faserkeramischen Armierung für heißdampfführende Stahlrohre zur Vermeidung von
Kriechverformung, first ed. PhD Thesis, Göttingen, Cuvillier Verlag.

sintering temperature below 7501C. A carbon containing polysiloxane was chosen for the matrix precursor,67 which allows the
formation of a 3D-network without gaseous crosslinking products or catalysts by applying the PIP-process.
When pressurized up to 300 bar and exposed up to 7001C, the lower CTE of the CMC jacket in comparison to the metallic pipe
results in tensile stresses in the reinforcement and reduces the stresses in the metal pipe, thus preventing it from ternary creep. The
stress distribution is dependent on several parameters, for example, the CTE mismatch between both components, the wall
thickness ratios and the stiffness of the CMC reinforcement, which can be varied by different winding angles a. These relationships
are displayed in Figs. 22–24. To demonstrate the influence of these parameters, finite element simulations have been conducted
where the overall temperature was set to 7001C and the pressure of the inner metal pipe was up to 300 bar. Fig. 22 shows the hoop
16 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs

Fig. 24 Circumferential stresses in the inner steel and outer ceramic pipe in dependence of the Young’s modulus of the ceramic matrix
composite (CMC) material. Reproduced from Spatz, C., 2015. Entwicklung einer faserkeramischen Armierung für heißdampfführende Stahlrohre zur
Vermeidung von Kriechverformung, first ed. PhD Thesis, Göttingen, Cuvillier Verlag.

Fig. 25 High temperature tests of the jacketed pipe in comparison with the metallic pipe. Reproduced from Spatz, C., 2015. Entwicklung einer
faserkeramischen Armierung für heißdampfführende Stahlrohre zur Vermeidung von Kriechverformung, first ed. PhD Thesis, Göttingen, Cuvillier Verlag.

stresses in the pipe and the reinforcement in dependence on a low (1.6  106K1) and high (7.4  106K1) CTE of the CMC
material in circumferential direction. All stresses are normalized to the inner hoop stress of the steel pipe. If no gap is applied
between the two components, the steel pipe is loaded with compressive stresses due to the higher CTE of steel in comparison to
CMC. CMC reinforcements with higher CTEs, which can be realized by lower winding angles with respect to the pipe axis, reduce
the tensile stress level of the CMC reinforcement due to the smaller CTE mismatch (CTESteel ¼ 13  106K1).
The tensile resp. compressive stresses of the jacketed pipe in dependence on the wall thickness ratio between pipe and CMC
jacket is shown in Fig. 23. With increasing CMC wall thickness, the tensile stresses in the CMC pipe decrease logically. If the tensile
stresses in the CMC reinforcement are higher than in the unjacketed pipe, the pipe is always kept under compression. Such a stress
distribution needs dimensions of the jackets, where the wall thicknesses are much higher than the corresponding metallic pipe.
However, to fulfill the requirements of the jacket, it is not necessary to keep the metal pipe under compression. Therefore, only
thin-walled CMC jackets must be applied to suppress the undesirable creep of the pipe.
Higher Young’s moduli of the CMC jacket in circumferential direction realized by higher winding angles, lead to higher tensile
stresses in the CMC material, as it is shown in Fig. 24.
Prototypes of the jacketed pipes were manufactured and tested successfully and demonstrated the feasibility of this hybrid
concept (Fig. 25). The reference metallic pipe without jacket failed after less than 500 h at 6001C and 35 MPa inner pressure, while
the jacketed pipe withstands these harsh mechanical conditions for more than 3500 h.

5.1.5 Summary

CMCs are a unique and relatively new class of structural materials. Depending on the processing route and the type of interphase
resp. interface, a wide variability of mechanical as well as thermophysical properties can be achieved and result in new applica-
tions. All CMC materials are characterized by a porous and/or micro-cracked matrix, by a high anisotropy and by a high fracture
toughness. These properties can be used in new designs of novel hybrid structures, where the favorable properties of CMC
materials are combined with those of conventional materials like metals or monolithic ceramics.
 Design Objectives and Design Philosophies, Interphases and Interfaces in Fiber-Reinforced CMCs 17

See also: 5.2 Non-Oxide/Non-Oxide Ceramic Matrix Composites – Composite Design for Tough Behavior. 5.3 Advanced SiC/SiC Composite
Systems. 5.4 Chemical Vapor Infiltration Processing of Ceramic Matrix Composites. 5.5 Oxide/Oxide CMCs – Porous Matrix Composite
Systems; Composites With Interface Coatings. 5.6 Mechanical Behavior of Non-Oxide Fiber-Reinforced CMCs at Elevated Temperature:
Environmental Effects. 5.7 Mechanical Behavior of Oxide–Oxide Fiber-Reinforced CMCs at Elevated Temperature: Environmental Effects. 5.8
Modeling Mechanical Behavior of Ceramic Matrix Composites. 5.9 Geopolymer-Based Composites. 5.10 Ultra-High Temperature Ceramic-Based
Composites. 5.11 Joining and Machining of CMCs. 5.12 Nondestructive Evaluation – Use of Acoustic Emission for CMCs. 5.13 Development
History of GE's Prepreg Melt Infiltrated Ceramic Matrix Composite Material and Applications

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