James A. DiCarlo

Glenn Research Center, Cleveland, Ohio

Advances in SiC/SiC Composites for

Aero-Propulsion

NASA/TM—2013-217889

July 2013
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James A. DiCarlo

Glenn Research Center, Cleveland, Ohio

Advances in SiC/SiC Composites for

Aero-Propulsion

NASA/TM—2013-217889

July 2013

National Aeronautics and

Space Administration

Glenn Research Center

Cleveland, Ohio 44135
Acknowledgments

Regarding the SiC/SiC advancements made over the last decade, the author gratefully acknowledges the funding support of the 

NASA Ultra Effi cient Engine Technology program, the NASA Glenn Director’s Discretionary Fund, and the NASA Fundamental 

Aeronautics program. As major contributors to these advancements, the author wishes to acknowledge the exceptional professional 

efforts of H.M. Yun, R.T. Bhatt, and G.N. Morscher, plus the excellent technical support of R. Phillips, R. Babuder, and R. Angus.

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Available electronically at http://www.sti.nasa.gov

Trade names and trademarks are used in this report for identifi cation 

only. Their usage does not constitute an offi cial endorsement, 

either expressed or implied, by the National Aeronautics and 

Space Administration.

Level of Review: This material has been technically reviewed by technical management. 
NASA/TM—2013-217889 1 

Advances in SiC/SiC Composites for Aero-Propulsion 

James A. DiCarlo 

National Aeronautics and Space Administration 

Glenn Research Center 

Cleveland, Ohio 44135 

Abstract 

In the last decade, considerable progress has been made in the development and application of 

ceramic matrix composites consisting of silicon carbide (SiC) based matrices reinforced by smalldiameter continuous-length SiC-based fibers. For example, these SiC/SiC composites are now in the early 

stages of implementation into hot-section components of civil aero-propulsion gas turbine engines, where 

in comparison to current metallic components they offer multiple advantages due to their lighter weight 

and higher temperature structural capability. For current production-ready SiC/SiC, this temperature 

capability for long time structural applications is ~1250 C, which is better than ~1100 C for the best 

metallic superalloys. Foreseeing that even higher structural reliability and temperature capability would 

continue to increase the advantages of SiC/SiC composites, progress in recent years has also been made at 

NASA toward improving the properties of SiC/SiC composites by optimizing the various constituent 

materials and geometries within composite microstructures. The primary objective of this chapter is to 

detail this latter progress, both fundamentally and practically, with particular emphasis on recent 

advancements in the materials and processes for the fiber, fiber coating, fiber architecture, and matrix, 

and in the design methods for incorporating these constituents into SiC/SiC microstructures with 

improved thermo-structural performance. 

1.0 Introduction 

Silicon carbide (SiC) ceramic matrix composites (CMC) reinforced by continuous-length, highperformance, small-diameter (~10 to 15 μm), polycrystalline SiC fibers are considered enabling materials 

for a variety of advanced applications where lightweight re-usable structural materials are required to 

operate for long time periods within extreme high-temperature environments. Today the first generations 

of SiC/SiC CMC with thermo-structural capability to ~1250 C are being introduced into the hot-section 

components of military and commercial gas turbine engines (Refs. 1 and 2). In comparison to superalloy 

metallic components, which at best operate as high as ~1100 C, these CMC components will not only 

offer reduced engine weight, but also reduced component cooling-air requirements. Reduction in cooling 

air would then have the additional engine benefits of improved thrust-to-weight ratio, reduced fuel burn, 

and reduced harmful exhaust emissions. In order to seek further increases in these engine benefits, 

research has been on-going within NASA aimed at developing SiC/SiC CMC with even higher structural 

reliability and temperature capability. 

The objective of this chapter is to detail, both fundamentally and practically, the key advances from 

recent SiC/SiC CMC research activities within NASA. Although these advances are primarily still under 

development and optimization, they do show various degrees of improvement over the materials and 

processes currently being implemented for SiC/SiC engine components. To this end, Section 2.0 of this 

chapter presents information on the general constituent materials and process requirements for structurally 

reliable high-temperature SiC/SiC engine components. Section 3.0 then discusses the primary fabrication 

routes currently being employed for production-ready SiC/SiC components in terms of their benefits and 

limitations. Section 4.0 then details recent advancements in the materials and processes for the SiC fiber, 

the fiber interfacial coating, the fiber geometry or architecture, the SiC-based matrix, and in the design 

methods for assembling these microstructural constituents into higher-temperature higher-performance 

SiC/SiC composites and components. Where possible, CMC property data will be presented and analyzed 
NASA/TM—2013-217889 2 

that directly compares the materials and processes of the advanced approaches against those of current 

approaches. This property comparison will focus primarily on such key CMC properties as in-plane and 

thru-thickness tensile strength, thermal conductivity, creep resistance, and rupture behavior since it is 

these properties that directly control the CMC structural and upper temperature capability. Finally 

Section 5.0 summarizes these results into current design guidelines for achieving SiC/SiC microstructures 

with improved intrinsic thermo-structural capability, and then discusses potential issues, such as thermomechanical shock, environmental attack, and creep-related residual stress development, which still require 

further research to understand whether they will be areas of concern for real engine components. 

2.0 Materials and Process Requirements for Structurally Reliable 

High-Temperature SiC/SiC Components 

As with any structural ceramic material, the first key performance requirement for a SiC/SiC CMC is 

the ability to display three-dimensional (3D) strength properties or allowables that are high enough to 

permit a component to achieve its desired service life. Like metallic hot-section components, ceramic 

components for aero-propulsion gas turbine engines will typically be designed to be thin-walled and 

cooled using conventional approaches such as film cooling (vanes, blades) and backside convective 

cooling (combustion liners, shrouds). Thus component stresses typically arise from a combination of 

mechanical, aerodynamic, and thermal-gradient loads that produce both in-plane and thru-thickness 

stress-temperature distributions that can vary with time within the component walls and attachment areas. 

Under these stress conditions over a time period generally equivalent to the component inspection cycle, 

the SiC/SiC composite’s in-plane Matrix Cracking Strength (MCS) should be sufficiently high enough to 

reliably avoid cracks that propagate into and across the component wall, thereby preventing aggressive 

application environments from reaching and degrading the crack-bridging fibers. Of additional 

importance is the ability of the CMC to avoid the occurrence of delamination cracks within the wall since 

these can significantly reduce the thru-thickness thermal conductivity, further increasing the thermal 

gradient stresses and the probability of wall spallation. Thus the constituents and geometries of the 

as-produced CMC should be optimized to assure that the composite’s matrix cracking strengths both inplane and thru-thickness sufficiently exceed the expected component service stresses in these directions. 

In terms of maintaining these 3D strength characteristics at high temperatures, the key requirements 

of the SiC/SiC fiber and matrix constituents are (1) to be thermally stable at the expected application 

temperatures and (2) to avoid fracture under the maximum time, temperature, and stress conditions that 

they will experience during component service. Today these constituents will most always have a 

polycrystalline SiC microstructure in part due to the fact that single crystal microstructures are too costly 

to produce. Also, as seen for single crystal oxide fibers in oxide matrices, single crystal microstructures 

may not offer any significant performance advantage over polycrystalline materials due to rapid atomic 

diffusion on the fiber surface and pitting-induced fiber strength loss caused by irregular contact points 

with the multiconstituents in the composite environment (Ref. 3). Furthermore, if the SiC fiber and matrix 

contain a pure stoichiometric composition and grain sizes near ~500 nm, their intrinsic microstructural 

and strength stability under zero stress conditions is not an issue for thousands of hours at temperatures up 

to at least 1600 C. However, as soon as stress is applied to a SiC/SiC component, the grains in the 

constituent microstructures will begin to slide past each other, giving rise to measurable time-dependent 

deformation or creep that could start as low as ~1100 C depending on the impurity phases in the grain 

boundaries. Two practical issues with constituent creep are undesirable dimensional changes for the CMC 

component and the growth of cavitation flaws at the grain triple points within the constituent 

microstructures. Eventually these flaws grow large enough to cause time-dependent constituent fracture 

or rupture, which in the worst case could cause rupture of the total SiC/SiC component. Unlike metals 

that can rupture at creep strains well over 1 percent so that undesirable dimensional changes are typically 

their life limiting criteria, SiC-based constituents will typically rupture at strains below 1 percent so that 

their use lives are primarily limited by rupture rather than by dimensional stability issues. Thus for 
NASA/TM—2013-217889 3 

structurally reliable high-temperature operation for long times, the constituents of SiC/SiC composites 

should have microstructures and compositions that display the optimum in creep and rupture resistance 

under the stress and temperature conditions that they will encounter within the final application. 

Another life-limiting and creep-related mechanism for high-temperature SiC/SiC components is that 

the constituents almost always differ in creep behavior due to different compositions and/or to different 

stresses related to their geometries within the composite. To appreciate this issue, one can consider the 

situation where a constant and uniform mechanical load is applied to a composite and the fiber and matrix 

within the composite display significantly different creep behavior. For Case A, where the matrix is more 

creep prone than fiber, the matrix will relax its initial stress and shed more of the CMC load onto the 

fibers, so that the composite upper use temperature will then typically depend on such life-limiting issues 

as excessive fiber deformation, fiber time-dependent fracture or rupture, and/or adverse residual stress 

development within the composite. Regarding this last issue, after creep under uniform CMC tensile 

loads, the internal residual stress on the matrix at lower temperatures will be compressive, so that the 

CMC in-plane MCS will beneficially increase in this temperature regime (Ref. 4). However, after CMC 

creep under compressive loads, tensile residual stresses on the matrix will develop at lower temperatures, 

thereby decreasing the CMC MCS at these lower temperatures and increasing the chance of thruthickness cracking if the local CMC load should become tensile. This reduction in MCS has been shown 

to be a likely scenario for components with non-uniform stresses in their walls, as for example when 

subject to thermal gradients across the walls (Ref. 5). As will be discussed, this Case A scenario is the 

situation with current SiC/SiC composites containing free silicon in the matrix. For Case B, where the 

matrix is more creep resistant than the fiber, the fiber in its primary creep stage will quickly shed its load 

to the matrix, so that initial composite tensile loads must be kept low to avoid the matrix stresses rising 

over time above their cracking limit. Bottom line is that, if possible, both the fiber and matrix should be 

highly creep resistant under the component service conditions; but if this is not the case, the matrix creep 

should be slightly greater than that of the fiber, so that during the component service life, the matrix 

carries some of the composite load and remains protective of the fiber without cracking. 

3.0 Current Fabrication Routes for SiC/SiC Engine Components 

Today two primary SiC/SiC fabrication routes have achieved to a large degree the material and 

process requirements needed to produce complex-shaped engine components with microstructures 

capable of reliable structural operation for long times at temperatures greater than those of metallic 

superalloys. The first route was developed under the NASA Enabling Propulsion Materials (EPM) 

Program and was demonstrated by the production of small and large combustor liners (Ref. 6). Here 

continuous-length multifiber tows of a small-diameter high-performance SiC fiber type are woven into 

continuous fabric pieces, typically with a 0/90 balanced five-harness satin (5HS) fiber architecture. These 

two-dimensional (2D) fabric pieces are then cut and judiciously stacked into component-shaped 3D preforms with fiber volume fractions up to 40 percent. Chemical vapor infiltration (CVI) is employed to first 

apply an interfacial coating or interphase on the fibers within the tows, and then to apply a thin protective 

SiC matrix layer over the interphase. To adequately achieve the key interfacial coating requirements of 

high environmental durability and high compliance for matrix crack deflection, the interphase 

composition is typically based on silicon-doped boron nitride (BN). The coating thicknesses can vary 

from ~0.1 to 0.5 μm depending on such factors as the fiber surface roughness and the degree of protection 

needed to avoid coating degradation during matrix formation and/or component service. The thin CVI 

SiC matrix coating is important in that it not only protects the interphase and fibers from subsequent 

matrix infiltration processes and from external environments, but it also provides a highly thermally 

conductive and creep-resistant matrix constituent. As will be discussed, NASA has determined that these 

latter two matrix properties, as well as the protective nature of the interface coating, can be further 

improved prior to the final matrix infiltration steps by a high temperature treatment of the CVI SiC-coated 

pre-form (Ref. 7). In the final steps of matrix fabrication, SiC particulate is slurry-cast into the component 

pre-form near room temperature, followed by a finishing step in which a molten metallic silicon alloy is 
NASA/TM—2013-217889 4 

infiltrated near 1400 C into any open remaining matrix porosity. The silicon bonds the SiC particulates 

together with little chemical reaction. This last step of this non-reactive melt infiltration (NRMI) route is 

important in that it results in a SiC-based matrix with low residual porosity, high thermal conductivity, 

and good environmental protection for the coated fibers. In addition, it minimizes the existence of large 

matrix flaws, thereby raising the stress levels required for formation of detrimental thru-thickness and 

delamination cracks in the composite. In following discussions, the matrices formed from the NRMI 

approach of CVI, slurry, and silicon melt infiltration (MI) will often be referred to as the CVI-MI 

matrices for simplicity. 

In the second fabrication route, which has been developed and demonstrated by General Electric for a 

variety of turbine hot-section components, multifiber tows of a small-diameter high-performance SiC 

fiber type are first spread and coated by chemical vapor deposition (CVD) with a layered interfacial 

coating consisting of BN and silicon nitride (Si3N4) compositions (Ref. 8). The coated tows are then prepregged into unidirectional 2D tapes using a polymer-based binder containing SiC and carbon 

particulates. The tapes are then cut, oriented, and stacked into 3D pre-forms with total fiber volume 

content less than 30 percent. To achieve the component structural and shape requirements, the 3D preforms are assembled using methods much like those used for tape lay-up of polymer matrix composites. 

In the final steps of this “prepreg” approach, the 3D pre-form is heated to high temperature to decompose 

the polymer and to form the final matrix composition by infiltration with molten silicon near. Unlike the 

NMI route, the SiC-based matrix is formed in-situ by the reaction of the silicon with the residual carbon 

in the preform. During this reactive MI (RMI) process, there is an exothermic reaction between silicon 

and carbon, resulting in an expansion of the matrix that puts the fibers under residual tension and the final 

matrix under residual compression, which is beneficial in increasing the composite in-plane cracking 

strength at low temperatures. However, the RMI route also causes a rise in local temperatures, which if 

not controlled could cause strength loss in those SiC fiber types that are produced at maximum temperatures 

near 1400 C. In addition, unless the interphase coating is thick enough (~1 μm) with the proper 

compositions, these temperature increases can increase the risk of fiber attack by the thermally-induced 

diffusion of excess free silicon across the coating. Thus SiC fibers with the highest production temperatures 

are generally preferred for strength retention during both the NRMI and RMI matrix fabrication routes. 

In terms of temperature capability, both current CMC fabrication routes always result in some excess 

silicon in the matrix, thereby allowing SiC/SiC components under their application loads to display 

undesirable matrix creep that begins at ~1100 C, and becomes excessive above ~1300 C. As described 

in Case A above, if the SiC fiber is more creep resistant than these matrices, effectively all the composite 

load will be shifted to the fiber, resulting in composite creep and rupture at the higher temperatures 

controlled by the fiber creep-rupture characteristics and its volume fractions. In addition, at temperatures 

approaching ~1400 C, not only does the excess silicon melt and tend to exude from the composite, but 

the beneficial residual stresses in the RMI route will tend to disappear. Thus for current production 

SiC/SiC composites fabricated by both MI routes and with high-temperature SiC fibers, it is the final 

silicon infiltration step and not the reinforcing SiC fiber that is the prime mechanism that limits the 

intrinsic CMC upper use temperature to only ~1250 C. 

4.0 Recent NASA Advancements in SiC/SiC Materials and Processes 

Although the current CMC fabrication routes described above are adequate for the introduction of 

SiC/SiC components into the hot sections of aero-propulsion turbine engines, NASA has determined that 

these production approaches not offer the optimum in SiC/SiC composite structural reliability or in high 

temperature service capability. The following sections will detail the various deficiencies of these routes 

and the advancement approaches recently developed by NASA and demonstrated with industry 

collaboration to address these short-comings. 
NASA/TM—2013-217889 5 

4.1 Advances in SiC-Based Fibers 

First-level property requirements for a SiC fiber in high-performance high-temperature SiC/SiC 

components are (1) to be produced in continuous lengths with a small diameter; (2) to display in its asproduced condition high tensile strength (~3 GPa) and high-stiffness or modulus (~400 GPa); and (3) to 

have the ability to retain as much of this strength and stiffness as possible at high temperatures for long 

times under high stresses. Today the best approach for achieving these requirements is by the spinning, 

curing, and pyrolysis of multifilament tows of polymer derived fibers (Ref. 9). For high strength and 

stiffness in their as-produced condition, key fiber microstructural requirements are, respectively, average 

grain sizes below ~500 nm and a near- stoichiometric SiC composition with high density (3.0 gm/cm3

) 

or low porosity (Ref. 10). Although decreasing grain size below 500 nm should theoretically increase 

tensile strength of polymer-derived SiC fibers, this is not observed because inherent in the polymer 

preparation and spinning processes, flaw-related particulates with sizes of the order of ~200 nm are 

allowed to pass into the green fibers. Thus for the finer grained SiC fibers, these particulates rather than 

the fiber grains are limiting the strength to ~3 GPa. For high strength and stiffness retention (creep 

resistance) during service at high temperatures, key microstructural requirements are a pure nearstoichiometric SiC composition and grain sizes as large as possible without exceeding ~500 nm. These 

properties in turn will also provide a SiC fiber with high intrinsic thermal conductivity (50 W/m.K), a 

property very important for cooled components. Finally, on the more practical side, the SiC fiber should 

be commercially available in large quantities and with low acquisition or life-cycle cost. 

Today the commercial SiC fiber type that is primarily being employed for both the NRMI and RMI 

SiC/SiC component fabrication approaches is the “Hi-Nicalon Type S” fiber (hereafter HNS fiber) that is 

produced by Nippon Carbon in Japan. This fiber meets most of the fiber practical requirements, 

particularly commercial availability in large quantities (Ref. 2), but still lacks some important 

microstructural requirements. The HNS fiber, like its “Hi-Nicalon” precursor fiber, is first spun from a 

organometallic “pre-ceramic” polymer precursor, followed by cross-linking or curing under electron 

irradiation. It is then subjected to high-temperature pyrolysis and heat treatment steps to remove the 

excess carbon present in the Hi-Nicalon fiber (C/Si~1.3) and to form a nearly stoichiometric (C/Si~1.05) 

-phase SiC fiber (Ref. 11). Some key chemical and physical properties for the as-produced HNS fiber 

are shown in Table 1, and compared with other SiC-based fibers with potential for high-temperature CMC 

(Refs. 9, 11 to 14). Even though irradiation rather than oxidative curing is employed, the HNS fiber still 

contains residual oxygen (~0.7 wt%), which typically is in the form of a silicon-oxycarbide impurity 

phase. When exposed for short times at temperatures greater than ~1600 C, this phase decomposes into 

SiO and CO gases, which in turn creates pores in the fiber, thereby reducing fiber strength (Ref. 15). In 

addition, the as-produced HNS (NOX grade) fiber has a carbon-rich surface, which can be detrimentally 

removed within the CMC when exposed to combustion environments (Ref. 16). Thus for turbine 

applications, a judicious practice is to utilize HNS (OX grade) fibers where the carbon surface has been 

removed by the fiber vendor or by the fiber end-user prior to CMC fabrication. However, as shown in 

Table 1, this will reduce fiber strength by ~10 percent by exposing flaws on the as-produced fiber surface. 

Adding a compliant interfacial coating to the fiber, such as BN, may act to blunt these flaws again, 

provided the coating is applied to all areas of the fiber surface. 

Regarding structural performance at high temperatures, because of its thermal limit of ~1600 C for 

strength retention, the average grain size within the as-produced HNS fiber does not grow beyond ~20 nm 

(Ref. 9) and the fiber still retains the oxide-based impurity phase in its grain boundaries (Ref. 12). Recent 

studies at NASA (Ref. 17) suggest that these two factors are the prime contributors to undesirable fiber 

creep and a structural limit for this fiber type of less than ~1450 C. To better understand this limitation, 

Figure 1 compares the creep behavior of various polymer-derived SiC fiber types when subjected as 

single fibers to a constant stress () of 270 MPa at a constant high-temperature (T) of 1400 C in 

laboratory air (Ref. 17). One can see that all polycrystalline SiC fibers display a time-dependent transient 

or primary creep strain (p) followed by a nearly steady-state secondary creep strain (s). NASA (Refs. 10 
NASA/TM—2013-217889 6 

and 17) has identified the underlying creep mechanism for the secondary stage of these fine grained SiC 

fibers as Interface Controlled Diffusional (ICD) creep in which the concentration of vacancies supporting 

fiber creep by grain-boundary diffusion is limited by the fine-sized grains (Ref. 18). In this case, the 

second stage creep strain rate (

.

s) can be described analytically as:

  C/d exp  Q/RT 3 s    (1) 

TABLE 1.—SOME KEY PROPERTIES OF SiC-BASED FIBERS FOR HIGH TEMPERATURE SiC/SiC CMC 

Fiber type 

(Trade-Name) 

Vendor Max 

temp.a

Key 

Non-SiC 

elements in 

bulk, wt.% 

Grain 

size b, 

nm 

Key 

surface 

phase 

Tensile 

strength c, 

GPa 

R.T. 

thermal 

cond., 

W/m.K 

Ref. 

Hi-Nicalon Nippon 

Carbon 

1450 C 0.5 Ox d

5 C 3.0 8 9, 11 

Hi-Nicalon 

Type-S (NOX grade) 

Nippon 

Carbon 

1650 C 0.7 Ox d

20 C 2.8 18 9, 12 

Hi-Nicalon 

Type-S (OX grade) 

Nippon 

Carbon 

1650 C 0.7 Ox d

20 SiC 2.6 18 12 

Sylramic ATK-COI 

Ceramics 

1850 C  0.1 Ox, 1.2 B, 

2.4 Ti 

100 B 3.2 46 13 

Sylramic-iBN ATK-COI 

Ceramics 

1800 C  0.1 Ox, ~0 B, 

2.4 Ti 

200 BN 3.1  50 13 

Super 

Sylramic-iBN 

NASA 1800 C  0.1 Ox, ~0 B, 

2.4 Ti 

200 BN 3.0  50 14 

Tyranno SA3 UBE 

Industries 

1900 C 0.2 Ox, 

0.6 Al 

400 SiC 2.8 65 9 

a

~Maximum temperature for loss of 10 percent of as-produced strength after 10 hr inert exposure b

~Average grain size measured by TEM 

c

Average as-produced single fiber strength at room temperature and 25 mm gauge length 

d

Spinning-related strength controlling flaws are ~200 nm in size 

Figure 1.—Creep strain behavior of various SiC fiber types in single fiber form (Ref. 17). 
NASA/TM—2013-217889 7 

Here (C) is a composition-dependent diffusion constant; (d) is average fiber grain size; (Q) is the 

activation energy controlling steady-state creep; and (R) is the universal gas constant (8.314 J/mol.K). 

Thus for a given applied stress and temperature, Equation (1) indicates that the HNS fiber secondary 

creep strain could possibly be reduced by an order of magnitude simply by increasing its grain size from 

20 to 200 nm. But due to decomposition of the silicon-oxycarbide phase pre-existent in this fiber, heat 

treatments above ~1600 C, although increasing the HNS fiber grain size and improving its creep 

resistance, causes an undesirable loss in fiber strength (Ref. 15). 

Besides limiting grain size, the oxide-based phase in the HNS fiber also gives rise to an undesirable 

large primary creep stage for this fiber type (Ref. 17). A convenient method for seeing this effect is by 

examining Table 2, which compares the parameter (m) measured in a Bend Stress Relaxation (BSR) test 

(Ref. 19) after 1 hr exposure at various temperatures in argon. The listed (m) data are not only for the 

HNS fiber, but also for two other high-performance SiC fiber types that possess less oxygen content after 

fabrication (Ref. 20). Since stress relaxation is caused by creep, the following equation for (m) values 

0.5 is a good approximation for a fiber’s primary creep strain (p)(Ref. 17): 

1/m 1 p  e  (2) 

Here e = /E = elastic strain on the fiber. Thus one can see in Table 2 that the as-produced HNS fiber 

shows primary creep as low as 1100 C where its (m) value begins to fall below unity after only 1 hr; 

whereas the other two fiber types do not show this behavior until higher temperatures. It is believed that 

this creep is caused by grain boundary sliding accommodated by the low viscosity silicon-oxycarbide 

phase within the HNS boundaries (Ref. 21). Furthermore, on the practical side, one can see by 

Equation (2) and Table 2 that the HNS primary strain is approximately equal to its elastic strain after only 

1 hr at ~1450 C (m = 0.5). This implies a 50 percent loss in fiber stiffness or modulus in just that 1 hr 

period which, depending on the matrix type, could cause a stress overload on the matrix, causing it to 

crack. Thus based on the observed limitations of the HNS fiber, in order to achieve SiC/SiC composites 

with improved structural performance and temperature capability, it is important to seek SiC fibers with 

little or no low-viscosity phases in the grain boundaries and with average grain sizes up to ~500 nm. The 

guideline for little or no low-viscosity phases in the grain boundaries is also important for creep of the 

SiC matrix. 

One proven approach for obtaining SiC fibers with greater temperature capability than the HNS fiber 

is to allow by thermal treatment the silicon-oxycarbide phases that are typically created during fiber 

fabrication to decompose as gases, resulting in a weak SiC fiber with fine-sized pores. Then by 

incorporating sintering aids in the fiber pre-ceramic polymer (Refs. 9 and 22) or by gaseous infiltration of 

the pores with sintering aids after decomposition of a SiCO fiber (Refs. 23 and 24), the pores are then 

removed by sintering treatments to 1800 C and above. This approach typically results in high-density 

near-stoichiometric SiC fibers with little remaining oxide phases and with average grain sizes greater than 

100 nm, two important microstructural conditions not available in the HNS fiber. Also this sintering 

approach has the major advantage of producing strong fibers at very high production temperatures, which 

can significantly open the temperature window for CMC fabrication and service without experiencing 

microstructural instabilities in the fiber reinforcement. 

TABLE 2.—AVERAGE BSR (M) VALUES FOR SINGLE NEAR-STOICHIOMETRIC SiC-BASED 

FIBERS MEASURED AFTER 1-HR EXPOSURE IN ARGON AT 

VARIOUS TEMPERATURES (REFS. 19 AND 20) 

Fiber type 1100 C 1200 C 1300 C 1400 C 1500 C 

Tyranno SA3 ---- 0.90 0.74 0.45 ---- 

Hi-Nicalon Type-S 0.96 0.89 0.83 0.64 0.40 

Sylramic-iBN 1.00 0.97 0.91 0.84 0.74 

 
NASA/TM—2013-217889 8 

Today two commercial SiC fiber types are using this sintering approach: the “Tyranno SA3” fiber 

from UBE Industries where an aluminum-based sintering aid is incorporated into a SiC-yielding preceramic polymer (Ref. 9); and the “Sylramic” fiber originally produced by Dow Corning and now by 

ATK-COI Ceramics in San Diego, California, where a boron-based sintering aid is infiltrated into the 

pores of a decomposed SiCO precursor fiber (Ref. 24). Table 1 lists some key as-produced properties of 

these two fiber types, and Figure 1 compares the tensile creep behavior of these fibers against HNS under 

the same test conditions of 1400 C, 270 MPa, and laboratory air. As can be seen, although the Tyranno 

SA and Sylramic fibers are near-stoichiometric with high strength and high thermal conductivity, their 

high-temperature tensile creep behavior is worse than that of the HNS fiber. This less than optimum creep 

behavior can be attributed to the aluminum and boron-based sintering aids that still remain in the grain 

boundaries of the as-produced fibers, thereby enhancing the rate of grain boundary sliding. Realizing this 

issue, NASA developed high-temperature thermal treatments in selected environments to remove or alter 

the sintering aids in the as-produced fibers. It was determined that whereas the grains grew and creepresistance improved, the low tensile strength of the as-produced Tyranno SA fibers degraded further with 

high-temperature thermal treatment. On the other hand, the boron in the Sylramic fiber sintering aids 

could be removed in both nitrogen and argon atmospheres with retention of high tensile strength 

(~3 GPa), and with significant improvement in fiber creep resistance (Ref. 25). When treated in one 

atmosphere of high-purity nitrogen, the resulting fiber has been labeled “Sylramic-iBN” because as the 

boron is removed from the as-produced Sylramic fiber, a thin protective in-situ grown BN (iBN) coating 

is formed with uniform thickness on the circumference of each fiber, even those in tight contact with its 

neighbors within a woven tow (Ref. 6). When treated in high-purity nitrogen above one atmosphere, 

fibers with even more creep resistance have been produced and have been labeled “Super Sylramic-iBN” 

(Ref. 14). Key properties and the improved tensile creep behavior of the Sylramic-iBN SiC fiber 

(hereafter iBN fiber) and the Super Sylramic-iBN fiber over their precursor Sylramic fiber are shown in 

Table 1 and Figure 1, respectively. 

It is interesting to note in Figure 1 that under creep testing at 1400 C, the iBN fiber is slightly less 

creep resistant than the HNS fiber; whereas the Super iBN fiber is slightly more creep resistant. By 

Equation (1), if grain size were the only difference, this result would be unexpected since the average 

grain size of both iBN fibers is at least an order of magnitude larger than that of the HNS fiber. On the 

other hand, the bend stress relaxation (BSR) results of Table 2 show a significant improvement in creep 

resistance in the primary stage for the iBN fiber over the HNS fiber. Recent NASA studies attribute these 

apparently contradicting results to the fact that after sintering, the Sylramic as well as the Tyranno SA 

fibers display a cross-sectional microstructure consisting of large grains in a thin shell at the fiber surface 

and a core region consisting of smaller grains and fine-size pores (Ref. 26). It would appear that the shell 

region is created prior to or during sintering by the concentration near the fiber surfaces of the aids 

responsible for grain growth. Thus, similar to the fiber-matrix creep conditions described above for 

Case A, at high temperatures the more creep-prone inner region within the fibers will shed its load to the 

more creep-resistant shell region, which only occupies ~30 percent of the total cross-sectional area. As 

such, the stress on the shell can be as much as ~3 times larger than that on the HNS fiber, which contains 

a more uniform cross-section (Ref. 26). Then by Equation (1), the tensile steady-state creep rate for the 

iBN would be ~33 times greater than the HNS fiber at the same grain size. Thus, the perceived tensile creep 

advantage of a larger grain size for the iBN fiber is effectively lost by its nonuniform cross-section. Also, it 

follows that since the BSR test primarily samples the creep-resistant shell region, the BSR data for the iBN 

fiber would be expected to show much better creep-resistant behavior than that of the HNS fiber. 

Because it can be more conveniently produced without the need for high-pressure nitrogen, the iBN 

fiber in contrast to the Super Sylramic-iBN fiber has gone beyond the developmental stage to commercial 

availability as multifilament tows. It also has been available for nearly a decade to be employed in 

SiC/SiC CMC panels as 2D-woven fabric for the NRMI fabrication route and for 3D-woven pre-forms 

using CVI-MI and other matrix infiltration approaches. This has allowed a direct property comparison 

between CMC reinforced with the HNS and iBN fibers. For example, during creep-rupture tests at 

1315 C in laboratory air (Ref. 27), the HNS CMC showed a significant initial increase in creep strain in 
NASA/TM—2013-217889 9 

comparison to the iBN CMC when the applied on-axis stress increased from 103 to 138 MPa. Each CMC 

type was fabricated by the NRMI route with 2D-woven balanced fabric and ~18 percent fiber content in 

the loading direction. It would appear that this behavior is caused by a lower MCS for the HNS CMC 

and/or by the larger primary creep stage of the HNS fiber. This later mechanism is supported by the BSR 

data in Table 2 that shows after 1 hr at 1300 C, the creep of the iBN fiber is much less than that of the 

HNS fiber. As a result, on initial CMC loading, the primary creep for the HNS fiber allowed the NRMI 

matrix to exceed its micro-cracking strain limit of ~0.05 percent (CMC modulus ~265 GPa) before it can 

stress relax and shift most of its load to the fiber. Thus it would appear that the higher HNS CMC creep 

behavior at 138 MPa is indicative of local matrix cracking, leaving the crack bridging fibers to carry 

much higher stress levels they would experience in an uncracked CMC (Ref. 4). On the other hand, 

matrix cracking was probably not occurring in the iBN CMC, allowing these fibers to creep at lower 

stress levels and thus providing a CMC with reduced creep. 

For MI CMC specimens at low stresses with no matrix cracking, steady state creep behavior can be 

predicted well by NASA’s analytical creep models determined from the creep behavior of single HNS 

and iBN fibers by assuming that all the fibers after long times are carrying the entire CMC load, (Ref. 17). 

This is not only true for CMC fabricated by the NMI route, but also for CMC fabricated with HNS fibers 

by the RMI route (Ref. 28). This implies that at least in laboratory air, the rupture life of these CMC 

should be predictable by understanding the intrinsic rupture characteristics of these two fiber types. 

However, although these characteristics have been measured for single HNS and iBN fibers as a function 

of stress and temperature (Ref. 29), one should realize that the fiber conditions within a cracked CMC 

during high-temperature application could be considerably different than those employed during simple 

creep-rupture testing of single fibers. For example, the condition of tensile stress being held constant both 

in time and along the fiber length will probably not exist within a CMC due to fiber curvature induced by 

the fiber architecture and to the typical occurrence of CMC matrix micro-cracking, where fiber stresses 

are highest within the matrix cracks and then drop off within the intact segments of the matrix. Also under 

these CMC conditions, not only is fiber gauge length considerably smaller than the 25 or 100 mm length 

typically used for single-fiber creep-rupture testing, but also rupture is controlled by bundles of fibers 

rather than single fibers. In addition, if the testing environment plays a role on fiber rupture, testing in 

laboratory air may give different results than testing within a given composite application environment. 

For this reason, NASA has found it more practically useful to understand and empirically model the 

general rupture characteristics of various fiber types by examining the rupture characteristics of CMC 

specimens in which the reinforcing fibers are (1) loaded on axis with the primary fiber direction at a known 

fiber content, (2) are essentially straight in the loading direction, and (3) the CMC matrix carries little or no 

load because it is more creep prone than the fiber. The results of this approach are discussed in Section 4.5. 

In summary, when compared to the HNS fiber, the iBN SiC fiber is more thermally stable due to its 

higher production temperature, thereby allowing higher SiC/SiC fabrication and service temperatures 

without fiber tensile strength degradation. Its higher production temperature also results in lower oxygen 

content for a reduced primary creep stage, plus larger grains for much improved thermal conductivity and 

the potential for improved creep if its microstructure can be made more uniform in grain size distribution. 

In addition, the in-situ grown BN layer provides more environmental resistance than the carbon-rich or 

non-carbon-rich surfaces of the HNS fiber as well as the CVI produced BN-based interfacial coatings that 

cannot cover the full fiber surface in high fiber fraction CMC. However, beyond its nonuniform 

microstructure, certain other issues are currently limiting optimum application of the iBN fiber. For 

example, its larger grain sizes in comparison to those of the HNS fiber typically result in rougher fiber 

surfaces, which although beneficial for high interfacial stress and tight crack opening, can be detrimental 

in that they can cause fiber-fiber abrasion and fiber fracture when multifiber tows are woven or braided 

into complex pre-forms. Also, as will be discussed in Section 4.5, the larger grains appear to reduce the 

ultimate creep strains that the fiber can reach without rupture. In addition, there is the general increase in 

the acquisition costs for sintered fibers that typically accompanies their higher production temperatures. 

 
NASA/TM—2013-217889 10 

Finally, the Sylramic precursor fibers for the iBN SiC fibers are currently produced commercially by a 

combination of batch and continuous processing, which limits their availability in sufficient quantities for 

production-ready components. 

4.2 Advances in Interfacial Fiber Coatings 

Although the SiC fiber controls many of the key structural properties of SiC/SiC composite systems, the 

other key composite constituents also need to be optimized in composition and geometry for durable 

component use. For example, in early SiC/SiC composites, the preferred interfacial fiber coating 

composition was carbon as formed by the CVI process. However, under certain conditions, it is likely that 

these coatings will be exposed to aggressive environmental conditions, for example within a matrix crack, 

but also at free or cut CMC edges. For this latter condition, NASA has shown that under the conditions of 

high-velocity combustion gases, these carbon-based coatings can be quickly removed by oxidation, and that 

this removal could occur throughout the composite because of fiber-fiber contact (Refs. 16 and 30). 

Working within the EPM program, NASA has demonstrated that a more stable coating would be based on a 

CVI-produced silicon-doped BN composition. This is the case because BN is not only structurally weak for 

crack deflection, but instead of forming a volatile gas upon oxidation, it forms in the presence of silicon a 

nonfugitive product (borosilicate glass), thereby preventing further oxygen ingress into the composite. 

Although silicon-doped BN coatings have proven useful for CMC with free or cut CMC edges, if 

stress-induced matrix crack deflection occurs at the fiber surface rather than on the outside of the 

interfacial coating, the environment entering the matrix crack could reach and degrade the SiC fiber. To 

address this concern, NASA developed CMC thermal treatment processes that would allow the BN 

coating to shrink towards the fiber surface and debond from the matrix, thereby assuring crack deflection 

to occur on the outside of the fiber-protective BN coating (Ref. 31). Even though the interphase coating 

remains attached to the fiber and debonded from the matrix, load transfer between the fibers and matrix is 

still maintained within complex-shaped fiber architectures that allow the interfacial coating to 

mechanically slide against the matrix during the application of stress. In comparison to multilayer 

concepts for interfacial coatings, this “outside debonding” approach avoids the fabrication of complex 

interphase compositions and structures, does not rely on uncertain microstructural conditions for matrix 

crack deflection outside of the interphase, and provides a more reliable approach for retention of the total 

interphase on the fiber surface. As mentioned earlier, this oxidation protective scheme is further enhanced 

by use of the Sylramic-iBN fibers that also have a thin in-situ grown environmentally-resistant BN 

coating on their surfaces. As a result of outside debonding, a SiC/SiC system with 2D fiber architectures 

will typically display a slightly lower thru-thickness thermal conductivity due to partial loss of matrix 

thermal contact with the fibers. However, NASA has shown that this effect can be mitigated by the use of 

3D architectures as discussed in the next section. 

4.3 Advances in SiC Fiber Architectures 

As with any composite component reinforced by high-performance continuous fibers, one of the 

primary requirements is to design the fiber architectures to provide sufficient fiber content in those 

directions that are structurally loaded during component service in order that the fiber-controlled 

structural properties of the component, such as MCS, will be sufficient to allow it to survive without 

failure during its required lifetime. For SiC/SiC components, life-limiting failure mechanisms are 

generally identified with (a) thru-thickness matrix cracking under in-plane tensile loading, and (b) 

delamination cracking under thru-thickness tensile loading. In the first case, thru-thickness cracking will 

allow the service environment to reach and degrade the crack-bridging fibers within time periods typically 

less than the required component lifetime. For aero-propulsion components, this situation would apply, 

for example, to the walls of pressure-cooled airfoils and to centrifugally loaded blades (Ref. 32). In the 

second case, delamination cracking can increase the thermal stresses and wall temperatures for internally 

cooled turbine components, increasing the risk of wall spallation and other component damage. Thus the 
NASA/TM—2013-217889 11 

general fiber architecture goals for SiC/SiC components are to achieve a starting pre-form with (1) a 

shape as close as possible to the final component shape and (2) a total fiber content as high as possible, 

and (3) a judicious splitting of this total fiber content in-plane, thru-thickness, and locally to allow the 

CMC directional MCS properties to sufficiently exceed the component service stresses. Furthermore, if 

the fiber is thermally conductive, it will aid in increasing the CMC directional thermal conductivities, 

thereby reducing service-related thermal stresses in these directions. 

When the current 2D fiber architecture designs of the NRMI and RMI fabrication routes are 

compared to the above goals, certain deficiencies can be observed. For example, regarding CMC in-plane 

properties, although the 2D woven fabric approach of the NRMI fabrication route (and other routes that 

use 2D fabric) have reached total fiber contents as high as 50 percent, their in-plane fibers are not straight. 

Depending on the weave pattern, the effective on-axis fiber modulus and strength will be less than those 

of straight fibers, which in turn will debit CMC thru-thickness matrix cracking strength (MCS), ultimate 

tensile strength (UTS), and rupture strength (to be discussed). In the case of the RMI or prepreg 

fabrication route, although the in-plane fibers are effectively straight, total fiber contents are not high 

(30 percent) due to the need to individually coat almost every fiber with a thick interphase. Thus all 

fiber-controlled properties are less than optimum for the RMI route. 

Regarding CMC thru-thickness properties, it is found that as the multiply 2D-woven fabric of for the 

NRMI route is compressed in tooling prior to interphase deposition, the resulting CMC thru-thickness 

tensile strength and thermal conductivity will generally increase. This effect is primarily due to increased 

fiber-fiber contact and intermingling or “nesting” of the fiber tows. But as might be expected, this nesting 

results in less opportunity to achieve uniform CVI-produced interfacial coatings on all the fiber surfaces 

in a tow or thru the thickness, plus increased mechanical degradation of the fibers and CMC in-plane 

strength properties. Whereas the NRMI route relies on increasing fabric nesting to improve the CMC 

thru-thickness or Z-direction properties, the RMI route relies on its low total fiber content to allow the 

dense SiC-based matrix between the fibers to provide monolithic-like thru-thickness tensile strength. For 

both the NRMI and RMI routes, it is difficult to obtain reproducible CMC thru-thickness properties due to 

local microstructural variances. This results in large scatter, for example, in the thru-thickness or 

Z-direction properties, which in turn results in significantly low thru-thickness tensile strength allowables 

(20 MPa) for current SiC/SiC components. Finally, both fabrication routes rely on the lay-up of finitesized 2D plies, which means that for components with continuous surfaces, such as combustor liners and 

airfoils, there must be regions where the plies terminate and fiber reinforcement is absent. Thus unless the 

ply ends are intermingled in some judicious manner, or the ply ends are located in low stress regions, 

there can be a significant knockdown in component strength properties at these locations (Ref. 33). 

Realizing the general architecture goals and the deficiencies of the two MI fabrication routes, NASA 

in recent years has been examining SiC/SiC CMC panels primarily consisting of multiple types of 3Dwoven pre-forms with thru-thickness fiber reinforcement. These types range from simple 2.5D-woven to 

more complex 3D-orthogonal architectures (Refs. 34 and 37). In most cases, iBN SiC fiber tows have 

been used as the in-plane warp and fill stuffers, and the interfacial coatings and matrices were formed by 

those used in the NRMI fabrication route. For the warp weaver or thru-thickness fibers, due to their high 

stiffness and tendency to fracture under high bend angles, iBN fibers were only used for preforms with 

low-angle thru-thickness tows; whereas low-stiffness low-performance SiC-based “ZMI” fibers were used 

for the high-angle 3D-orthogonal pre-forms. Table 3 lists the key room temperature in-plane and thruthickness properties for some 2.5D and 3D SiC/SiC panels and compares them with panels fabricated by 

the standard NRMI route using the lay-up of 2D-woven iBN fabric (NASA N24A SiC/SiC). For the 

Table 3 panels, the in-plane iBN stuffer tows were nearly straight with total fiber contents of ~36 percent, 

and were nearly balanced at ~18 percent in each of the warp and fill directions or were un-balanced at 

~10 percent in the warp and ~25 percent in the fill direction. The indicated test data were measured inplane in the fill or Y direction, as well as in the thru-thickness or Z direction. 

 
NASA/TM—2013-217889 12 

TABLE 3.—TYPICAL DIRECTIONAL STRENGTH AND THERMAL CONDUCTIVITY PROPERTIES OF 

SiC/SiC PANELS WITH VARIOUS WOVEN FIBER ARCHITECTURESa

SiC/SiC 

Fiber architecture 

Fiber content, 

percent 

20 C CMC directional 

strengths, 

MPA 

CMC directional thermal 

conductivity, 

W/m.K Ref. 

Y Z X 

20 C 

Z 

20 C 

Z 1400 C 

X Y Z MCS UTS UTS 

A 2D fabric 18 18 0 180 460 15 50 25 18 6, 32 

B 2.5D 17 19 3 140 360 28 -- 50 25 35, 36 

C 3D Orthogonal 18 17 5 120 320 -- -- -- -- 37 

D 3D Orth-Unbal 10 25 5 240 600 -- -- -- -- 37 a

X, Y stuffer fibers = Sylramic-iBN fiber; Matrix fabrication = CVI-MI; 

2.5D Z weaver = Sylramic-iBN fiber; 3D Z weaver = ZMI SiC Fiber 

Table 3 shows that for the NRMI fabrication route, significant advances can be achieved in SiC/SiC 

CMC structural and thermal properties by the use of 3D-woven pre-forms rather than by the current 2Dwoven fabric lay-up approach. For example, when the 2D-woven Panel A is compared against the 2.5Dwoven Panel B, although a small knock-down occurred for the in-plane strengths, one can see a doubling 

in the Z-direction properties of thru-thickness tensile strength (Z-UTS) and thermal conductivity (Z-TC) 

at room temperature, but also a significant increase in Z-TC at 1400 C (extrapolated). Indeed this latter 

Z-TC value of 25 W/m.K may be the highest ever measured at high temperature for SiC/SiC panels. 

These improvements were achieved with only ~3 percent iBN fibers in the Z direction, so that further 

improvements should be expected with higher Z-fiber content. Whereas the Z-UTS improvement can be 

explained by the use of strong iBN fibers, the improvement in Z-TC was probably due in part to the high 

conductivity of these fibers and in part to the fact that weaver fibers probably allowed better infiltration 

channels for the NRMI matrix. 

Then when one compares Panel A with the balanced 3D-orthogonal Panel C, one finds a further 

decrease in in-plane strengths, even though the Y-stuffer fibers in Panel C are straighter than those in 

Panel A. This degradation in the warp and fill stuffers of Panel C may have been caused by excessive 

abrasion with the ZMI weaver fibers and/or to reaction with these fibers during high-temperature CMC 

fabrication. Finally, when one compares the balanced 3D-orthogonal Panel C with the un-balanced 3Dorthogonal Panel D, one sees an approximate doubling of the in-plane Y-direction strengths, even though 

the fiber content in this direction only increased from ~18 to 25 percent. Thus, although these results are 

preliminary, one might conclude that the demonstrated Z-direction property benefits of thru-thickness 

fibers need not necessarily exclude improvement at the same time of the in-plane properties, provided one 

makes a judicious choice of the SiC fiber and its 3D architecture. 

Beyond improved thru-thickness strength and conductivity, NASA has shown that the 2.5D-woven 

architectures with iBN fibers also provide SiC/SiC CMC with improved delamination resistance and 

impact resistance in comparison to 2D architectures (Ref. 38). As discussed above, another potential 

advantage for 3D architectures can be enhanced matrix infiltration down the weaver fibers, resulting in 

reduced matrix porosity. Furthermore, the possibility may exist that the automation aspects of 3D 

weaving or braiding will result in reduced manufacturing costs and reduced scatter in composite 

properties by eliminating the human element that is typically involved in 2D fabric and tape lay-up 

techniques. However, for the highest temperature SiC/SiC components, sintered SiC fibers, such as the 

iBN and Tyranno SA fibers, are near-stoichiometric in composition with large nano-sized grains, 

resulting in high bending stiffness and rough fiber surfaces leading to poor abrasion resistance. As 

described for the panels in Table 3, these issues can enhance the probability of fiber fracture during 3D 

pre-forming and thus limit the available 3D-woven architectural designs and thru-thickness fiber 

fractions. Nevertheless, NASA has recently demonstrated with a commercial textile processing vendor 

that complex 3D-woven pre-forms for SiC/SiC panels as well as for SiC/SiC turbine airfoils can be 

fabricated using polymer-based “serving” fiber tows that are wound around sintered SiC fiber tows prior 

to weaving, thereby significantly reducing their abrasion issues (Ref. 39). In subsequent CMC processing, 

these fibers will decompose leaving a negligible carbon char, which should be of little issue for the key 
NASA/TM—2013-217889 13 

CMC properties. Another potential approach developed at NASA for complex-shaped components is to 

first form 3D pre-forms with simpler shapes that are close to those required by the component, but do not 

contain the extreme bends that can cause fiber fracture. Then one can heat treat the architecture while 

applying continuous or discrete levels of pressure to creep-form the fibers without fracturing them into 

the tighter bends required by the component shape. NASA has demonstrated that this can be done at the 

same time as the iBN treatment with minimal loss in fiber strength (Ref. 25). 

4.4 Advances in SiC-Based Matrices 

Since the free silicon in the NRMI and RMI matrices limits their long-term SiC/SiC CMC upper use 

temperatures to less than 1300 C, NASA and the U.S. Air Force have recently been researching and 

developing SiC/SiC systems by using alternate matrix infiltration processes into stacked 2D-fabric preforms that result in SiC compositions that are effectively silicon-free. To date, three such infiltration 

processes have been extensively evaluated (Ref. 34): (1) CVI processing where SiC deposition is 

continued as long as possible to fill the porosity within an interface-coated 3D pre-form (Full CVI 

approach); (2) polymer-infiltration and pyrolysis (Full PIP approach) in which multiple infiltration cycles 

are performed, each consisting of infiltration of a SiC-yielding polymer into the remaining 3D pre-form 

porosity, followed by polymer decomposition at high temperatures into crystalline SiC particulates; and 

(3) a hybrid of these two processes, where typically the CVI approach is first used for partial filling of the 

pre-form porosity, and then the remaining porosity is filled as much as possible by subsequent multiple 

PIP cycles (CVI-PIP approach). Although achieving the goal of an effectively silicon-free SiC matrix, 

each of these three processes have various advantages and issues that may or may not limit their 

application for higher temperature SiC/SiC components. 

In terms of advantages, the Full CVI method for matrix infiltration is amenable to the fabrication of 

SiC/SiC components having irregular, large, and complex shapes with near-net shape capability. High 

temperature properties are generally good because high purity matrices with controlled microstructures 

can be obtained. Aside from trapped porosity (10 to 15 vol%), this approach typically results in a SiC 

matrix material with high thermal conductivity and high creep resistance in multiple directions. Also CVI 

is carried out at relatively low temperatures (1000 C), so that there is little risk of fiber damage. In 

addition, the fiber-matrix interphase can be tailored by pre-coating the fiber preform during an initial CVI 

step using different reagents, but the same reactor. However, one of the major limitations of the Full CVI 

method is the long deposition times needed for the densification process (typically greater than 100 hr). 

This raises the cost of the component fabrication in comparison to processes that require much shorter 

times to complete the matrix infiltration. Furthermore, although the CVI matrix initially forms fairly 

uniformly within the fiber tows and then on the surface of the tows, as deposition continues nonuniform 

regions of matrix with large and sharp pores typically form between the tows due to their separation and 

due to their depth within the CMC thickness. Because of their size and irregular shapes, these trapped 

pores reduce the CMC MCS in multiple directions. Thus it is not possible to obtain fully dense matrices 

using the Full CVI approach. In addition, the typical presence of residual closed porosity may shorten the 

useful life of the CMC component particularly after matrix cracking in an oxidizing environment. Finally, 

in relation to fully dense CMC, the trapped or closed porosity will not allow CMC with maximum 

thermal conductivity. 

Regarding the Full PIP approach, its main advantage is its use of methods typically employed and 

well established in the fabrication of polymer matrix composites. Thus in contrast to the Full CVI 

approach, the size of CMC components is not limited by the size of the CVI reactor. However, one of its 

main drawbacks is the long processing times and repeated infiltration/pyrolysis cycles in order to obtain 

sufficiently high matrix density. Due to polymer shrinkage during pyrolysis, another drawback is that 

matrix micro-cracking generally results within the CMC in the form of fine cracks between the 

crystallized particulates. Re-infiltration fills these cracks to some degree, but matrix cracking strength and 

thermal conductivity in multiple directions generally remain low in comparison with the Full CVI 

approach. Finally, due to oxygen in the polymer and also multiple cycles from room to high temperature, 
NASA/TM—2013-217889 14 

oxygen is typically picked up within the fine cracks, which then can attack the interphase and fibers 

unless oxygen getters are also introduced into the matrix (Ref. 34). 

For recent research and development studies aimed at silicon-free matrices, NASA has been 

emphasizing the CVI-PIP hybrid approach because it avoids to a large degree the drawbacks of the Full 

CVI and Full PIP approaches while incorporating some of their advantages. For example, just as in the 

NRMI fabrication route, only partial CVI deposition is utilized, thereby affording much shorter deposition 

times and reduced risk of large inter-tow pores, while retaining a partial SiC matrix with high strength, 

high creep-resistance, and high thermal conductivity in multiple directions. In addition, the partial CVI 

SiC matrix provides environmental protection for the interphase and fibers during subsequent matrix 

infiltration processes. Filling the remaining matrix by the PIP approach allows the introduction of high 

conductivity SiC particulates without the concern of oxygen pick-up during multiple PIP cycles. 

Although fine cracks will still exist between the particulates, they are much finer than those that form 

during cracking of the CVI matrix so that they can seal much more rapidly in high-temperature oxidizing 

environment. Nevertheless, although the hybrid approach has these multiple advantages, it like the other 

approaches does not as yet yield matrices with zero or near-zero porosity. 

Table 4 compares some key in-plane and thru-thickness properties at 20 and 1450 C for SiC/SiC 

CMC panels fabricated with matrices infiltrated by the CVI-MI, Full CVI, Full PIP, and hybrid CVI-PIP 

approaches. The first set of panels employed 2D-woven iBN fabric; while the second set employed 2.5D 

woven iBN tows. To achieve maximum performance in terms of reduced matrix defects and increased 

matrix creep resistance and thermal conductivity, all the panels with silicon-free matrices were postprocessed at NASA by annealing treatments above 1500 C (Refs. 6 and 35). Examining Table 4, one can 

see that currently all the indicated processes for Si-free matrices do not achieve the low ~5 percent 

porosity of the CVI-MI approach. The high porosity levels of ~15 percent typically result in lower matrix 

cracking strengths depending on fiber content and the character of the porosity, and in variable ultimate 

strengths, also depending on fiber content and possibly on the resistance of the iBN fibers to attack during 

CMC processing and post treatment. One approach to acquire better information about the matrix 

structural quality is to test the CMC in an off-axis direction where the matrix carries more of the CMC 

load. Thus when tested at 45 to the balanced X and Y directions, Table 4 shows a significant knockdown 

in the X/Y-UTS for the Full PIP panel, indicating that this matrix approach does not have the structural 

quality provided by the partial CVI or Full CVI approaches. On the other hand, the Full PIP displays a 

relatively high Z-UTS. As previously described, this is caused primarily by increased fabric compression 

and nesting of the fiber tows between fabric plies. Furthermore, due to the fine cracks, the Full PIP panel 

does not compare in thermal conductivity with the other matrix approaches. This low Z-TC property is 

probably due in part to the need for boron-based oxygen getters purposely incorporated in the matrix to 

avoid the effects of oxygen pickup during the multiple PIP cycles (Ref. 34). 

TABLE 4.—TYPICAL EFFECTS OF VARIOUS MATRIX INFILTRATION APPROACHES ON DIRECTIONAL 

STRENGTH AND THERMAL CONDUCTIVITY PROPERTIES OF SiC/SiC PANELSa

SiC/SiC 

Matrix 

approach 

Y 

Fiber, 

percent 

Matrix 20 C CMC directional strengths, 

MPa 

Z thermal 

conductivity, 

W/m.K 

Ref. 

~Free 

Silicon, 

percent 

~Total 

porosity, 

percent 

Y 

MCS 

Y 

UTS 

XY 

UTS 

Z 

UTS 

20 C 1400 C 

 2D-woven Sylramic-iBN fabric 

CVI-MI 18 13 5 180 460 240 15 25 18 6, 32 

Full CVI 18 0 15 120 350 --- 7 28 8 34, 35 

Full PIP 21 0 <15 140 400 100 23 8 5 34, 32 

CVI-PIP 18 0 14 150 360 --- 10 28 10 34, 35 

 2.5D-woven Sylramic-iBN Tow 

CVI-MI 18 13 5 140 360 --- 28 50 25 36 

CVI-PIP 18 0 14 130 380 --- 24 45 22 36 a

All silicon-free SiC/SiC CMC were post treated above 1500 C 
NASA/TM—2013-217889 15 

Regarding the hybrid CVI-PIP approach, although it can provide SiC/SiC that can operate at higher 

temperatures, Table 4 shows that with conventional 2D-woven architectures, it does not as yet provide 

comparable properties with CMC fabricated with the CVI-MI matrix approach. As might be expected 

from their high residual porosity, the 2D-woven CVI-PIP panels show a small knockdown in Y-MCS and 

Z-UTS, as well as high temperature Z-TC, but their values are still higher than those of the Full CVI 

panels. Nevertheless, as also shown in Table 4, when the hybrid matrix approach is combined with a 

2.5D-woven architecture, the resulting panels, while still retaining a small deficient in in-plane properties, 

display significant improvement in Z-UTS and Z-TC in comparison to those fabricated by the 2D-woven 

route and CVI-MI matrix. Thus if the porosity can be reduced and advanced 3D-architectures with iBN 

fibers are employed, the hybrid CVI-PIP matrix approach has the potential of not only displaying better 

structural and thermal properties than the current NRMI or CVI-MI route for SiC/SiC components, but 

also higher temperature structural capability as discussed in the following section. 

4.5 Advances in SiC/SiC Microstructural Design Methods 

As discussed above, of key importance for SiC/SiC component implementation in high-temperature 

applications is the ability to select the optimum constituents and their geometries so that the component 

microstructure has the intrinsic structural capability to achieve the required component life under its 

projected mechanical, aerodynamic, and thermal loads. For this selection process, a method often used for 

metals and ceramics is to perform stress-rupture tests in which a constant stress is applied to a material at 

a constant high temperature under laboratory conditions. The time for total material rupture is then 

recorded for various combinations of stress and temperature. One simple and convenient empirical 

approach for representing the data from these tests is by use of “Larson-Miller” (LM) plots. In these plots, 

master curves are generated for various materials, which describe the applied stress at rupture (material 

rupture strength) versus the time-temperature dependent parameter q given by 

q Tlog t D  r  (3) 

Here T (Kelvin) is the absolute temperature for the rupture test, tr (hours) is the material rupture time, 

and D is a constant that best allows the rupture data at different temperatures to fall nearly on one master 

curve. The designer can then compare the rupture test results for prospective materials for an initial downselect of those materials that can best meet the component’s service requirements. NASA has determined 

that this LM approach with D = 22 can be effective method for SiC-based fibers (Ref. 29) and also for 

SiC/SiC CMC (Ref. 40), not only for showing the effects of various constituents and architectures on the 

structural life of various SiC/SiC CMC, but also for demonstrating the benefits of the recent constituent 

advances discussed above. 

In Figure 2, the LM approach has been used to plot the in-plane rupture strength data for various 

SiC/SiC composites versus the time-temperature at which they failed by complete rupture (Ref. 40). All 

CMC were fabricated as panels with by the NRMI route and tested on-axis with one of the primary fiber 

directions (content of ~18 percent), and were reinforced with balanced 2D-woven 0/90-balanced fiveharness satin (5HS) fabrics of different SiC fiber types as indicated by the different curves. Also included 

is the predicted curve if an NRMI CMC were fabricated in the same manner using the developmental 

Super Sylramic-iBN fiber. Thus all CMC contained free silicon so that their long term creep and rupture 

behaviors were effectively controlled by the reinforcing fiber tows under the total CMC load. In Figure 2, 

one can see that the LM approach appears to be a good empirical method for predicting SiC/SiC rupture 

life in that the rupture strength data for various fiber types not only fall consistently near one best fit curve 

for each type, but also the curves closely follow the creep ranking of the different SiC fiber types as 

measured on single fibers in Figure 1. Although the data for each fiber type show some scatter from its 

best fit curve, the curves can still be clearly distinguished from one another. It is envisioned that some of 

data scatter for a specific fiber type may be due to whether or not matrix cracking occurred during initial 

CMC loading. Also it should be realized that with the LM approach, the scatter in rupture time data is 
NASA/TM—2013-217889 16 

Figure 2.—Effects of SiC fiber type on on-axis rupture strength in air for SiC/SiC CMC with 2D-woven 

0/90-balanced 5HS fabric, on-axis fiber content of ~18 percent, and CVI-MI matrix (Ref. 40). 

greatly reduced by the time insensitivity of the LM (q) parameter. Thus the curves of Figure 2 are 

practically useful for understanding the projected rupture life of an NRMI-fabricated SiC/SiC component 

that needs to experience a maximum stress at a maximum temperature. For example, if the maximum onaxis stress for an un-cracked woven iBN/SiC component was 100 MPa, Figure 2 predicts a maximum (q)

value of ~39400, which in turn would allow an intrinsic rupture life for the component of 100 hr at 

1642 K (1369 C), or 1000 hr at 1576 K (1303 C).

Besides its practical implications, there are also various mechanistic conclusions one can draw from 

the Figure 2 results. For example, some of the data were taken at high enough CMC stress to initially 

crack the matrix, and yet the rupture strength data fall very close to the data where the CMC stresses did 

not cause initial matrix cracking. Thus, at least in laboratory air, the rupture mechanism for NRMI CMC 

appears to not be related to environmental attack, but to the fiber tows reaching a certain ultimate rupture 

strain during creep. For example, no matter the test conditions, the HNS CMC typically ruptures at 

~0.6 percent total strain; whereas the iBN CMC ruptures at ~0.4 percent (Ref. 41). Thus in comparison to 

the iBN CMC, the slightly better rupture behavior of the HNS CMC in Figure 2 can be explained by its 

enhanced rupture strain, which in turn is probably related to its smaller grain size. It is also interesting to 

note that the rupture strains of the two fiber types are effectively equivalent to the ultimate fracture strains 

for their CMC in a room temperature fast-fracture test (Ref. 17).

The LM approach has also proved useful in understanding fiber architecture effects. For example, 

Figure 3 plots rupture strength data versus (q) data for a variety of different iBN-reinforced architectures 

within SiC/SiC panels fabricated with CVI-MI matrices (Ref. 40). The first set of data on the left 

reproduces the data in Figure 2 for NRMI-fabricated CMC reinforced by 2D-woven 5-harness satin fabric 

of iBN tows with on-axis content of ~18 percent. This curve is then projected to the right by assuming 

that if the same woven architecture could be produced with ~25 percent on-axis iBN fiber content, the 

final stress on the fiber would be reduced by a factor of 18/25 for the same CMC applied stress. Finally, 

40

50

30

20

10
NASA/TM—2013-217889 17 

Figure 3.—Effects of fiber architecture on on-axis matrix cracking and rupture strength in air for SiC/SiC 

CMC with Sylramic-iBN woven tows and CVI-MI matrix (Ref. 40). 

in the upper center of Figure 3 are the CMC rupture data (Ref. 41) measured in the fill direction for some 

of the 2.5D and 3D architectures described in Table 3. As previously indicated, in the fill direction the 

iBN tows were straight with an average fiber content of ~25 percent. When one compares this data set 

with the first set, one can see a significant improvement in CMC rupture strength for the same timetemperature test conditions. As indicated, this improvement is still much better than the middle curve 

which has the same fiber content, but incorporates bent fibers as a result of the woven fabric. Bottom line 

from Figure 3 is that the structural life of SiC/SiC components can be significantly improved by (1) 

increasing the fiber content in the key component directions and (2) assuring as much as possible that the 

fibers are in their as-produced straight condition. 

Finally, one can use the LM approach in Figure 4 to understand the effects of different matrix 

infiltration approaches on SiC/SiC rupture behavior (Refs. 40 and 42). Here all CMC are again fabricated 

with 2D-woven 5-harness satin fabric of iBN tows, and the first data set on the left were fabricated by 

CVI-MI or the NRMI route. The next curve at higher (q) values is for CMC with Full PIP matrices, but 

with ~25 percent fiber content due enhanced fabric compression after lay-up. This data set is almost 

exactly 25/18 times the CVI-MI data set indicating that although containing no free silicon, the PIP 

matrix carries no long term load, leaving the fibers to again carry the entire CMC load. This effect can be 

attributed to little bonding between the PIP-generated SiC particulates caused by particle shrinkage during 

pyrolysis. Thus as previously discussed, the Full PIP matrix approach picks up its main advantage of 

higher fiber content in plane, but suffers in other key CMC properties. On the other hand, the postannealed CMC with Full CVI matrix and ~18 percent fiber content does show improved behavior over the 

silicon-containing CVI-MI matrix with the same fiber content, suggesting that this creep-resistant matrix 

can carry some of the CMC load at temperatures where the CVI-MI and Full PIP matrices cannot. Finally, 

Figure 4 shows that although the post-annealed hybrid CVI-PIP approach does not have the CVI SiC 

content of the Full PIP, this could be an advantage by allowing more load to be put on the CVI SiC, 
NASA/TM—2013-217889 18 

Figure 4.—Effects of matrix infiltration approach on on-axis rupture strength in air for SiC/SiC CMC 

with 2D-woven 0/90-balanced Sylramic-iBN fabric (Refs. 40 and 42). 

causing it to creep at a rate closer to that of the fiber. In effect, this would be similar to increasing the 

fiber content. Thus for balanced 2D fiber architectures and post-annealed Si-free matrices, the hybrid 

CVI-PIP matrix approach currently yields the best high temperature SiC/SiC rupture behavior when 

compared to any other current approach (Ref. 42). 

5.0 Current Microstructural Design Guidelines and Potential Service Issues 

for Higher Temperature SiC/SiC Components 

The advancements described in this chapter offer a better theoretical and experimental understanding 

of the effects of various SiC-based constituent compositions and geometries on the key mechanical and 

thermal properties of SiC/SiC CMC. As such, they suggest the following general microstructural design 

guidelines for achieving a SiC/SiC component with high structural reliability and higher temperature 

capability than current SiC/SiC fabrication routes: 

 Select (or develop) a polycrystalline small-diameter near-stoichiometric SiC fiber with average grain 

size between 200 to 500 nm, with grain boundaries free of such impurity elements as free silicon, free 

carbon, oxygen, boron, aluminum, and other elements that act as SiC sintering aids, and with a crosssectional microstructure as uniform as possible in grain size and composition. 

 Select (or develop) a fiber interfacial coating system that is compliant enough to deflect thruthickness matrix cracks before reaching the fiber, is based on a composition that does not volatilize in 

an oxidative environment like carbon, and fully covers every fiber surface such as the iBN coating. 

 Select, if possible, a fiber architecture that allows the final component shape to be textile formed, 

maximizes total fiber content in the component, allows sufficient fiber content in the key component 

stress directions, and eliminates or minimizes fiber curvature (straight fibers) in these directions. 

 Select (or develop) a matrix process that yields a near-dense Si-free, tough, and SiC-based matrix that 

is highly thermally conductive and creeps slightly more than the fiber. 
NASA/TM—2013-217889 19 

Although these guidelines may provide a SiC/SiC component with constituents that are intrinsically 

able to withstand quasi-static structural loads under uniform stress conditions, there are certain situations 

that can arise during component service that may still limit its performance and life and thus require 

further investigation. As discussed in Section 2.0, one such situation is related to creep in that, even if the 

creep mismatch between the fiber and matrix can be minimized, there will certainly be stress and 

temperature gradients within in the component leading to local amounts of differential creep strain, which 

in turn will result in local residual stress development that is not relieved at lower temperatures where the 

CMC behaves elastically (Ref. 5). These residual stresses can in turn reduce the local matrix cracking 

strength, so that during some follow-up situation, the weakened area could be unpredictably subject to a 

tensile stress that could exceed the local cracking strength. Hopefully the resulting cracking or damage 

will not alter affect the component performance or allow the service environment to enter and degrade the 

CMC microstructure. Clearly one approach for avoiding this situation is to only employ the component at 

temperatures and times where the CMC exhibits little or no creep. Thus if a certain high temperature 

performance is required for a SiC/SiC component, one should seek SiC fibers and matrices that display 

minimum long term creep at this temperature. 

Another possible life-limiting issue is flaw or crack creation in the SiC/SiC component generated by 

unexpected mechanical and/or thermal shock. Mechanical shock, for example by foreign object impact, 

cannot only occur during component service, but also during general handling and assembly of the 

component. Currently, without sufficient engine experience, it is very difficult to predict the flaw severity 

that can arise from this effect or its degrading influence on component residual strength and structural 

life. For this reason, the general design approach should be to only seek SiC-based matrices that in their 

as-produced condition display high fracture toughness or high resistance to crack propagation from 

surface flaws, so that the as-produced matrix cracking strength of the as-produced component will be 

minimally compromised during service. Similar to mechanical shock, thermal shock can also lead to 

cracking, for example in a component wall as the colder side experiences an unexpectedly high tensile 

stress that is proportional to the CMC thru-thickness thermal conductivity. If cracking should occur under 

either shock situations, the reinforcing fibers will bridge the cracks providing they have sufficient tensile 

strength and volume content to withstand the mechanical over-load. Thus for each engine component, it is 

important to understand the probability of shock, and where it might occur in order to assure that the 

CMC will have sufficient overload capability in these locations to maintain component integrity and 

allow continuous engine operation. As discussed in Section 4.3, the use of 3D fiber architectures not only 

improves the mechanical shock or impact resistance of SiC/SiC panels, but also can improve the panel 

thru-thickness thermal conductivity, thereby reducing the stresses from any thermal shock situation. 

Finally, a predictable and very important life-limiting issue for SiC/SiC components is erosion and/or 

corrosion of the component surface during high-temperature gas-turbine service. This effect, which can 

be severe for all silicon-based ceramics, is typically due to the water vapor formed in the combustion 

process, which reacts with the oxides, such as silica, that continually form on the ceramic surface, causing 

them to volatilize. This in turn results in surface recession of the ceramic at a rate that increases with 

water vapor pressure, combustion gas velocity, and most importantly with CMC surface temperature 

(Ref. 43). To minimize this effect, environmental barrier coatings (EBC) have been developed for both 

Si-based monolithics and composites (Refs. 44 and 45). For industrial gas turbines, these EBC have been 

very instrumental in extending the life of SiC/SiC CMC combustor liners and shrouds at temperatures up 

to 1250 C (Ref. 46), but little experience exists for using EBC under more aggressive engine 

conditions. Because EBC themselves are typically produced from oxide compositions, such as barium 

aluminosilicate and mullite, they generally display very low thermal conductivity, thereby also acting as 

thermal barrier coatings (TBC). Thus EBC are not only important for maintaining ceramic component 

cross-section, they can also aid in reducing the surface temperature of a cooled ceramic component by 

providing another source of thermal resistance between the combustion gas temperature and the 

component surface. Generally EBC are required for Si-based ceramic components that will require long- 

 
NASA/TM—2013-217889 20 

time service and surface temperatures above ~1000 C. As efforts are ongoing to increase the upper use 

temperatures of SiC/SiC components beyond 1250 C, recession becomes the primary concern for long 

component life. Thus high-temperature EBC development is a key objective of current studies at NASA 

(Ref. 47). 

6.0 Concluding Remarks 

This chapter has detailed the results of various studies conducted primarily at the NASA Glenn 

Research Center aimed at advancing the current state of experimental and mechanistic knowledge 

concerning the constituent materials and processes for SiC/SiC composites. When compared to current 

approaches for SiC/SiC aero-propulsion components, this knowledge has the potential of generating 

components with enhanced structural reliability and higher temperature performance. It should be realized 

that these advancements have been demonstrated primarily using single SiC fibers, woven SiC/SiC preforms, and SiC/SiC test coupons removed from panels that contain the improved approaches. These 

specimens were then tested under simple conditions such as laboratory air, constant temperature, and 

constant stress along the primary fiber direction. Thus considerable more effort is required before these 

advancements can be comfortably implemented into real SiC/SiC engine components. Nevertheless, when 

compared to current approaches for SiC/SiC components, they clearly show substantial improvements, 

particularly for achieving the multiple engine benefits that can be accrued by components with durable 

higher temperature performance. Because of these benefits, efforts at NASA are on-going to further 

understand and demonstrate the practical viability of these advancements. 

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Cylinders Fabricated from Spliced 2D Fabric, Cer. Eng. Sci. Proc., 24 [4], 407–413 (2003). 
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34. H.M. Yun, J.A. DiCarlo, and R.T Bhatt, Advanced SiC/SiC Ceramic Composites for Air-Breathing 

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14. ABSTRACT

In the last decade, considerable progress has been made in the development and application of ceramic matrix composites consisting of 

silicon carbide (SiC) based matrices reinforced by small-diameter continuous-length SiC-based fibers. For example, these SiC/SiC 

composites are now in the early stages of implementation into hot-section components of civil aero-propulsion gas turbine engines, where in 

comparison to current metallic components they offer multiple advantages due to their lighter weight and higher temperature structural 

capability. For current production-ready SiC/SiC, this temperature capability for long time structural applications is ~1250 °C, which is 

better than ~1100 °C for the best metallic superalloys. Foreseeing that even higher structural reliability and temperature capability would 

continue to increase the advantages of SiC/SiC composites, progress in recent years has also been made at NASA toward improving the 

properties of SiC/SiC composites by optimizing the various constituent materials and geometries within composite microstructures. The 

primary objective of this chapter is to detail this latter progress, both fundamentally and practically, with particular emphasis on recent 

advancements in the materials and processes for the fiber, fiber coating, fiber architecture, and matrix, and in the design methods for 

incorporating these constituents into SiC/SiC microstructures with improved thermo-structural performance. 

15. SUBJECT TERMS

Ceramic composites; SiC fibers; Fiber architectures; Fiber coatings; SiC matrices; SiC/SiC design 

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