S.V. Raj

Glenn Research Center, Cleveland, Ohio

M. Singh and R. Bhatt

Ohio Aerospace Institute, Brook Park, Ohio

High Temperature Lightweight Self-Healing 

Ceramic Composites for Aircraft 

Engine Applications 

NASA/TM—2014-218352

October 2014
	
	
	








	



	

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S.V. Raj

Glenn Research Center, Cleveland, Ohio

M. Singh and R. Bhatt

Ohio Aerospace Institute, Brook Park, Ohio

High Temperature Lightweight Self-Healing 

Ceramic Composites for Aircraft 

Engine Applications 

NASA/TM—2014-218352

October 2014


	



	


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NASA/TM—2014-218352 1 

High Temperature Lightweight Self-Healing 

Ceramic Composites for Aircraft Engine Applications 

S.V. Raj 

National Aeronautics and Space Administration 

Glenn Research Center 

Cleveland, Ohio 44135 

 

M. Singh and R.T. Bhatt 

Ohio Aerospace Institute 

Brook Park, Ohio 44142 

Abstract 

The present research effort was undertaken to develop a new generation of SiC fiber-reinforced 

engineered matrix composites (EMCs) with sufficient high temperature plasticity to reduce crack 

propagation and self-healing capabilities to fill surface-connected cracks to prevent the oxygen ingress to 

the fibers. A matrix engineered with these capabilities is expected to increase the load bearing 

capabilities of SiC/SiC CMCs at high temperatures. Several matrix compositions were designed to match 

the coefficient of thermal expansion (CTE) of the SiC fibers using a rule of mixture (ROM) approach. 

The CTE values of these matrices were determined and it was demonstrated that they were generally in 

good agreement with that of monolithic SiC between room temperature and 1525 K. The parameters to 

hot press the powders were optimized, and specimens were fabricated for determining bend strength, 

CTE, oxidation and microstructural characteristics of the engineered matrices. The oxidation tests 

revealed that some of the matrices exhibited catastrophic oxidation, and therefore, these were eliminated 

from further consideration. Two promising compositions were down selected based on these results for 

further development. Four-point bend tests were conducted on these two promising matrices between 

room temperature and 1698 K. Although theses matrices were brittle and failed at low stresses at room 

temperature, they exhibited high temperature ductility and higher stresses at the higher temperatures. The 

effects of different additives on the self-healing capabilities of these matrices were investigated. The 

results of preliminary studies conducted to slurry and melt infiltration trials with CrSi2 are described. 

1.0 Introduction 

Nickel-based superalloys have been successfully used in aircraft engines as blades, combustor liners 

and vanes for several decades [1]. The use of reliable, high temperature, lightweight materials in the 

manufacture of aircraft engines is expected to result in lower fossil and bio-fuel consumption, thereby 

leading to cost savings and lower carbon emissions due to air travel. Designing these hot sections of gas 

turbine engines with lighter weight silicon carbide fiber-reinforced silicon carbide (SiC/SiC) ceramic 

matrix composites (CMCs) allows them to operate at higher temperatures and higher pressure ratios with 

reduced cooling air while leading to lower fuel burn and reduced CO2 emissions [2,3,4]. As a result, a 

considerable amount of research has been conducted over the last three decades to develop SiC/SiC 

CMC technology for aerospace applications [5,6,7,8]. Thus, the technology has maturated sufficiently 

that most major aircraft engine companies in the United States and Europe plan to introduce SiC/SiC 

CMCs in their next generation aircraft engines [9]. 

Currently, there are four broad methods for fabricating silicon carbide-based CMCs depending on the 

matrix processing methods: (a) Melt infiltration (MI), (b) Chemical Vapor Infiltration(CVI), (c) Polymer 

Infiltration and Pyrolysis (PIP), and (d) Hot Pressing (HP) of SiC powders with additives [8]. In all CMC 

manufacturing, the starting material is a one-dimensional (1-D) SiC fiber tow, a two-dimensional (2-D) 

woven cloth, or three-dimensional (3-D) woven fiber preforms. The starting material is then coated with 
NASA/TM—2014-218352 2 

single or multi-layered interface coatings, such as BN, to reduce strength degradation of the fibers during 

composite fabrication, as well as to provide a crack deflection path at or along the interface during overloading of the composite. A SiC or Si3N4 layer is often deposited on these fiber coatings by CVI to protect 

these coatings from reacting with either molten silicon for MI-CMCs or SiC producing polymers for PIPCMCs. 

1.1 Limitations of the Current Technology 

The present generation SiC/SiC CMCs possess a SiC matrix which results in the CMCs possessing a 

relatively low proportional limit, PL, due to the through-thickness cracking of the brittle matrix thereby 

limiting the design stresses, D, to much less than PL [5,6]. Although the SiC matrix fulfills a useful role 

of filling space between the SiC fiber tows in the woven fabric or preform to provide a thermally 

conductive path for heat transfer, these CMCs rely almost entirely on the SiC fibers to carry the load 

owing to the premature cracking of the matrix during loading. Thus, the high temperature usefulness of 

traditional non-oxide CMCs falls well below their theoretical capabilities. Fracture toughness 

improvement over monolithic SiC is achieved by crack bridging and crack deflection mechanisms due to 

the fibers [10]. Above the matrix cracking stress, these composites show poor durability because of 

oxygen ingression from the external surfaces through interconnected matrix cracks to the fiber/matrix 

interface thereby leading to the oxidation of the boron nitride interfacial coating and premature strength 

degradation of the reinforcing fibers [6,11]. Therefore, the brittleness of the matrix severely constraints 

the design of CMCs for highly stressed components, such as high-pressure turbine blades. It is noted that 

in practical applications, CMCs will be protected by several layers of environmental barrier coatings 

(EBCs) so that the useful life of traditional CMCs are greatly dependent on the life of these coatings and 

the matrix cracking strength. 

Irrespective of the other manufacturing details, the final step of MI SiC/SiC CMCs traditionally 

involves infiltrating the space between the SiC fiber tows with molten silicon [12]. Although CMCs 

fabricated by this technique are denser than those prepared by non-MI methods, the presence of 

unreacted silicon in the matrix restricts their use to below 1588 K (2400 °F) due to its low creep strength 

and high reactivity above this temperature. Even when the CMCs are fabricated entirely by CVI, the 

presence of unreacted silicon in the grain boundaries of the CVI deposited SiC can react with the 

protective boron nitride coating on the SiC fibers. One of the objectives of NASA’s Aeronautical 

Science’s Project is to “develop high-temperature materials for turbine engines that enable a 6 percent 

reduction in fuel burn for commercial aircraft, compared to current SOA materials” [13]. Key among its 

goals is to develop CMC for use up to 1755 K (2700 °F). The presence of unreacted silicon in current 

generation CMCs makes this goal unattainable without the development of ceramic composites with Sifree advanced matrices with high crack resistance. 

1.2 Objectives 

The overall objective of the proposed research under Phases I and II was to develop a new class of 

high temperature, lightweight, self-healing, SiC fiber-reinforced ceramic composites termed Engineered 

Matrix Composites (EMCs). The proposed research was focused on developing a new matrix material 

designed for applications up to 1755 K (2700 °F) with an ability to blunt and self-heal matrix cracks, and 

to match the thermal expansion of the reinforcing SiC fibers. A matrix engineered with these capabilities 

is expected to increase the load bearing capabilities of SiC/SiC CMCs at high temperatures. The concept 

is expected to advance ceramic composite technology for applications up to 1755 K (2700 °F). 

1.3 Potential Impact on NASA and National Aeronautics Challenges 

Carbon dioxide emissions due to global air travel are currently 2 percent of total global CO2

emissions [14]. Clearly, a reduction in the global emission of greenhouse gases and an increased engine 
NASA/TM—2014-218352 3 

efficiency by increasing combustion temperature capabilities is of immediate importance for NASA, the 

United States and the global community. It is expected that a successful development of the proposed 

concept to TRL 9 would significantly increase engine efficiency, and considerably decrease fuel 

consumption and CO2 greenhouse gas emissions due to air travel. 

2.0 Theoretical Considerations in Designing Engineering Matrices 

The design of engineered matrices (EMs) involves several steps: 

(a) Introduction of chemically stable and ductile second phase particles to increase the local 

plasticity of the matrix. 

(b) Adding constituents to self-heal cracks with low viscosity oxides or silicates. 

(c) Ensuring that the melting point of engineered matrix is higher than 1755 K (2700 °F). 

(d) Matching the thermal expansion of the engineered matrix with that of the SiC fibers to ensure 

compressive residual stresses in the matrix to delay or prevent crack initiation and propagation 

in the matrix. 

(e) Melt infiltrating with a metallic silicide instead of Si to either eliminate or minimize the presence 

of unreacted Si. 

2.1 Second Phase Toughening 

It is well-known that many ceramics can be toughened by ductile second phase particles, where the 

fracture toughness is dependent on the nonlinear deformation of the particles and their volume fraction, 

Vf [15]. The critical energy release rate for crack propagation, GR, which is the work required per unit 

crack length to propagate the crack from a to a+a, is given by 



     ij

0 GR 2hVf ijd ij 2 s (1) 

where 2h is the width of the process zone around the crack, ij and ij are the local stresses and strains at 

the crack tip, and 2s is surface energy required for creating two new surfaces as the crack tip propagates 

to increase the length of the crack from a to a+a. 

Thus, if the brittle matrix in the traditional SiC/SiC CMC is suitably reformulated by adding ductile 

second phase particles to enhance its local plasticity at operating temperatures, then an approaching 

crack tip would blunt when it meets a ductile particle (Figure 1). The crack tip will be blunted if D >> r, 

where D is the dimension of the ductile phase and r is the crack tip radius. Theoretical calculations reveal 

that the force required to propagate a blunted crack is 15 to 20 percent that to propagate a sharp crack 

[16]. A blunted crack is likely to stabilize and prevent the matrix from prematurely shedding the load to 

fibers so that the matrix cracking strength and the proportional limit of the composite should increase. 

2.2 Engineering the Matrix for Self-Healing Cracks 

A second consideration in the design of engineered matrices is to add suitable additives to the matrix 

to allow the formation of low viscosity oxides to flow into cracks and heal them (Figure 2). There is a 

considerable amount of work in the literature on additives that self-heal cracks in monolithic ceramics 

and CMCs [4,17,18,19,20,21,22]. In particular, carbon-carbon composites have been successfully 

developed with self-healing characteristics for commercial applications [4]. These studies have 

demonstrated that self-healing of cracks at high temperatures results in a corresponding improvement in 

the mechanical properties and structural life of the ceramics and ceramic composites. 
NASA/TM—2014-218352 4 

Figure 1.—Schematic showing crack tip blunting in 

an engineered matrix composite due to the 

presence of a ductile second phase. 

Figure 2.—Schematic showing a self-healed 

crack in an engineered matrix composite, 

which minimizes oxygen ingress. 

2.3 Matching the Thermal Expansion of the Fibers and Matrix 

Although refractory metals, such as Nb and Ta, have been used to study the ductile phase toughening 

of ceramics [23,24,25], these metals are unsuitable for use in SiC/SiC CMCs due to the possibility that 

they will react with SiC to form metallic carbides and silicides. In the present investigation, several 

metallic silicides were considered as potential additives to the SiC matrix to enhance the crack blunting 

capabilities of the matrix due to the fact that these silicides exhibit ductile-to-brittle transition 

temperatures (DBTT) [26]. Figure 3 compares the thermal expansions of Si [27], several silicides 

[28,29],SiC [28] and Si3N4 [28]. It is evident that the thermal expansions of metallic silicides are much 

larger than those of SiC and deviate from it with increasing temperature. An inspection of Figure 3 

demonstrates that the thermal expansion of Si3N4 is lower than SiC thereby suggesting the possibility 

that a mixture of silicides and Si3N4 could be result in a matrix with a thermal expansion closer to SiC. 

Hebsur [30] decreased the thermal expansion of MoSi2 to match that of SCS-6 large diameter SiC fibers 

by adding 30 to 50 vol.% Si3N4 to the matrix, and demonstrated that the resulting composite could be 

thermally cycled between 473 and 1473 K for 1000 cycles without cracking. 

Similarly, an engineered matrix can be designed to theoretically match its thermal expansion to that 

of the fibers using a simple rule of mixtures (ROM) approach or a more sophisticated model (Figure 4). 

Referring to Figure 5, an engineered matrix is designed by mixing different volume fractions of three or 

more constituents so that the following condition is fulfilled 

 (	L/L0)fiber = (	L/L0)EM = V1(	L/L0)1 + V2(	L/L0)2 + V3(	L/L0)3 + ... (2) 

where (	L/L0)fiber is the thermal strain in the fiber, (	L/L0)EM is the thermal strain in the engineered 

matrix, (	L/L0)i is thermal expansion of the ith constituent (i = 1,2,3...n), Vi is the volume fraction of the 

i

th constituent, and V 1 n

i 1 
 i   . For the specific case of an engineered matrix consisting of a silicide, SiC 

and Si3N4, Equation (2) can be expressed as 

 (	L/L0)fiber = (	L/L0)EM = Vsilicide(	L/L0)silicide + VSiC(	L/L0) SiC + VSi3N4 (	L/L0) Si3N4 (3) 
NASA/TM—2014-218352 5 

Figure 3.—Comparison of the temperature dependence of the average 

thermal expansion behavior of hot-pressed CrSi2 [29], Cr-30%Mo-30%Si 

[29], MoSi2 [28], Si [27], SiC [28], Si3N4 [28], TiSi2 [29] and WSi2 [29] 

between 293 and 2000 K. 

Figure 4.—Schematic showing a matrix consisting of a silicide, SiC and a 

nitride engineered to match its thermal expansion with that of the SiC 

fibers. 
NASA/TM—2014-218352 6 

Figure 5.—Schematic showing the general concept used for designing engineered matrices 

with three constituents with thermal expansions matching that of SiC. 

3.0 Experimental Procedures 

3.1 Engineered Matrix Constituents 

Commercially produced silicide powders (–325 mesh) of CrSi2, TiSi2, WSi2 and a Cr-30(at.%)Mo30%Si (CrMoSi) alloy were used in this investigation. The CrMoSi alloy was procured from ATI 

Powder Metals, Pittsburgh, Pennsylvania, as gas atomized powder. Additionally, attrition milled Mo50(vol.%) Si3N4 from an earlier investigation [30] was used as the starting material. High purity Si3N4

(grade SN-E10) powders were procured from Ube America, Inc., New York, New York. High purity -

SiC (HSC059) submicron powders were procured from the Superior Graphite Co., Chicago, Illinois. 

3.2 Matrix Fabrication 

Figure 6 shows the various processing steps used in the fabrication of engineered matrices. The major 

constituents, SiC, Si3N4 and the metallic silicide were weighed in the required amount as determined using 

Equation (3). Unfortunately, extensive thermal expansion data as a function of temperature were not 

always available for some of the silicides (e.g., CrSi2) used in this research so that the calculated volume 

fractions, and the corresponding weight fractions, for these silicides were estimated at the start of the 

investigation. The Mo-50(wt.%) Si3N4 powder used in an earlier investigation [30] was reformulated by 

adding appropriate amounts of SiC. The powders were wet ball ground in ethanol using SiC grinding 

media for 48 h. After drying the powder cake, it was hand ground to powder before being hot-pressed 

under pure argon into the desired shape. The hot pressing temperatures varied between 1183 and 2073 K 

while the hot pressing stresses varied between 69 and 105 MPa for times varying between 0.25 and 4 h. 

Specimens were machined from the hot-pressed material for microstructural evaluation, thermal expansion 

measurements, oxidation studies and four-point bend testing. 

3.3 Composite Fabrication 

Attempts to fabricate EMCs using conventional slurry and melt infiltration techniques have not been 

completely successful. Two-dimensional fabrics woven out of Sylramic or Tyranno SA tows were used 

in this investigation. The Sylramic fabrics had been originally procured from Dow Corning under 

NASA’s Enabling Propulsion Materials (EPM) program while the Tyranno SA fabrics were procured 

from Ube Japan under the current program. The fibers were coated with a 0.5 m thick layer of BN and 

further coated with a thicker coating of SiC by CVI at Rolls-Royce Hyper-Therm HTC, Inc., Huntington, 
NASA/TM—2014-218352 7 

California. Figure 7 shows the various steps involved in the fabrication of the EMCs for infiltrating the 

fiber-reinforced preforms. The powders were attrition milled in ethanol using SiC or Si3N4 balls for 48 h. 

After drying the powders, a slurry was prepared using either an Epofix cold setting resin and hardener or 

an aqueous solution with a pH of 9 to 10 with surfactants and water soluble epoxy binders. In the case of 

the epoxy resin-slurry mixture, the preforms were put in a high pressure gas chamber and evacuated with 

a mechanical pump before pressuring the chamber with argon gas to above 900 psi (Figure 8). In this 

case, the epoxy infiltrated preforms were allowed to harden overnight. 

Figure 6.—Schematic showing the various processing steps used in the 

fabrication of the engineered matrices. 

Figure 7.—Schematic showing the various processing steps used in the 

fabrication of the engineered matrix composites. 
NASA/TM—2014-218352 8 

Figure 8.—Photographs of the high pressure epoxy infiltration equipment and a slurry-infiltrated 

SiC/SiC preform. 

A second method was also attempted to infiltrate the preforms using an aqueous medium. In the 

absence of a dispersant, the isoelectric points (IEP), which is defined as the pH of aqueous dispersion at 

which the surface charge is zero, is 3.3 to 4.9 for SiC [31,32,33,34] and 6.7 for Si3N4 [35], respectively. 

Lower values of pH result in a positive surface charge on the particles while higher pH values result in a 

negative surface charge. Thus, particles repel each other and remain suspended in the aqueous slurry 

when the pH significantly deviates from the IEP for the particles. Since IEP depends on both pH and 

particle size, it is important to ensure that the solution is highly acidic (i.e., low pH) or highly alkaline 

(high pH). Based on the results of earlier studies [31,32,33,34,35], it was preferable to prepare highly 

alkaline slurries with pH 9 to 10. In order to avoid contamination by alkaline ions, ammonia was used to 

increase the pH of the slurry. 

A contract was issued to NanoSperse, LLC, Kettering, Ohio, to develop aqueous slurry formulations. 

Attrition milled 20%CrSi2-55%SiC-20%Si3N4-5%CrB2 and 22.5%CrMoSiGe-49%SiC-22.5%Si3N4-

5%CrB2 powders were first soaked overnight in a 29 (wt.%) ammonia solution. The particles were 

washed in water using a Buchner funnel, where the particles were filtered using a filter paper. The 

particles were dried in a vacuum oven maintained at 283 K under a pressure of 338.6 Pa for 2h. The 

slurry was prepared using 10 mL of a 10 percent Triton X100 dispersant. A 3 percent water suspended 

epoxy binder was added to the final formulation of the slurry, which contained 30 (wt.%) of the 

engineered matrix powder. Initial suspension tests revealed that the particles remained in suspension for 

at least 24 h. However, preforms infiltrated with these slurries under a mechanical pump vacuum did not 

show sufficient solids loading after one infiltration attempt. Attempts to increase the volume fractions of 

the particles by both particulate infiltration techniques are continuing. 

The specimens were then pyrolized to convert the organic material to carbon before melt infiltrating 

with either Si or Si-1(at.%)Ge alloy. Melt infiltration was conducted in a chamber evacuated to a 

vacuum level of about 1.3
10-4 Pa (~1
10-6 torr). Alternate to the traditional melt infiltration with Si, 

attempts were made to infiltrate the preforms with CrSi2 or Cr-25(wt.%)Si. Here too, more work is 

needed to completely infiltrate the preforms. 
NASA/TM—2014-218352 9 

3.4 Materials Characterization 

3.4.1 Microstructural and Non-Destructive Characterization 

Specimens were cut, mounted and polished for metallographic examination by optical and scanning 

electron microscopy (SEM). Computerized tomography (CT) scans were used to determine the extent of 

particle and melt infiltration of the preforms, and these images were compared with the microstructural 

results. 

3.4.2 Thermal Expansion Measurements 

Thermal expansion measurements were conducted on hot-pressed cylindrical specimens 25.4-mm 

long and 9.5-mm in diameter using a NETZSCH Dilatometer Model DIL 402C equipped with a high 

purity alumina as a calibration standard, where the two faces were machined to ensure that they were flat 

and parallel. Measurements were made over three heat-cool cycles to (a) minimize the effects of 

compositional, microstructural and processing variables on the data, (b) determine the extent of scatter in 

the data, and (c) to evaluate statistical averages of the coefficients for the regression curve. The specimen 

was placed in a sample holder and aligned with a single push-rod under an applied constant load of 0.2 

N. The specimens were heated from 293 to 1523 K at 10 K/min. in the first cycle and cooled to 373 K at 

10 K/min. in the first cool-down cycle. Subsequent cycles consisted of heating and cooling between 373 

and 1523 K. All measurements were conducted in a He atmosphere flowing at 60 cc/min. The length 

changes were recorded by a computerized data acquisition system. The experimental strain, 	L/L0, 

where 	L is the differential change in length, L-L0, and L0 is the original length of the specimen at room 

temperature, were measured. 

3.4.3 Isothermal Oxidation 

The engineered matrices that were formulated with different metallic silicide constituents were 

isothermally oxidized at 1600 K for 100 h. The nominal dimensions of the oxidation specimens were 

13.5 mm in diameter and 3 mm thick. Each specimen had a 2.5 mm diameter hole machined by electrodischarge machining (EDM) about 3.2 mm from the edge to allow it to hang from a quartz hook in a 

thermogravimetric analyzer (TGA) equipped with a weighing balance to periodically monitor the weight. 

The data were logged by a computer data acquisition system. Oxidation was conducted in dry air flowing 

at 100 sccm. 

3.4.4 Mechanical Testing 

Four-point bend tests were conducted on engineered matrices where the nominal specimen 

dimensions were 50 x 4 x 3-mm. These specimens were machined flat and parallel from hot-pressed 

plates 50 x 50 x 3-mm in dimensions using a diamond saw. Bend tests were conducted in air between 

room temperature and 1698 K in a servo-hydraulic uniaxial testing machine at a speed of 8.3
10-3 mm/s 

corresponding to an approximate engineering strain rate of 2
10-3 s-1. 

3.4.5 Self-Healing Studies 

Several CrSi2/SiC and CrMoSi/SiC matrices were formulated without Si3N4 to study the effect of 

CrB2, Ge, Y and ZrSiO4 additives on self-healing cracks and holes. These compositions were prepared as 

described in Section 3.2 and hot-pressed into disks 12.7 mm diameter and 3.2 mm thickness. Initial 

attempts to indent these specimens with a Vickers hardness indent to generate cracks were unsuccessful. 

Instead, a 1 mm diameter hole was drilled in each specimen before annealing them in air in a rapid 

heating box furnace. The specimens were oxidized either at 1600 or 1700 K for a cumulative oxidation 

time of 100 h, and the oxidized specimens were microstructurally examined to study the extent of scratch 

healing and hole closure. Other tests that are underway involve studying the effects of the additives on 

healing cracks developed from hardness indents (Figure 9). 
NASA/TM—2014-218352 10 

Figure 9.—Photograph of a Vickers hardness indent showing cracks 

(red arrows) emanating from it. These specimens will be oxidized at 

high temperatures to study the effect of additives on the self-healing 

of these cracks. 

Figure 10.—Photograph of a bend specimen showing a SEPB notch. The notched specimens will be 

tested in air and inert gas to demonstrate that the air-tested specimens are stronger than those 

tested in inert gas due to self-healing of cracks as shown in the schematic figure on the left. 

Based on these observations, several CrSi2/SiC/Si3N4 and CrMoSi/SiC/Si3N4 engineered matrices 

were formulated and hot-pressed into 50 x 50 x 3 mm plates. Several bend specimens with dimensions 

50 x 4 x 3-mm were machined for fracture toughness testing in argon and air using the single edge Vnotched beam (SEVNB) method [36]. These tests are still underway and only preliminary results are 

reported in this report. Sharp notches were cut into the specimens to help initiate cracks during loading 

(Figure 10). These specimens will be tested in air and argon to demonstrate the self-healing of the notch 

as shown in the schematic stress-strain curve. Elastic moduli, EMOD, measurements were conducted on 

some specimens before and after heat treatment in air and argon at 1700 K using the Impulse Excitation 
NASA/TM—2014-218352 11 

Figure 11.—Relative weight change versus oxidation time for slurry and 

melt-infiltrated SiC/SiC preforms. 

technique [37,38,39]. The desired loading profile shown in Figure 10 could not be precisely followed 

since the elastic moduli measurements indicated that the specimens were porous. Thus, the SEVNB 

specimens were pre-annealed at 1698 K for 1800 s prior to loading. 

Four-point bend tests were conducted on the slurry infiltrated and melt infiltrated composite 

specimens at room temperature and 1645 K. Tests were also conducted on uninfiltrated preforms. In 

order to study the effect of self-healing without additives, some infiltrated composite specimens were 

pre-cracked at room temperature on a three-point bend fixture, heat-treated at 1600 K for 50 h and retested at 1645 K. The self-healing times were determined by oxidizing small specimens at 1600 K for 

different lengths of time and measuring the relative weight change as a function of oxidation time 

(Figure 11). As shown in Figure 11, the curves flatten out at about 50 h. 

4.0 Results and Discussion 

4.1 Thermal Expansion and Proof-of-Concept 

4.1.1 Microstructures of Thermal Expansion Specimens 

Microstructural observations of the hot-pressed matrix specimens revealed that most of them were 

well-consolidated although the extent of homogeneity varied from one composition to another. Figure 

12(a) shows a photograph of a hot-pressed 10(wt.%)CrMoSi-60%SiC-30%Si3N4 (CrMoSi-EM) 

specimen while Figure 12(b) and (c) show the corresponding optical, and back-scattered electron 

microscopy (BSE) images and energy dispersion spectra (EDS), respectively, where it was observed that 

the white particles were CrMoSi. The microstructure was fairly homogenous. Figure 13(a) to (c) show 

the optical micrographs of 10(wt.%)CrSi2-70%SiC-20%Si3N4 (CrSi2-EM), 10(wt.%)TiSi2-70%SiC20%Si3N4 (TiSi2-EM) and 10(wt.%)WSi2-70%SiC-20%Si3N4 (WSi2-EM), respectively, where the degree 

of homogeneity varied between the three engineered matrices. These microstructures did not contain any 

significant amount of cavities. Figure 14 shows the BSE image for a TiSi2-EM specimen, where the EDS 

show that the gray particles are Ti-poor (Region C) while the white particles are TiSi2 (Region D). The 

WSi2-EM specimen unexpectedly had several white W-rich (Region A) and gray Si-rich (Region B) 

particles (Figure 15 and Figure 16). The source of the tungsten particles was attributed to the WSi2

powder used in fabricating these matrices. 
NASA/TM—2014-218352 12 

Figure 12.—(a) Photograph of hot-pressed 10(wt.%)CrMoSi-60%SiC-30%Si3N4 engineered matrix; (b) optical 

micrograph; and (c) back-scattered electron micrograph and energy dispersion spectra. 

 

Region J 

Region H 

Region I 
NASA/TM—2014-218352 13 

Figure 13.—Optical micrographs of hot-pressed engineered matrices: (a) 10(wt.%)CrSi2-70%SiC-20%Si3N4; 

(b) 10(wt.%)TiSi2-70%SiC-20%Si3N4; and (c) 10(wt.%)WSi2-70%SiC-20%Si3N4. 

 
NASA/TM—2014-218352 14 

Figure 14.—Back-scattered image of hot-pressed 10(wt.%)TiSi2-70%SiC-20%Si3N4 engineered matrix and 

energy dispersion spectra of regions C and D showing gray Si-rich and white TiSi2 particles, respectively. 

 
NASA/TM—2014-218352 15 

Figure 15.—Back-scattered image of hot-pressed 10(wt.%)WiSi2-70%SiC-20%Si3N4 engineered matrix and energy 

dispersion spectra of regions A and B showing W and WSi2 particles, respectively. 

 
NASA/TM—2014-218352 16 

 SEM C 

 Si W 

Figure 16.—Scanning electron image of hot-pressed 10(wt.%)WiSi2-70%SiC-20%Si3N4 engineered 

matrix and C, Si and W x-ray maps showing W and WSi2 particles, respectively. 

4.1.2 Thermal Expansion of Silicides 

Details of the thermal expansion data generated on the disilicides and the CrMoSi alloy in this 

investigation are reported in detail elsewhere [29]. For some silicides, the thermal expansion during the 

first heat-up cycle was different than those in the subsequent cool-down and heat-up cycles. This 

difference in behavior was attributed to the effect of residual stresses generated during cool down from 

the hot-pressing temperatures, which are annealed out after the first heat-up cycle [40]. Figure 17(a) 

shows an example of the thermal expansion data for the three heat-cool cycles for CrSi2. Neglecting the 

data from the first heat-up cycle, the data from the first cool-down to the third cool-down cycles were 

fitted with equation (4): 

 (	L/L0)thermal = A(T-293)3

 + B(T-293)2 + C(T-293) + D (4) 

where, (	L/L0)thermal is the magnitude of the 	L/L0 without any residual processing strains, T is the 

absolute temperature, A, B, C and D are regression constants. Table I compiles the values of the 

regression constants and the corresponding coefficients of determination, Rd

2

, for the CrSi2 specimens 

from three hot-press runs.1

 Figure 17(b) compares the regression plots for the three specimens, where it 

is seen that the curves are reasonably close thereby suggesting that batch-to-batch variability is relatively 

small. Unfortunately, a sparsity of published data on polycrystalline CrSi2 did not permit a meaningful 

comparison with the present results. 

 

1

 The specimen from Run 582 had more porosity than those from Runs 619 and 620. 
NASA/TM—2014-218352 17 

TABLE I.—VALUES OF THE REGRESSION CONSTANTS 

FOR CrSi2 SPECIMENS FROM THREE HOT-PRESS RUNS 

Hot press run 

no. 

A, 

K–3

B, 

K–2

C,

K 

D, 

% 

Rd

2

582 2.0
10–10 1.2
10–7 1.0
10–3 –1.8
10–1 0.993 

619 4.2
10–10 –1.4
10–7 1.1
10–3 –1.4
10–1 0.999 

620 4.2
10–10 –1.8
10–7 1.1
10–3 –9.3
10–2 0.999 

Figure 17.—Temperature dependence of the thermal expansion behavior of (a) a 

CrSi2 specimen during three heat up-cool down cycles between 293 and 1523 K; 

and (b) comparison of the regression equations for three CrSi2 specimens 

between 303 and 1523 K. 
NASA/TM—2014-218352 18 

Figure 18(a) to (c) show the thermal expansion behaviors of the CrMoSi, TiSi2 and WSi2, 

respectively, for three heat-cool cycles. There was no significant difference in the thermal expansion 

behavior for the CrMoSi and TiSi2 between the first heat-up and the other thermal cycles. In contrast, the 

WSi2 exhibits somewhat higher values of 	L/L0 in the first heat-up cycle than in the subsequent cycles 

although the difference was not as large as for the CrSi2 specimens. Table II gives the values of the 

regression constants and Rd

2

 for these three silicides. Since the CrMoSi alloy was prepared by replacing 

Cr with Mo in Cr3Si [41], it is reasonable to compare its thermal expansion with that for Cr3Si [28]. The 

addition of Mo appears to decrease the thermal expansion of Cr3Si slightly at the higher temperatures. 

The present values of 	L/L0 are in reasonable agreement with the literature values for TiSi2 and WSi2

although the deviation in the two sets of data slightly increases with increasing temperature [28]. 

Figure 18.—Continued. 
NASA/TM—2014-218352 19 

Figure 18.—Concluded. Temperature dependence of the thermal expansion 

behaviors of (a) a Cr-30(at.%)Mo-30%Si specimen, (b) TiSi2 and (c) WSi2 during 

three heat up-cool down cycles between 293 and 1523 K. The literature data for 

Cr3Si, TiSi2 and WSi2 are also shown in the figure for comparison [28]. 

TABLE II.—VALUES OF THE REGRESSION CONSTANTS AND 

COEFFICIENTS OF DETERMINATION FOR Cr-30%Mo-30%Si, TiSi2 AND WSi2

Alloy A, 

K–3

B, 

K–2

C, 

K 

D, 

% 

Rd

2

CrMoSi 1.3
10–10 –1.7
10–8 9.2
10–4 –2.5
10–2 0.9993 

TiSi2 1.8
10–10 –1.7
10–7 1.2
10–3 –2.0
10–2 0.999 

WSi2 1.2
10–10 –1.2
10–7 8.3
10–4 –8.0
10–2 0.999 

4.1.3 Effect of Si3N4 Addition on Thermal Expansion 

As demonstrated in Figure 3, the magnitudes of 	L/L0 for silicides are larger than that for SiC and 

Si3N4 with the extent of deviation increasing with increasing temperature. In order to minimize this 

difference in 	L/L0, it was proposed that the addition of 50(vol.%) of Si3N4 to MoSi2 can decrease its 

thermal expansion to values close to those of SiC [30]. However, a close examination of the heat-cool 

cycles for this material suggested that its thermal expansion behavior was somewhat more complex due 

to the development of residual stresses during thermal cycling [42] (Figure 19). Similarly, the thermal 

expansion of a TiSi2-80(vol.%) Si3N4 material is close to that of SiC [28] only during the first heat-up 

cycle (Figure 20). However, its thermal expansion deviated from that of SiC during subsequent cooldown and heat-up cycles due to the development of thermal residual stresses in the matrix. 
NASA/TM—2014-218352 20 

Figure 19.—Temperature dependence of the thermal expansion behavior of a 

MoSi2-50(vol.%)Si3N4 specimen during three heat up-cool down cycles between 

293 and 1688 K [42]. The data for SiC is included for comparison [28]. 

Figure 20.—Temperature dependence of the thermal expansion behavior of 

TiSi2-80(vol.%)Si3N4 during three heat up-cool down cycles between 293 

and 1523 K. The literature data for SiC is also shown in the figure for 

comparison [28]. 

4.1.4 Thermal Expansion Data for Engineered Matrices 

As noted above, the design of an engineered matrix cannot simply involve mixing a silicide and 

Si3N4 in appropriate amounts in order to match the thermal expansion with that of the SiC fibers. The 

resulting magnitudes of the thermal stresses can eventually become significant after several cycles and 

result in the catastrophic failure of the matrix. However, formulating the matrix in accordance with 

Equation (3) as described in Section 2.3 generally results in a thermally stable matrix capable of 

withstanding several thermal cycles (Figure 21(a) to (e)). Figure 21(a) demonstrates the almost close 

correspondence of the 	L/L0 values for CrSi2-EM and SiC [28] between 293 and 1523 K during both the 
NASA/TM—2014-218352 21 

Figure 21.—Continued. 
NASA/TM—2014-218352 22 

Figure 21.—Concluded. Temperature dependence of the thermal expansion behaviors of 

(a) 10%CrSi2-70%SiC-20%Si3N4; (b) 10%CrMoSi-60%SiC-30% Si3N4; (c) 10%MoSi2-

70%SiC-20%Si3N4; (d) 10%TiSi2-70%SiC-20%Si3N4; and (e) 10%WSi2-70%SiC20%Si3N4; engineered matrices during three heat up-cool down cycles between 293 and 

1523 K. The literature data for SiC is also shown in the figure for comparison [28]. 

heating and cooling cycles. The fact that the 	L/L0 values are lower for the CrSi2-EM than SiC suggests 

that compressive residual stresses are likely to be present in the matrix at room temperature, which 

should enhance room temperature fracture strength of the matrix. Similar thermal expansion behavior is 

observed for the CrMoSi-EM specimen (Figure 21(b)). Figure 21(c) shows the thermal expansion values 

for 10%MoSi2-70%SiC-20%Si3N4 (MoSi2-EM), where the composition of the original material used to 

generate the data shown in Figure 19 was reformulated. On the first heat-up cycle, the material exhibited 

a large thermal expansion. However, after the material was reheated again and thermally cycled between 

293 and 1523 K, the magnitudes of 	L/L0 match those of SiC almost exactly. As shown in Figure 22, the 

specimen had developed several circumferential and longitudinal cracks on the surface during the 

thermal expansion measurements. Owing to the potential thermal instability of this material during 

thermal cycling, the MoSi2-EM was no longer considered for further development. The thermal 

expansion responses for the TiSi2-EM (Figure 21(d)) and WSi2-EM (Figure 21(e)) were similar to the 

observations on the CrSi2-EM matrix (Figure 21(a)). Once again, the thermal expansions for these 
NASA/TM—2014-218352 23 

engineered matrices match those of SiC very closely and do not exhibit any significant change during 

thermal cycling. 

It is clear from Figure 21 that except for MoSi2-EM, which exhibited thermal instability during the 

first heat-up cycle, all other engineered matrices show little change in the magnitudes of 	L/L0 during 

the heat-cool cycles. Noting that these matrices were formulated using the concepts described in Section 

2.3, it is remarkable that the magnitudes of 	L/L0 are very close to those for SiC [28] as predicted by 

Equation (3) (Figure 23). These observations provide a complete proof of the proposed concept of 

ensuring thermal compatibility between the engineered matrix and the SiC fibers postulated in item (d) 

of Section 2.0. 

Figure 22.—Macrograph of the 10%MoSi2-70%SiC-20%Si3N4 specimen 

after thermally cycling between 293 and 1523 K showing several 

longitudinal and circumferential cracks. 

Figure 23.—Comparison of the temperature dependence of the thermal strains for 

disilicides, CrMoSi, engineered matrices with those for SiC [28] and Si3N4 [28]. 
NASA/TM—2014-218352 24 

Figure 24.—Comparison of the isothermal oxidation behavior of several 

engineered matrices with those for SiC [43] and Si3N4 [43] showing the change 

in the specific weight change with oxidation time. 

4.2 Isothermal Oxidation Behavior of Engineered Matrices 

Figure 24 compares the isothermal oxidation behaviors of CrSi2-EM, CrMoSi-EM, MoSi2-EM, 

TiSi2-EM and WSi2-EM, at 1600 K, where the weight change, 	W, per unit area, A, is plotted against 

oxidation time, t. The isothermal oxidation behavior of CVD SiC [43] and Si3N4 [43] at 1573 K are 

shown for comparison. In contrast to SiC and Si3N4, the engineered matrices oxidize at a much higher 

rate, which suggests that these matrices would be able to easily getter any oxygen that infiltrates the 

composite matrix therefore adding an additional layer of protection for the BN-coated SiC fibers. 

The engineered matrices oxidize rapidly on heating in air as the silicides begin to form a metallic oxide 

layer. However, as the more protective silica layer begins to form, the oxidation rate slows down in 

accordance with the parabolic oxidation kinetic law. An examination of Figure 24 reveals that the initial 

oxidation rates for TiSi2-EM and WSi2-EM are higher than those for CrSi2-EM, CrMoSi-EM and MoSi2-

EM but the magnitudes of specific weight gain, 	W/A, for these engineered matrices are similar at 

t = 100 h. Curiously, the oxidation rates for the two CrSi2-EM specimens were quite different, where one of 

the specimens exhibited relatively little specific weight gain after the initial oxidation while the other 

specimen continued to gain weight even after 100 h. This difference in behavior is attributed to an 

inhomogeneous distribution of the CrSi2 particles. Since the TiSi2-EM spalled after 100 h, this matrix 

composition was eliminated from further consideration in the investigation. 

4.3 Bend Properties of Engineered Matrices 

As postulated in item (a) Section 2.0 and discussed in Section 2.1, increasing the local plasticity of 

the matrix should result in an increase in the strength of the matrix as the crack tip is blunted and the 

fracture toughness of the material increases. Figure 25 shows the bend stress-strain curves for the CrSi2-

EM between room temperature and 1588 K. On close examination, it is evident that the bend strength 

increases with increasing temperature above 898 K corresponding to an increase in plasticity of the 

material. Above 1273 K, general plasticity increases and the strength decreases as the material deforms 

readily. However, the specimen tested at 1588 K exhibits the highest strength presumably because of the 

dominance of local plasticity over general plasticity which serves to blunt the crack. The advent of local 

plasticity rather than general plasticity is attributed to an inhomogeneous distribution of the CrSi2
NASA/TM—2014-218352 25 

particles. Clearly, designing the matrix in order to enhance local plasticity while reducing the advent of 

general plasticity would be advantageous in a composite. 

Figure 26 shows the four-point bend stress-strain curves for the CrMoSi-EM matrix. Once again, the 

bend strength increases with increasing temperature up to 1573 K as the plasticity of the material 

increases. Above 1573 K, the strength drops at the higher temperatures although the data does not vary 

in a consistent manner with increasing temperature. The reason for a lack of a consistent trend is unclear. 

Nevertheless, the results shown in Figure 26 are consistent with the discussion in Section 2.0 that 

increasing the local plasticity of the matrix results in a corresponding increase in its strength. 

Figure 25.—Four-point bend stress-strain curves for 10%CrSi2-70%SiC20%Si3N4 tested in air between room temperature (RT) and 1588 K showing 

that the bend strength increases with increasing temperature due to an 

increase in plasticity. 

Figure 26.—Four-point bend stress-strain curves for 10%CrMoSi-60%SiC30%Si3N4 tested in air between room temperature (RT) and 1698 K showing 

that the bend strength increases with increasing temperature due to an 

increase in plasticity. 
NASA/TM—2014-218352 26 

High temperature four point bend stress-strain data could not be determined for the WSi2-EM since 

the specimens catastrophically oxidized (i.e., “pest”) and crumbled to powder during heat-up to test 

temperature. This catastrophic oxidation is due to the fact that the microstructure unexpectedly contained 

W-rich particles (Figure 15 and Figure 16). Therefore, this engineered matrix was dropped from further 

consideration in this investigation. 

Based on the bend, isothermal oxidation and thermal expansion results discussed earlier, the CrSi2-

based and CrMoSi-based compositions were down selected for further development for 1588 (2400 °F) 

and 1755 K (2700 °F) composite matrix applications, respectively. 

4.4 Composite Fabrication Trials 

As mentioned in Section 3.3, attempts to fabricate EMCs with a high and well distributed engineered 

matrix and minimizing the void volume fraction in the preforms to a minimum has not been completely 

successful. More research and development is continuing under the Aerosciences Project and the current 

description relates to the outcome of research conducted under Phases I and II of the ARMD Seedling 

Program. 

Infiltrating the preforms with particle slurry and silicide melt proved to be particularly challenging 

for several reasons. Figure 27(a) shows the top view of a Sylramic fiber-reinforced preform while 

Figure 27(b) shows the CT scan image of its cross-section. The void area fraction varied between 21 and 

23 percent. As shown in the schematic Figure 27(c), the voids are seldom directly connected to the free 

surface so that complete infiltration of the preform by the particulate slurry and molten silicon or silicide 

to fill these cavities is difficult to achieve due to the tortuous path that the slurry and molten metal have 

to follow. Additionally, particles in the slurry are likely to be filtered by the woven SiC tows so that it is 

desirable that the particle size be small enough to be able to enter the voids and the particles be well 

separated and bound to the carrier medium. With both particulate slurry and melt infiltration, it is 

important to adjust the viscosity of the fluid and to ensure that it adequately wets the SiC fiber surface so 

that it penetrates the interior of the preform. 

4.4.1 Slurry Infiltration 

The particle size distribution was determined for attrition milled 20%CrSi2-55%SiC-20%Si3N4-

5%CrB2 and 22.5%CrMoSiGe-49%SiC-22.5%Si3N4-5%CrB2 powders2

 (Figure 28). The average 

diameters were 0.389 and 0.530 m, respectively, with the maximum powder diameters being 0.611 and 

0.572 m, respectively. For comparison, the void sizes in the preforms were typically several hundred 

microns. During the early stages of this investigation, the slurries were prepared by mixing ethanol and 

the powder, and infiltration was conducted under a mechanical pump vacuum. Although the preforms 

showed an increase in weight after each infiltration3

 (Figure 29), microstructural observations of the 

cross-sections of the preforms revealed that the extent of particulate infiltration was relatively small. 

Thus, weight change measurements of the infiltrated preforms proved to be unreliable in predicting the 

amount of solids loading. 

Infiltration trials using powder-epoxy slurries at high pressure were more successful (Figure 30). 

Although the epoxy successfully completely filled the voids, the amount of particulates that could be 

infiltrated into the preforms was much less than the desired amount of filling about 50 percent of the void 

space between the woven tows. When the particulate-to-epoxy mixture ratio was increased, it was difficult 

to completely infiltrate the preforms (Figure 31). Even in regions where the epoxy-particulate mixture had 

infiltrated, the distribution of the particles was inhomogeneous (Figure 32). These observations suggest that 

the particles were not well separated and they were not bonded to the epoxy which effectively allowed the 

fibers to act as filters and separate them from the epoxy. Therefore, a considerable amount of effort was 

expended in trying to find a suitable binder but this effort has not been successful to date. 

 

2

 These measurements were conducted by NanoSperse, LLC, Dayton, Ohio. 

3

 The initial slurry infiltration trials were conducted with a 10%TiSi2-70%SiC-20%Si3N4. 
NASA/TM—2014-218352 27 

Figure 27.—(a) Optical view of the top surface of a Sylramic 

preform; (b) CT scan of the cross-section of the preform; and 

(c) schematic of the cross-section showing the distribution of 

voids. The area fraction of the voids was estimated by image 

analysis of the CT scan was 21 to 23%. 
NASA/TM—2014-218352 28 

Figure 28.—Particle size distribution for attrition milled 20%CrSi2-55%SiC20%Si3N4-5%CrB2 and 22.5%CrMoSiGe-49%SiC-22.5%Si3N4-5%CrB2

powders. 

Figure 29.—Changes in the relative weight, density and estimated 

porosity of a Sylramic preform infiltrated with a 10%TiSi2-70%SiC20%Si3N4 powder-ethanol slurry after each infiltration trial. 
NASA/TM—2014-218352 29 

Figure 30.—Optical micrographs showing (a) low and (b) high magnification views of the 

cross-section of a Sylramic preform infiltrated with an epoxy-10%TiSi2-70%SiC20%Si3N4 slurry at high pressure. 

Figure 31.—Optical micrographs showing the cross-section of a Sylramic preform infiltrated with 

an epoxy-10%TiSi2-70%SiC-20%Si3N4 slurry at high pressure. (a) Composite view of the 

preform cross-section in fluorescent light; (b) bright field and (c) fluorescent light views of the 

square region shown in (a). 
NASA/TM—2014-218352 30 

Figure 32.—Optical micrographs showing the cross-sections of a Sylramic preform 

infiltrated with an epoxy-10%CrMoSi-60%SiC-30%Si3N4 slurry at high pressure. 

In view of the fact that there is a considerable amount of research on preparing SiC [31,44,45,46,47] 

and Si3N4 [35,44] particles dispersions in water, where the particles are well-separated due to the 

addition of surfactants, more recent effort has concentrated on developing aqueous-based slurries. As 

noted in Section 3.3, preliminary infiltration trials did not result in sufficient solids loading in the 

preforms. This research is currently in progress. 

4.4.2 Melt Infiltration 

Several bend specimens, which had been infiltrated with epoxy-particulate slurry at high pressure 

and subsequently heated in flowing argon to pyrolize the epoxy to carbon, were infiltrated with either Si 

or Si-1(at.%)Ge. Figure 33 shows the CT images, the processed images and the estimated porosity in a 

preform specimen at each stage of the composite fabrication process, where the red areas in the 

processed images represent voids. Although the CT images suggested that the area fraction of the 

porosity was 0.9 percent after epoxy-particulate infiltration, these images could not distinguish between 

the epoxy and the solids. As noted earlier, the solids loading after the epoxy-particle infiltration was low 

(Figure 31 and Figure 32). The porosity content after Si melt infiltration was about 1.8 percent. Figure 

34 shows the microstructures of the cross-sections of the specimen, where it is clear that some regions of 

the preform were not infiltrated. In a second trial, where the amount of Si or Si-1(at.%)Ge was reduced, 

it was observed that large areas of the preform were not infiltrated (Figure 35(a) to (d)). 

In order to develop a silicon-free composite matrix as postulated in item (e) in Section 2.0, melt 

infiltration was conducted using CrSi2 and Cr-25(wt.%)Si. Figure 36 shows the Cr-Si binary phase 

diagram [48]. It is clear that CrSi2 and Si have melting points of 1763 and 1687 K, respectively. The Cr25(wt.%)Si has a solidus temperature of 1778 K. Thus, melt infiltrating with CrSi2 and Cr-25(wt.%)Si 

alloys should increase the ability of the matrix to withstand temperatures higher than composites melt 

infiltrated with Si. 
NASA/TM—2014-218352 31 

Figure 33.—CT scan (left) and processed (right) images of each stage of fabrication of 

an EMC. (a) As-received preform; (b) preform infiltrated with an epoxy-10%TiSi2-

70%SiC-20%Si3N4 slurry; (c) preform after pyrolizing the epoxy; and (d) preform after 

melt infiltration with silicon. The red areas in the processed images represent voids. 

The area fractions of the voids are shown after each processing stage. 

Figure 34.—Optical micrographs showing the cross-sections of a Sylramic preform 

infiltrated with an epoxy-10%CrMoSi-60%SiC-30%Si3N4 slurry and melt infiltrated 

with silicon. 
NASA/TM—2014-218352 32 

Figure 35.—Optical micrographs showing the cross-sections of Sylramic 

preforms infiltrated with an epoxy-90%SiC-10%Si3N4 slurry and melt infiltrated 

with (a) and (c) silicon and (b) and (d) Si-1(at.%)Ge. 

 

Figure 36.—Cr-Si binary phase diagram [48]. 
NASA/TM—2014-218352 33 

The Sylramic and Tyranno SA fibers contain TiB2 and Al present as SiC1.08Al0.005 [49], respectively. The 

Sylramic fibers contain 2.3 (wt.%) B, 2.1 (wt.%) Ti and 0.8(wt.%) O [50]. The Tyranno SA fibers contain 

less than 2(wt.%) Al and about 0.3(wt.%) O [50,51]. Thus, it is important to verify that preforms melt 

infiltrated with these chromium silicides will be chemically stable both during melt infiltration as well as in 

the actual applications. Thermodynamic calculations using FactSage (Thermfact/CRCT (Montreal, Canada) 

and GTT-Technologies (Aachen, Germany)) were conducted to determine whether there would be any 

reactions with either the CVI SiC or the BN coatings. As shown in Table III, the thermodynamic calculations 

predicted that the CVI SiC and BN coatings are unlikely to be attacked by either molten CrSi2 or Cr-25%Si 

under an infiltration pressure of 0.1 MPa but the SiC, Si3N4 and the silicides decompose at a lower pressure of 

1.3 x 10-10 MPa . The decomposition of SiC under low pressures of 1.3 x 10-10 MPa is expected to leave a 

layer of carbon on the fiber surfaces readily available to react with either the molten silicides. The FactSage 

calculations also predicted that the presence of small amounts (0.1%) impurities, such as B, Hf, Ti and Zr, 

would react to form CrB2, HfC, TiSi2 and ZrC, respectively, at 1808 K. While trace amounts Al were not 

expected to react, the thermodynamic calculations suggested that it would form liquid Al at these 

temperatures. 

Figure 37(a) and (b) show optical micrographs of a Tyranno SA preform melt infiltrated with CrSi2

in the first trial. The preform was not completely infiltrated due to poor wetting action as evidenced by 

the distinct meniscus seen in Figure 37(b). A careful examination of the CrSi2-SiC interfaces revealed 

that there was no reaction between the two phases (Figure 38(a) to (c)). Raman spectra from the CrSi2

matrix confirmed the absence of free C and Si, which was one of the requirements postulated in item (e) 

in Section 2.0. A small amount of free Si was detected in the CVI SiC coating on the fibers. 

Figure 37.—(a) Low and (b) high magnification optical micrographs of the cross-section of a 

Tyranno SA preform melt infiltrated with CrSi2. 
NASA/TM—2014-218352 34 

TABLE III.—RESULTS FROM THE FACTSAGE CALCULATIONS 

Reaction T 

(K) 

P 

(MPa) 

Products Amount 

(mol.) 

Activity/ 

fugacity 

(MPa) 

Remarks 

-SiC (s)+ Si3N4 (s) 1808 1.3x10-10 

C (s) 

Si (g) 

N2 (g) 

0.8 

3.9 

2.0 

1.0 

8.5x10-11 

4.3x10-11

SiC and Si3N4 likely 

to decompose 

-SiC (s)+ Si3N4 (s) +CrSi2 (s) 1808 1.3x10-10 

C (s) 

Cr (g) 

N2 (g) 

Si (g) 

0.7 

1.0 

2.0 

5.9 

1.0 

1.4x10-11 

2.3x10-11

8.5x10-11 

SiC and Si3N4 likely 

to decompose 

-SiC (s)+ Si3N4 (s) +CrSi2 (s) 1808 0.1 

-SiC (s) 

Si3N4 (s) 

CrSi2 (s) 

1.0 

1.0 

1.0 

1.0 

1.0 

1.0 

No reaction with 

SiC and Si3N4

-SiC (s)+Si3N4 (s) CrSi2

(s)+BN (s) 1808 1.3x10-10

C (s) 

B (g) 

Cr (g) 

N2 (g) 

Si (g) 

0.7 

1.0 

1.0 

2.5 

5.9 

1.0 

1.2x10-11 

1.2x10-11

3.1x10-11

7.3x10-11

SiC and Si3N4 likely 

to decompose 

-SiC (s)+Si3N4 (s) CrSi2

(s)+BN (s) 1808 0.1 

BN (s) 

-SiC (s) 

Si3N4 (s) 

CrSi2 (s) 

1.0 

1.0 

1.0 

1.0 

1.0 

1.0 

1.0 

1.0 

No reaction with 

SiC and Si3N4

-SiC (s)+ Si3N4 (s) 

CrMoSi (s)+CrSi2 (s) +BN

1808 1.3x10-10

C (s) 

MoB (s) 

Cr (g) 

N2 (g) 

Si (g) 

0.7 

0.3 

1.4 

2.5 

6.1 

1.0 

1.0 

1.7x10-11

3.0x10-11

7.3x10-11

Reaction of Mo and 

BN; SiC and Si3N4

likely to decompose 

-SiC (s)+ Si3N4 (s) 

CrMoSi (s)+CrSi2 (s) +BN 

1808 0.1 

BN (s) 

-SiC (s) 

Si3N4 (s) 

CrSi2 (s) 

Cr5Si3 (s) 

MoSi2 (s) 

1.0 

1.1 

1.0 

0.5 

0.2 

0.3 

1.0 

1.0 

1.0 

1.0 

1.0 

1.0 

No reaction with 

SiC and Si3N4

 
NASA/TM—2014-218352 35 

Figure 38.—(a) Scanning electron and (b) and (c) back scattered electron images of the cross-section of a Tyranno 

SA preform melt infiltrated with CrSi2. The corresponding energy dispersive spectra confirm that there was no 

reaction between the CrSi2 and SiC. 
NASA/TM—2014-218352 36 

 

Figure 39.—Raman spectra for C and CrSi2 from regions 1 to 6 identified 

in the optical micrograph above demonstrating the absence of free C 

and Si. 
NASA/TM—2014-218352 37 

Since the results from the first CrSi2 melt infiltration trials were encouraging, subsequent trials were 

conducted at higher temperatures and longer times in order to increase the fluidity of the silicide and 

improve the degree of infiltration of the preforms. However, an examination of the microstructures 

showed considerable evidence that the fibers had been attacked during melt infiltration (Figure 40 and 

Figure 41). Energy dispersive spectroscopy from four areas of Sylramic preforms revealed the following 

(Figure 40): The EDS pattern from Region A showed a large Si peak with much smaller B and Cr peaks; 

well-defined B and Ti peaks as well as a diffused Si peak are observed in Region B; Region C shows 

only a large Si peak; and Region D within the fiber shows only Ti and Si peaks. Surprisingly, except in 

Region A, Cr peaks are largely absent from the other regions. Similarly, the Tyranno SA preform in a 

second melt infiltration trial showed evidence of reaction with the fiber CVI SiC and BN coatings 

(Figure 41). Figure 40 and Figure 41 suggest that the CrSi2 has reacted with the CVI SiC and the BN 

coatings, which is inconsistent with the thermodynamic predictions detailed in Table III. The exact 

reason for this inconsistency is unclear and it is being investigated further. 

4.5 Bend Properties of Engineered Matrix Composites 

Figure 42 and Figure 43 compare the four-point bend stress-strain curves for several Sylramic 

composites with and without CrMoSi at room temperature and 1645 K, respectively. Data for an 

uninfiltrated preform are included for baseline comparisons. These composites were first slurry 

infiltrated with either 90%SiC-10%Si3N4 or 10%CrMoSi-80%SiC-10%Si3N4 and then melt infiltrated 

with either Si or Si-1(at.%)Ge. Not surprisingly, the composites are stronger than the preforms. 

However, as discussed in Section 4.4.2, the voids in these composites were not completely filled (Figure 

34 and Figure 35). Therefore, the data shown in Figure 42 and Figure 43 must be interpreted with a 

certain degree of caution. 

Figure 40.—Back-scattered electron image and energy dispersive spectra from Regions A, B, C and D for a 

Sylramic preform melt-infiltrated with CrSi2. 
NASA/TM—2014-218352 38 

Figure 41.—Back-scattered electron image and energy dispersive spectra from Regions A, B and C for a Tyranno 

SA preform melt-infiltrated with CrSi2. 

Figure 42.—Room temperature four-point bend tests for engineered 

matrix composites comparing the effects of Ge and CrMoSi 

additions. 
NASA/TM—2014-218352 39 

Figure 43.—Four-point bend tests for engineered matrix composites 

comparing the effects of Ge and CrMoSi additions at 1645 K. 

4.6 Self-Healing Studies 

Composite specimens were pre-cracked and oxidized as discussed in Section 3.4.5. These specimens 

were re-tested in four-point bend tests at 1645 K. Figure 44(a) to (d) compare the bend stress-strain 

curves for these specimens with those shown in Figure 43. These preliminary studies revealed that the 

bend strengths of the pre-cracked specimens significantly increased in almost all instances after 

oxidation. 

Initial assessment of additives on the self-healing characteristics of engineered matrices revealed that 

they were effective in healing scratches and closing a 1 mm diameter hole to various degrees. After 

oxidation at 1600 K for 24 h, CrMoSi-SiC specimen without any additives showed no significant healing 

of the scratches and the hole (Figure 45(a)). However, adding 5(wt.%) CrB2 had a significant effect on 

healing the scratches and partially closing the hole (Figure 45(b)). Continued oxidation at 1700 K for a 

cumulative oxidation time of 100 h showed that the hole on the exposed top face had significantly closed 

relative to that on the rear face (Figure 46(a) to (d)). The presence of 1(wt.%)Ge in the CrMoSi alloy as a 

solid solution was fairly effective in healing the scratches, but it had a smaller effect on closing the hole 

(Figure 45(c)). The amount of Y in the CrMoSi alloy was 0.1 wt.%, which was insufficient to 

significantly heal the scratches (Figure 45(d)). The addition of 5(wt.%) ZrSiO4 was quite effective in 

healing the scratches although the hole showed good closure, the oxide scale was layered and brittle. 
NASA/TM—2014-218352 40 

 

 

Figure 44.—Four-point bend tests for engineered matrix composites at 1645 K comparing the effects of self-healing 

experiments. (a) 90(wt.%) SiC-10%Si3N4-Si; (b) 90(wt.%) SiC-10%Si3N4-SiGe; (c) 10(wt.%) CrMoSi-80% SiC10%Si3N4-Si; (d) 10(wt.%) CrMoSi-80% SiC-10%Si3N4-SiGe. 

Figure 45.—Comparison of the effective of additives on the self-healing of scratches and a 1 mm diameter hole 

before and after oxidation at 1600 K for 24 h. (a) 20(wt.%)CrMoSi-80%SiC; (b) 20%CrMoSi-75%SiC-5%CrB2; 

(c) 20%CrMoSiGe-80%SiC; (d) 20%CrMoSiY-80%SiC; and (e) 20%CrMoSi-75%SiC-5%ZrSiO4. 

(c) (d) 

(a) (b) 
NASA/TM—2014-218352 41 

Figure 46.—Optical macrographs showing (a) and (c) the top face exposed to air and 

(b) and (d) the rear face on the tray bottom showing the extent of hole closure after 

oxidation of a 20%CrMoSi-75%SiC-5%CrB2 at 1700 K for 100 h. Views of the holes 

(a) and (b) before oxidation; and (c) and (d) after oxidation. 

Energy dispersive spectra of the oxide scales formed on these specimens after oxidation at 1700 K 

for a cumulative oxidation time of 100 h revealed that they were silica (Figure 47 to Figure 50). 

However, the morphology of the scale was influenced by the additive. Without any additives, the 

morphology of the scale was granular (Figure 47). Silica globules were observed with the addition of 

5(wt.%) CrB2 (Figure 48). In contrast, Ge addition resulted in a “mud flat” scale morphology thereby 

suggesting a low viscosity scale, which would be advantageous in filling cracks (Figure 49). The effect 

of Y on the scale was similar to that shown in Figure 47 but with more rounded and fused features 

(Figure 50). The specimen with ZrSiO4 resulted in catastrophic oxidation at 1700 K. 

Similarly, 5(wt.%) CrB2 and 5(wt.%) ZrSiO4 additions to 20(wt.%) CrSi2-75%SiC were effective in 

partially closing a 1 mm diameter hole, where the scale for the specimen with ZrSiO4 was brittle (Figure 

51(a) to (d)). The relative change in the hole dimensions, 	d/d0 = (df – d0)/d0, where df and d0 are the final and 

initial diameters of the hole, significantly decreased with the addition of 5(wt. %) of either CrB2 or ZrSiO4

(Figure 52(a) and (b)). For example, the hole dimensions closed by about 60% for CrMoSi/SiC oxidized at 

1700 K for 100 h. However, specimens with ZrSiO4 suffered catastrophic oxidation and the specimens fell 

apart. 
NASA/TM—2014-218352 42 

Notched specimens from compositions identical to the specimens shown in Figure 45(b) to (d) were 

isothermally oxidized at 1700 K for 24 h. Figure 53 shows examples of micrographs of the oxidized and 

unoxidized notches for CrMoSi-SiC-1(wt.%) CrB2 (Figure 53(a) and (c)) and CrMoSi-SiC-5(wt. %) 

CrB2 (Figure 53(b) and (d)), where it is evident that the surfaces have undergone extensive oxidation for 

the specimens with a 5 (wt. %) addition of CrB2. 

Additional compositions were prepared based on the results shown in Figure 45 to Figure 53 with 

additions of sufficient amounts of Si3N4 to compensate for the thermal expansion of CrSi2 and CrMoSi. 

Figure 54 and Figure 55 show the micrographs, energy dispersive spectra and elemental x-ray maps for 

one of these new compositions, which are currently being investigated for self-healing properties. The 

microstructures showed a lot of porosity (Figure 54). The magnitudes of EMOD measured for several 

specimens with a composition of 25(wt. %)CrMoSi-45%SiC-25Si3N4-5%CrB2 increased after they were 

annealed in air at 1698 K for 1800 s compared to those annealed in argon under identical conditions 

(Figure 56). The average values of EMOD for the as-received and specimens annealed in argon were 75 ± 

4.1 GPa and 72.4 ± 4.7 GPa, respectively. In comparison, the average value for the specimens annealed 

in air was 112 ± 4.5 GPa, which represents a 49 percent increase over the as-received values. As shown 

in Figure 56, there was a corresponding increase in the densities of the specimens with the average value 

increasing from 2.62 to 2.80 g/cc for specimens annealed in argon and air, respectively. Fracture 

toughness tests conducted on SEVNB specimens annealed in air and argon at 1698 K for 1800 s were 3.3 

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*-healing of the voids in the 

specimens. Microstructural analyses will be performed on these specimens to confirm this hypothesis. 

Figure 47.—Back-scattered electron images and energy dispersive spectra of the silica scale formed on 

20(wt.%)CrMoSi-80%SiC after oxidation at 1700 K for a cumulative oxidation time of 100 h. 
NASA/TM—2014-218352 43 

Figure 48.—Back-scattered electron images and energy dispersive spectra of 

the silica scale formed on 20%CrMoSi-75%SiC-5%CrB2 after oxidation at 

1700 K for a cumulative oxidation time of 100 h. 

Figure 49.—Back-scattered electron images and energy dispersive spectra of the silica scale formed 

on 20%CrMoSiGe-80%SiC after oxidation at 1700 K for a cumulative oxidation time of 100 h. 
NASA/TM—2014-218352 44 

Figure 50.—Back-scattered electron images and energy dispersive spectra of the silica scale formed on 

20%CrMoSiY-80%SiC after oxidation at 1700 K for a cumulative oxidation time of 100 h. 

Figure 51.—Comparison of the effective of additives on the selfhealing of scratches and a 1 mm diameter hole before and after 

oxidation at 1600 K for 24 h. (a) and (c) 20(wt.%)CrSi2-75%SiC5%CrB2; (b) and (d) 20(wt.%)CrSi2-75%SiC-5%ZrSiO4. 
NASA/TM—2014-218352 45 

Figure 52.—Effective of self-healing additives on the relative change in the hole dimensions as a function of 

oxidation time. (a) CrMoSi/SiC oxidized at 1700 K; (b) CrSi2/SiC oxidized at 1600 K. 

 

Figure 53.—Micrographs of (a) unoxidized and (c) oxidized 20(wt.%)CrMoSi-79%SiC1%CrB2 notched specimens; and (b) unoxidized and (d) oxidized 20(wt.%)CrMoSi75%SiC-5%CrB2 notched specimens. The specimens were oxidized at 1700 K for 24 h. 

(a) (b) 
NASA/TM—2014-218352 46 

Figure 54.—(a) Secondary and (b) and (c) backscattered electron micrographs of 25(wt.%) CrMoSi-45%SiC25%Si3N4-5%CrB2 with corresponding energy dispersive spectra of the phases of a hot-pressed specimen. 

Figure 55.—Secondary electron micrograph and elemental x-ray maps of 25(wt.%) CrMoSi45%SiC-25%Si3N4-5%CrB2. 
NASA/TM—2014-218352 47 

Figure 56.—Plot of Young’s modulus versus density for 25(wt.%) CrMoSi45%SiC-25%Si3N4-5%CrB2 demonstrating an increase in modulus and 

density after annealing in air at 1498 K for 1800 s. The data for the asreceived specimens and those annealed in argon under identical conditions 

are also included. 

5.0 Summary and Conclusions 

A concept for developing a new generation of SiC fiber reinforced engineered matrix composites is 

proposed for applications at 1755 K (2700 °F). The design of the engineered matrices involves several 

steps: Increasing the local plasticity of the matrix in a SiC/SiC composite to blunt cracks; adding 

additives to self-heal cracks and voids; matching the thermal expansion of the matrix with that of the SiC 

fibers; and minimizing the presence of unreacted silicon. Several engineered matrices containing CrSi2, a 

Cr-30(at.%)Mo-30%Si alloy, MoSi2, TiSi2 and WSi2 were designed and formulated by ball milling 

mixtures of the metallic silicides, SiC and Si3N4 powders and hot-pressing them into desired shapes. 

Thermal expansion measurements of the engineered matrices were conducted to demonstrate that they 

matched those of SiC as proof-of-concept. The oxidation behavior of these engineered matrices were 

studied and based on these results, two compositions were down selected for further development. These 

compositions were 10(wt.%)CrSi2-70%SiC-20%Si3N4 and 10(wt.%)CrMoSi-60%SiC-30%Si3N4 for 

1589 K (2400 °F) and 1755 K (2700 °F) applications, respectively. Four-point bend tests conducted on 

these engineered matrix compositions demonstrated that the bend strength increased with increasing 

temperature due to an increase in the local ductility of the matrix. 

Having demonstrated the validity of the proposed concepts, several new engineered matrices were 

designed with additives to enhance their self-healing properties. Early results from self-healing studies 

conducted on these matrices were encouraging. It was demonstrated that additions of 5(wt.%) CrB2 to 

CrSi2-SiC and CrMoSi-SiC matrices were effective in healing scratches and in partially closing 1 mm 

diameter holes. Elastic moduli measurements on a 25(wt.%) CrMoSi-45%SiC-25%Si3N4-5%CrB2

SEVNB specimens showed that the Young’s modulus increased when the specimens were annealed in 

air at 1698 K for 1800 s as compared to specimens annealed in argon. The fracture toughnesses of 
NASA/TM—2014-218352 48 


		
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respectively. 

Attempts to fabricate engineered matrix composites by infiltrating woven preforms proved to be 

challenging. The engineered matrices were infiltrated into SiC/SiC woven BN and CVI SiC-coated 

preforms either using an aqueous or an epoxy carrier media. Despite several attempts, these efforts have 

been largely unsuccessful in infiltrating the preforms with sufficient loading of solids to fill at least 

50 percent of the available voids (~ 20 to 25 vol.%) in the preforms. Future research will concentrate on 

increasing the solids loading in the preforms. A few engineered matrix composites were fabricated by 

particulate and melt infiltrating the preforms, where either Si or a Si-1(at.%)Ge alloy was used for melt 

infiltration, for fabricating specimens for bend testing. The microstructures of these composites revealed 

that the voids were not completely filled. Four-point bend stress-strain data on these composites obtained 

at room temperature and 1645 K are reported. Interestingly, the four-point bend strengths of pre-cracked 

and oxidized specimens at 1645 K were generally higher than the uncracked specimens. 

In an attempt to develop a matrix for 1755 K with little or no unreacted silicon, melt infiltration trial 

runs were conducted on BN-coated Sylramic and Tyranno SA fiber preforms using CrSi2 and a Cr25(wt.%)Si alloy for melt infiltration. Thermodynamic calculations predicted that these alloys would not 

react with either the BN or the CVI SiC coatings. Microstructural observations of infiltrated preforms 

revealed that molten metal had not infiltrated large sections of the preforms. The first trial showed no 

evidence of any reaction between the CrSi2 s and the BN and CVI SiC coatings, and Raman 

spectroscopy of the CrSi2 matrix did not reveal the presence of C and Si. In subsequent melt infiltration 

trials, the microstructural evidence revealed that the CrSi2 and the Cr-25(wt.%)Si had reacted with the 

BN and CVI SiC coatings contrary to the thermodynamic predictions. The reason for this discrepancy is 

unclear and it is being investigated. Developing conditions for melt infiltration with high temperature 

metallic silicides will be continued under the Aerosciences project. Alternative traditional methods of 

fabricating the composites (e.g., tape casting) will be investigated. 

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