An affordable processing technique for the fabrication of silicon carbide-based interpenetrating phase composites (IPCs) is presented. This process consists of the production of microporous carbon preforms and subsequent infiltration with liquid silicon or silicon-refractory metal alloys. The microporous preforms are made by the pyrolysis of a polymerized resin mixture for which methods to control pore volume and pore size have been established. The process gives good control of microstructure and morphology of silicon carbide-based composite materials. Room and high temperature mechanical properties (flexural strength, compressive strength, and flexural creep) of low and high silicon-silicon carbide composites will be discussed

Singh, Mrityunjay

English

NASA Technical Reports Server

AFFORDABLE FABRICATION AND PROPERTIES OF SILICON CARBIDE-BASED
INTERPENETRATING PHASE COMPOSITES
Mrityunjay SINGH
FDC-NYMA, Inc., MS 106-5, NASA Lewis Research Center Group, Cleveland, OH
44135 (USA)
An affordable processing technique for the fabrication of silicon carbide-based
interpenetrating phase composites (IPCs) is presented. This process consists of the
production of microporous carbon preforrns and subsequent infiltration with liquid silicon
or silicon-refractory metal alloys. The microporous preforrns are made by the pyrolysis
of a polymerized resin mixture for which methods to control pore volume and pore size
have been established. The process gives good control of microstructure and
morphology of silicon carbide-based composite materials. Room and high temperature
mechanical properties (flexural strength, compressive strength, and flexural creep) of
low and high silicon-silicon carbide composites will be discussed.
1. INTRODUCTION
Silicon carbide-based ceramic and composite materials have good strength, high
thermal conductivity, and good creep and environmental resistance at high temperatures
[1]. These attractive thermomechanical properties have led to their use in a number of
applications in aeronautics, energy, process, electronic and nuclear industries. In the
aeronautical arena, composite materials are being considered for applications in jet engine
components. Applications in the energy industries include radiant heater tubes, heat
exchangers, heat recuperators, and components for land based turbines for power
generation. These materials are also being considered for use in the first wall and blanket
components of fusion reactors, in fumace linings, and bricks, and in components for
diffusion fumiture (boats, tubes) in the microelectronics industry.
There are a number of critical issues related to commercially available silicon carbide-
based ceramic materials. The majority of these materials are fabricated by a reaction
bonding process which typically consists of molten silicon infiltration into silicon carbide
and carbon containing green bodies or preforms. The resulting materials, also referred to
as reaction bonded silicon carbide (RBSC), contain silicon carbide grains bonded
together by a silicon phase. There is a wide-variation in the microstructural features of
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final materials e.g., interconnectivity of the silicon carbide network, and the morphology
and amount of the silicon phase. These features are detrimental to high temperature
themomechanical properties, and limit the temperature capability of these materials.
Some of these issues can be addressed by using siliconcarbide-based interpenetrating
phase composites in which the silicon carbide phase f_rms a continuous interpenetrating
network through the body with good control over the interconnectivityof the second phase.
In an extensive review, Clarke [2] has discussed various design philosophies and
processing approaches for interpenetrating phase composites. These materials have a
three dimensional microstructure with multifunctional characteristics. A number of silicon
carbide-based interpenetrating phase composites (IPCs) have been produced by reactive
infiltration of interconnected porous carbon preforms [3-6]. These preforms have been
fabricated by the phase separation and pyrolysis process [3-10]. This is a low cost
processing technique which requires fabrication times and/or temperatures significantly
lower than other conventional processing techniques. The materials produced are fully
dense, with excellent control of the pdmary phase microstructure, the amount and
distribution of second phases, and the detrimental effects of sintering aids are avoided [11-
23].
The overall objective of this program is to develop and fabricate affordable silicon
carbide-based IPC materials with tailorable microstructure and therrnomechanical
properties. In this paper, the fabrication approach, microstructure, and mechanical
properties (flexural strength, compressivestrength, and flexural creep) of a wide variety of
silicon carbide-based interpenetratingphase compositeswill be presented.
2. EXPERIMENTAL PROCEDURES
A schematic description of the fabrication process is given in Fig. 1. This process
consists of the production of a microporous carbon preform [7-10] and its subsequent
infiltration with liquid silicon or silicon-refractorymetal alloys. The microporous preforms are
made by the pyrolysis of a polymerized resin mixturewith good control of pore volume and
pore size. During the polymerization process, phase separation of the resin and pore
former takes place. The pore volume and pore size can be controlled by changing the
amount and type of pore formers, curing time, and temperature [8-9]. The interconnectivity
of the porous network is critical in achievingthe desired properties in the composites.
The silicon carbide-based interpenetratingphase composites (IPCs) produced by this
process demonstrate a high degree of control over the primary phase morphology and the
amount and distribution of second phases. Materials with variable volume percent of silicon
can be produced within a continuous network of silicon carbide. The nature and amount of
second phases can be controlled by proper selection of infiltrants and infiltration conditions.
Microstructures of the microporous preforrn.sused for the fabrication of silicon carbide-
based composites are shown in Fig. 2 (a) and (b). Mercury porosimetry data on these
preforms indicate that virtually all the porosity is open to infiltration. All the composite
materials used in this study were fabricated by the reactive infiltrationof molten silicon into
these preforms. The details of carbon preform fabrication [8-9] and other infiltration
conditions have been described elsewhere [4-5]. The control of processing conditions is
critical to avoid microstructural coarsening, silicon vein and lake formations, and cracking
due to thermal expansion mismatch and volume change. After infiltration, samples were
cross-sectioned and polished for metallographic studies. Powder x-ray diffraction analyses
was used to identify different phases in these materials.
Polymer Mixture(Resin, Pore Former, Catalyst),)
Melt Infiltration(Silicon or SI Alloys)
SIC -Based Interpenetratlng_
hase Composites (IPCs) with i
Tailorable Properties J
Fig. 1: Process flow chart for the fabrication of silicon carbide-based materials.
Flexural specimens measured 50 mm long, 4 mm wide and 3 mm thick. Four point
flexural testing was carried out per MIL-STD-1942 (MR) using configuration B specimens
with 20 mm inner and 40 mm outer spans, and a loading rate of 0.5 mm/min [16-19]. At
least, six to ten specimens were tested at room temperature and five specimens were
tested in air at each high temperature. However, in the case of NC-430 only two samples
were tested at each temperature and the data represents the average strength. After
flexure testing, fracture surfaces were examined in optical and scanning electron
microscopes to identify the failure origins.
For the compression test, rectangular samples 2.3 x 2.3 x 4 mm were used. The
specimens were tested in compression at a constant strain rate of 2 xl0 ' s1 at
temperatures ranging from 1250 to 1450 °C in air. Further details of the test conditions are
discussed in other publications [23-24].
(a)
(b)
Fig. 2: Microstructures of porous carbon preforms used for composite fabrication.
Flexural creep tests [22] were conducted with a lever arm machine (Applied Test
Systems, Butler, PA) using a silicon carbide four-point bend fixture. The displacement was
measured using an LVDT with a three-point extensometer made of alumina gage rods with
silicon carbide tips. The silicon carbide tips were in contact with the specimen tensile
surface at the center and beneath the inner load points Relative deflection was recorded
by a data acquisition system. After creep experiments, :he tensile and compression sides
of the crept specimens were examined in SEM.
3. RESULTS AND DISCUSSION
3.1 Microstructure
Optical micrographs of two types of silicon carbide composite materials, designated as
low silicon (LSSC) and high silicon (HSSC), are shown in Fig. 3 (a) and (b). These
microstructures indicate the uniform distribution of silicon phase (white) throughout a silicon
carbide matrix (gray). The silicon contents of low and high silicon composites, shown in Fig.
3 (a) and (b), were estimated to be -14-16 % and -24-26% by x-ray diffraction and image
analyses. For comparison, the microstructure of a reaction bonded silicon carbide
(REFEL-SiC) is shown in Fig. 3 (c). This material has a silicon content similar to the low
silicon-silicon carbide composite (LSSC), but its fabrication route is different (reaction
bonding of SiC grains).
(a).
(b)
(c)
Fig. 3: Microstructure of silicon carbide-silicon composites (a) low silicon (LSSC), (b) high
silicon (HSSC), and (c) REFEL reaction bonded silicon carbide (Si is white and SiC is
gray).
3.2 FLEXURAL STRENGTH
The flexural strengths of low silicon-silicon carbide composite (LSSC) material along
with two types of commercial silicon carbides at different temperatures are given in Table
1. The average room temperature strengthof this composite material is 371 +_28MPa. The
data represents mean results for ten tests at room temperature and five tests at elevated
temperatures. In general, flexural strengths of LSSC are higher than the two commercial
silicon carbides. At 1371°C, the flexural strength of NC-433 reaction bonded silicon carbide
decreases ~ 20% while there is no strength loss in low silicon-silicon carbide composite
(LSSC) materials. The apparent increase in strength of all the specimens at 1100°C may
be due to silica formation and healing of mach.iningdamage. Further experimental studies
are in progress to confirm this phenomena.
Table. 1: Flexural Strength of low silicon-siliconcarbide composite (LSC) and commercial
silicon carbide materials.
Test
Temperature
(oc)
25
Flexural Strength (MPa)
LSSC NC 433 NC430
371+_28 253+_26 146
1100 545±31 274±15 150
1260 422_-1:35 295±15 187
1371 395±30 200-J:23 167
3.3 COMPRESSIVE STRENGTH
The stress-strain curve of constant strain rate experiments for low and high silicon-
silicon carbide composites (LSSC and HSSC) and for REFEL reaction bonded silicon
carbide, are shown in Fig. 4. The experiments were stopped when the stress on the
sample started to decrease. This occurred at strains c,f 4 to 17 %. The strains at which
the decrease in strength initiates are considerably kLrger at lower temperatures. The
deformed samples showed microcracking, but catastrophic crack propagation did not
occur. The compressive strength of low silicon-silicon carbide composite (LSSC) was
considerably higher than the strength of REFEL reaction bonded silicon carbide material
although both materials had quite similar silicon contents. In addition, the compressive
strength of LSSC are also higher than high silicon-silicc,n carbide composites (HSSC).
The stress-strain curves for composites in Fig. 4 show three regions: a linear elastic
region, a non stationary region of plastic deformation where a continuous increase of
stress is necessary in order to advance the deformation of the material at the strain rate
imposed by the experiment, and a third region in which the material accommodates
plastic deformation by formation and coalescence of cavities with a subsequent decrease
in strength.
¢t3
1:3_
v
03
03
oo
_ 1271 RC
/ ! U^ 1448 "C __ 1323 =C
,ool// :ooc
I II IA' ,,o5,c
171 / '"°'°
otf =l / b a
0 5 0 5 10 0 5
Strain (%)
Fig. 4: Plot of compressive stress versus strain for (a) low silicon-silicon carbide composite
(LSSC), (b) high silicon-silicon carbide composite (HSSC), and (c) REFEL-reaction bonded
silicon carbide at high temperatures.
Both silicon carbide interpenetrating phase composites (LSSC and HSSC) used in
this study are composed of a continuous silicon carbide (SIC) network with silicon filling
the pores. This microstructure cannot be deformed without the deformation of both
components, so the total strength and plastic deformation have contributions of all the
phases present. The lower strength of the HSSC over the LSSC material at all
temperatures (in Fig. 4), is related to the lower amount of SiC present (attributed to
different initial carbon preform density), indicating clearly that the strength is dominated
by the silicon carbide network. It is important to point out the high deformation resistance
of the LSSC material, which reaches a value of 500 MPa at temperatures over the
melting point of silicon (Fig. 4), will only be possible when the interpenetrating silicon
carbide network is bearing the load. The compressive strength of REFEL reaction
bonded silicon carbide is lower than low silicon-silicon carbide composite material
(LSSC) due to absence of continuous silicon carbide network and accelerated
deformation of silicon which is used as reaction bonding phase.
3.4 FLEXURAL CREEP
Creep strain vs time plots for three temperatures at different flexural stresses were
generated. Final percent strain was calculated from measurements made on unfailed
test bars. The creep rates and calculated strains are given in Table 2. The comparison of
creep rates of the low silicon-silicon carbide composite (LSSC) with literature data
indicate that the flexural creep resistance of this material is higher than that of most
reaction-bonded silicon carbide materials [24].
Table 2: Status at 100 hrs. for flexural creep tests on low silicon-silicon carbide
composite (LSSC).
Temperature,°C Stress (MPa) Creep rate (s-1) Creep strain (%)
1200 200 1.34 xl0 '9 0.13
1260 200 2.59 xl0 "9 0.18
1320 200 7.8 x10 -9 0.30
1320 150 0.294.43 xl0 '9
The post-test examination of bars revealed the presence of cracks in the surface
silica layers. The microstructures of compression and tension regions were studied in
order to gain insight into the creep mechanism. The scanning electron micrographs of
regions under compressive and tensile stresses show :_ignificant differences. There were
more silicon precipitates on the tensile side of the specimens than on the compression
side. The tensile side had more cavities than the compression side. These cavities
primarily formed at multiple grain junctions in preference to grain boundaries. The
preliminary microscopic data indicate a mechanism ot cavitation during creep [22]. The
detailed TEM analysis of crept specimens is underway and will be reported elsewhere.
4. CONCLUSIONS
The findings of this work suggest that the reaction forming process can be used to
fabricate silicon carbide-based interpenetrating phase composite materials with tailorable
properties. The materials produced by this process are fully dense, with excellent control of
microstructure, second phases, and eliminate the detrimental effects of sintering aids.
These materials have superior high temperature flexural strength, compressive strength,
and creep properties compared to most reaction bonded silicon carbide materials. The
high temperature mechanical properties of composites can be controlled by controlling
the pore volume fraction within the carbon preforms, thus controlling the volume fraction
of the silicon carbide network in the interpenetrating phase composite body.
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Affordable Fabrication and Properties of Silicon Carbide-Based Interpenetrating Phase Composites

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