U.S. patent application number 10/261475 was filed with the patent office on 2005-08-18 for fiber-reinforced ceramic composite material comprising a matrix with a nanolayered microstructure.
Invention is credited to Steffier, Wayne S..
Application Number | 20050181192 10/261475 |
Document ID | / |
Family ID | 46457129 |
Filed Date | 2005-08-18 |
United States Patent
Application |
20050181192 |
Kind Code |
A1 |
Steffier, Wayne S. |
August 18, 2005 |
FIBER-REINFORCED CERAMIC COMPOSITE MATERIAL COMPRISING A MATRIX
WITH A NANOLAYERED MICROSTRUCTURE
Abstract
A fiber-reinforced ceramic matrix composite material exhibiting
increased matrix cracking strength and fracture toughness is
produced by sequentially depositing a plurality of 5-500
nanometer-thick layers of a primary ceramic matrix material phase
periodically separated by 1-100 nanometer-thick intermediate layers
of a secondary matrix material phase onto the reinforcing fibers
upon their consolidation. The resultant nanolayered matrix enhances
the resistance to the onset of matrix cracking, thus increasing the
useful design strength of the ceramic matrix composite material.
The nanolayered microstructure of the matrix constituent also
provides a unique resistance to matrix crack propagation. Through
extensive inter-layer matrix fracture, debonding and slip, internal
matrix microcracks are effectively diverted and/or blunted prior to
their approach towards the reinforcing fiber, thus increasing the
apparent toughness of the matrix constituent. This unique
toughening mechanism serves to dampen energetic co-planar
macrocrack propagation typically observed in conventionally
manufactured ceramic matrix composites wherein matrix cracks are
usually deflected at the fiber/matrix interphase region.
Inventors: |
Steffier, Wayne S.;
(Huntington Beach, CA) |
Correspondence
Address: |
MORLAND C FISCHER
2030 MAIN ST
SUITE 1050
IRVINE
CA
92614
|
Family ID: |
46457129 |
Appl. No.: |
10/261475 |
Filed: |
September 30, 2002 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
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10261475 |
Sep 30, 2002 |
|
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09764809 |
Jan 16, 2001 |
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Current U.S.
Class: |
428/293.4 |
Current CPC
Class: |
C04B 2235/428 20130101;
C04B 35/62897 20130101; Y10T 428/249945 20150401; C04B 2237/368
20130101; Y10T 428/249931 20150401; D04H 1/42 20130101; Y10T
428/249924 20150401; B32B 18/00 20130101; C04B 2235/77 20130101;
Y10T 428/24993 20150401; Y10T 428/2933 20150115; Y10T 428/249928
20150401; C04B 2235/5256 20130101; C04B 2235/5244 20130101; C04B
35/80 20130101; C04B 35/62894 20130101; C04B 2235/5268 20130101;
C04B 2237/365 20130101; C04B 2237/361 20130101; C04B 2235/96
20130101; Y10T 428/24994 20150401; C04B 2237/363 20130101; C04B
2235/80 20130101; C04B 35/62868 20130101; Y10T 428/2918 20150115;
Y10T 428/2938 20150115; C04B 35/565 20130101; C04B 35/62884
20130101; C04B 2237/38 20130101; C04B 2237/704 20130101; C04B
35/62873 20130101; C04B 2235/614 20130101 |
Class at
Publication: |
428/293.4 |
International
Class: |
B32B 017/12 |
Claims
1-12. (canceled)
13. A method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material comprising: fabricating a fibrous
preform of refractory reinforcing fibers; depositing a fiber
coating material which fully encapsulates the refractory fibers
thereof; and depositing a ceramic matrix material with a
nanolayered microstructure comprising a plurality of primary phase
layers of a first material and a plurality of secondary phase
layers of a second material, wherein the secondary phase layers are
interposed between the primary phase layers, and wherein the
nanolayered ceramic matrix material encapsulates the coated
refractory fibers of the fibrous preform and consolidates the
preform into a densified composite.
14. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein said
fibrous preform is fabricated from an assemblage of refractory
fibers produced by a textile fabrication process selected from a
group of fabrication processes comprising weaving, braiding,
knitting, fiber placement, filament winding, felting, and
needling.
15. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein the
fiber coating deposited on said fibrous preform has a thickness of
0.05-5.0 micrometers and is produced by a process selected from the
group of fiber coating processes comprising chemical vapor
infiltration (CVI), polymer precursor impregnation/pyrolysis (PIP),
reaction formation, and combinations thereof.
16. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein the
fiber coating deposited on said fibrous preform is carbon produced
by chemical vapor infiltration using a carbon-forming precursor
selected from a group of chemical precursors comprising methane,
propane, propylene, and mixtures thereof, which is pyrolytically
decomposed into carbon at an elevated temperature of
950-1250.degree. C. and at a reduced pressure of 1-250 Torr.
17. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein the
fiber coating deposited on said fibrous preform is boron nitride
produced by chemical vapor infiltration using a boron
nitride-forming precursor selected from a group of chemical
precursors comprising boron trichloride, boron triflouride,
diborane, and mixtures thereof, which is reduced with a reductant
selected from a group of chemical reductants comprising nitrogen,
hydrogen, ammonia, and mixtures thereof to form boron nitride at an
elevated temperature of 700-1200.degree. C. and at a reduced
pressure of 1-250 Torr.
18. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein said
fibrous preform is first coated with carbon followed by a boron
carbide coating produced by chemical vapor infiltration using a
boron carbide-forming precursor selected from a group of chemical
precursors comprising boron trichloride, boron triflouride,
diborane, and mixtures thereof, which is reacted with a
carbon-forming precursor selected from a group of chemical
reactants comprising methane, propane, propylene, and mixtures
thereof to form boron carbide at an elevated temperature of
800-1100.degree. C. and at a reduced pressure of 1-250 Torr.
19. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein said
fiber coated fibrous preform is consolidated with a nanolayered
ceramic matrix material produced by a process selected from a group
of matrix consolidation processes comprising chemical vapor
infiltration (CVI), polymer precursor impregnation/pyrolysis (PIP),
reaction formation, and combinations thereof, which fully
encapsulates said coated reinforcing fibers of said fibrous preform
for transforming said fibrous preform into a dense, ceramic matrix
composite material.
20. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein said
nanolayered ceramic matrix material is manufactured by the steps
comprising: depositing a first layer of a primary phase material
with a thickness of between 5 and 500 nanometers; depositing a
second layer of a secondary phase material with a thickness of
between 1 and 100 nanometers; and sequentially repeating steps
(a-b) to create a material with a nanolayered microstructure
comprising a plurality of primary phase layers of a first material
and a plurality of secondary phase layers of a second material, the
secondary phase layers interposed between the primary phase layers,
wherein the nanolayered ceramic matrix material encapsulates the
coated refractory fibers of the fibrous preform and consolidates
the preform into a densified composite.
21. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 20, wherein said
primary phase material layers and said secondary phase material
layers comprising said nanolayered ceramic matrix material are
produced by chemical vapor infiltration.
22. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 20, wherein said
primary phase layers and said secondary phase layers comprising the
nanolayered ceramic matrix material are produced by periodic
throttling of the respective chemical precursor materials at
discrete time intervals during chemical vapor infiltration matrix
deposition processing, whereby the selected throttling frequencies
are dependent on the respective primary and secondary material
deposition rates.
23. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 20, wherein the
chemical vapor infiltration process for depositing said primary
phase layers and said secondary phase layers comprising said
nanolayered ceramic matrix material is maintained at a constant
temperature, whereby the selected processing temperature is
dependent on the compatibility for producing both said primary and
said secondary material phases.
24. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 20, wherein the
chemical vapor infiltration process for depositing said primary and
said secondary phase layers comprising said nanolayered ceramic
matrix material is maintained at a constant pressure, whereby the
selected processing pressure is dependent on the compatibility for
producing both said primary and said secondary material phases.
25. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein said
plurality of primary phase layers comprising said nanolayered
ceramic matrix are silicon carbide produced by chemical vapor
infiltration using a silicon carbide-forming precursor selected
from a group of chemical precursors comprising
methyltrichlorosilane, dimethyldichlorosilane, silicon
tetrachloride with methane, and mixtures thereof, which is reacted
to form silicon carbide at an elevated temperature of
850-1150.degree. C. and at a reduced pressure of 1-250 Torr.
26. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein said
plurality of primary phase layers comprising said nanolayered
ceramic matrix are silicon nitride produced by chemical vapor
infiltration using a silicon nitride-forming precursor selected
from a group of chemical precursors comprising silicon
tetrachloride, silicon tetraflouride, dichlorosilane,
trichlorosilane, and mixtures thereof, which is reduced with a
reductant selected from a group of chemical reductants comprising
ammonia, nitrogen, hydrogen, and mixtures thereof to form silicon
nitride at an elevated temperature of 800-1100.degree. C. and at a
reduced pressure of 1-250 Torr.
27. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein said
plurality of primary phase layers comprising said nanolayered
ceramic matrix are boron carbide produced by chemical vapor
infiltration using a boron carbide-forming precursor selected from
a group of chemical precursors comprising boron trichloride, boron
triflouride, diborane, and mixtures thereof, which is reacted with
a carbon-forming precursor selected from a group of chemical
reactants comprising methane, propane, propylene, and mixtures
thereof to form boron carbide at an elevated temperature of
800-1100.degree. C. and at a reduced pressure of 1-250 Torr.
28. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein said
plurality of secondary phase layers comprising said nanolayered
ceramic matrix are carbon produced by chemical vapor infiltration
using a carbon-forming precursor selected from a group of chemical
precursors comprising methane, propane, propylene, and mixtures
thereof, which is pyrolytically decomposed into carbon at an
elevated temperature of 950-1250.degree. C. and at a reduced
pressure of 1-250 Torr.
29. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein said
plurality of secondary phase layers comprising the nanolayered
ceramic matrix are silicon produced by chemical vapor infiltration
using a silicon-forming precursor selected from the group of
chemical precursors comprising dichlorosilane, trichlorosilane and
mixtures thereof, which is reduced with hydrogen to form silicon at
an elevated temperature of 950-1250.degree. C. and at a reduced
pressure of 1-250 Torr.
30. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein said
plurality of secondary phase layers comprising the nanolayered
ceramic matrix are boron nitride produced by chemical vapor
infiltration using a boron nitride-forming precursor selected from
a group of chemical precursors comprising boron trichloride, boron
triflouride, diborane, and mixtures thereof, which is reduced with
a reductant selected from a group of chemical reductants comprising
nitrogen, hydrogen, ammonia, and mixtures thereof to form boron
nitride at an elevated temperature of 700-1200.degree. C. and at a
reduced pressure of 1-250 Torr.
31. The method for manufacturing a fiber-reinforced nanolayered
ceramic matrix composite material recited in claim 13, wherein said
plurality of secondary phase layers comprising the nanolayered
ceramic matrix are silicon nitride produced by chemical vapor
infiltration using a silicon nitride-forming precursor selected
from a group of chemical precursors comprising silicon
tetrachloride, silicon tetraflouride, dichlorosilane,
trichlorosilane, and mixtures thereof, which is reduced with a
reductant selected from a group of chemical reductants comprising
ammonia, nitrogen, hydrogen, and mixtures thereof to form silicon
nitride at an elevated temperature of 800-1100.degree. C. and at a
reduced pressure of 1-250 Torr.
32. A fiber-reinforced ceramic matrix composite material having
enhanced matrix cracking strength, comprising: a fibrous preform of
refractory fibers; a fiber coating material which fully
encapsulates the refractory fibers of said fibrous perform; and a
high-strength, nanolayered ceramic matrix which fully encapsulates
and consolidates the refractory fibers of said fibrous preform into
a densified composite, said nanolayered ceramic matrix having a
microstructure comprising a plurality of primary phase layers being
formed from a first material and having a thickness of between 5
and 500 nm, and a plurality of secondary phase layers being formed
from a second material and having a thickness of between 1 and 100
nm, wherein said plurality of secondary phase layers are interposed
between said plurality of primary phase layers, such that the
tensile matrix cracking strength of said ceramic matrix composite
material is greater than 100 Mpa, said fiber coating material
debonding said nanolayered ceramic matrix from the refractory
fibers of said fibrous perform.
33. The fiber-reinforced ceramic matrix composite material recited
in claim 32, wherein the first material from which said plurality
of primary phase layers of said nanolayered ceramic matrix formed
is selected from a group of materials consisting of silicon,
silicon carbide, boron carbide, tantalum carbide, hafnium carbide,
zirconium carbide, silicon nitride, tantalum nitride, hafnium
nitride, zirconium nitride, titanium nitride, silicon boride,
tantalum boride, hafnium boride, zirconium boride, titanium boride,
zirconium silicide, titanium silicide, molybdenum silicide,
aluminum oxide, silicon oxide, hafnium oxide, zirconium oxide, and
titanium oxide.
34. The fiber-reinforced ceramic matrix composite material recited
in claim 32, wherein the second material from which said plurality
of secondary phase layers of said nanolayered ceramic matrix is
formed is selected from a group of materials consisting of carbon,
silicon, silicon carbide, boron carbide, tantalum carbide, hafnium
carbide, zirconium carbide, silicon nitride, boron nitride,
tantalum nitride, hafnium nitride, zirconium nitride, titanium
nitride, aluminum nitride, silicon boride, tantalum boride, hafnium
boride, zirconium boride, titanium boride, zirconium silicide,
titanium silicide, molybdenum silicide, aluminum oxide, silicon
oxide, hafnium oxide, zirconium oxide, and titanium oxide.
35. The fiber-reinforced ceramic matrix composite material recited
in claim 32, wherein each of the plurality of primary phase layers
of said nanolayered ceramic matrix is formed from the same
material.
36. The fiber-reinforced ceramic matrix composite material recited
in claim 32, wherein each of the plurality of secondary phase
layers of said nanolayered ceramic matrix is formed from the same
material.
37. The fiber-reinforced ceramic matrix composite material recited
in claim 32, wherein the first material of said plurality of
primary phase layers of said nanolayered ceramic matrix is a
crystalline material, the second material of said plurality of
secondary phase layers being capable of interrupting the growth of
the crystallites of the first crystalline material of said primary
phase layers.
38. The fiber-reinforced ceramic matrix composite material recited
in claim 32, wherein the refractory fibers of said fibrous preform
are selected from a group of fiber materials consisting of silicon
carbide, silicon nitride, aluminum oxide, and any other matrix
thermal expansion-compatible fiber material that is capable of
withstanding a temperature in excess of 800.degree. C.
39. The fiber-reinforced ceramic matrix composite material recited
in claim 32, wherein the fiber coating material which encapsulates
the refractory fibers of said fibrous perform is a coating having a
thickness of between 0.05 and 5.0 micrometers and being selected
from a group of fiber coating materials consisting of carbon,
silicon carbide, boron carbide, tantalum carbide, hafnium carbide,
zirconium carbide, silicon nitride, boron nitride, tantalum
nitride, hafnium nitride, zirconium nitride, titanium nitride,
aluminum nitride, silicon boride, tantalum boride, hafnium boride,
zirconium boride, titanium boride, zirconium silicide, titanium
silicide, molybdenum silicide, aluminum oxide, silicon oxide,
hafnium oxide, zirconium oxide, and titanium oxide.
40. The fiber-reinforced ceramic matrix composite material recited
in claim 39, wherein said fiber coating material consists of a
single-layer phase of uniform composition.
41. The fiber-reinforced ceramic matrix composite material recited
in claim 39, wherein said fiber coating material consists of a
single-layer phase of mixed composition.
42. The fiber-reinforced ceramic matrix composite material recited
in claim 39, wherein said fiber coating material consists of a
multilayered phase including alternating coating layers having
respective fiber coating compositions.
43. A fiber-reinforced ceramic matrix composite material having
enhanced matrix cracking strength, comprising: a fibrous preform of
refractory fibers; a high-strength, nanolayered ceramic matrix
which fully encapsulates the refractory fibers of said fibrous
preform into a densified composite, said nanolayered ceramic matrix
having a microstructure including a plurality of primary phase
layers being formed from the same first crystalline material and a
plurality of secondary phase layers being formed from the same
second material, said second material being different from said
first material and being capable of interrupting the growth of the
crystallites of said first crystalline material, and said plurality
of secondary phase layers being interposed between said plurality
of primary phase layers such that the tensile matrix cracking
strength of said ceramic matrix composite material is greater than
100 Mpa; and a debonding layer to debond said nanolayered ceramic
matrix from the refractory fibers of said fibrous perform.
44. The fiber-reinforced ceramic matrix composite material recited
in claim 43, wherein the first crystalline material from which said
plurality of primary phase layers are formed is silicon carbide.
Description
CROSS-REFERENCE TO RELATED APPLICATIONS
[0001] This application is a continuation-in-part of application
Ser. No. 09/764,809 filed Jan. 16, 2001.
FIELD OF THE INVENTION
[0002] The present invention relates to a ceramic matrix composite
material composed of a refractory fiber reinforcement, a fiber
coating or fiber coating system, and a nanolayered ceramic matrix
having increased matrix cracking strength and methods of producing
same.
BACKGROUND OF THE INVENTION
[0003] Fiber-reinforced ceramic matrix composite materials are
actively being developed for a variety of high-temperature
military, aerospace and industrial applications. While possessing
high specific strength and toughness, the utility of current
ceramic matrix composites are severely limited by their
susceptibility to oxidation embrittlement and strength degradation
when stressed at or beyond their matrix cracking strength and
exposed to high-temperature oxidation. Thus, for the current state
of technology, the linear-elastic region represents the "useful"
design stress-strain region due to the negative effects caused by
environmental degradation of the fiber coating and/or reinforcing
fiber at elevated temperatures following the onset of matrix
cracking.
[0004] Ceramic materials have long been considered potentially
beneficial for hot structural component applications in advanced
gas turbine and rocket engines, and future high-speed aircraft and
atmospheric re-entry vehicles. In general, ceramics have superior
high-temperature strength and modulus while having a lower density
than metallic materials. The principal disadvantages of ceramics as
structural materials are their low failure strain, low fracture
toughness and catastrophic brittle failure characteristics. Because
of these inherent limitations, monolithic ceramics lack the
properties of reliability and durability that are necessary for
structural design acceptance. The emerging technology of
fiber-reinforced ceramics, or ceramic matrix composites is one
promising solution for overcoming the reliability and durability
problems associated with monolithic ceramics. By incorporating
high-strength, relatively high-modulus fibers into brittle ceramic
matrices, combined high strength and high toughness composites can
be obtained. Successfully produced ceramic matrix composites
exhibit a high degree of non-linear stress-strain behavior with
ultimate strengths, failure strains and fracture toughness that are
substantially greater than that of the otherwise brittle ceramic
matrix.
[0005] In order to exploit the benefits of fiber reinforcement in
brittle ceramic matrices, it is well recognized that relatively
weak fiber/matrix interfacial bond strength is essential for
preventing catastrophic failure from propagating matrix cracks. The
interface must provide sufficient fiber/matrix bonding for
effective load transfer, but must be weak enough to debond and slip
in the wake of matrix cracking, leaving the fibers to bridge the
cracks and support the far-field applied load. Fiber-reinforced
ceramic matrix composites with very high fiber/matrix interfacial
bond strengths (usually the result of chemical interaction during
manufacture) exhibit brittle failure characteristics similar to
that of unreinforced monolithic ceramics by allowing matrix cracks
to freely propagate directly through the reinforcing fibers. By
reducing the interfacial bond strength, the fiber and matrix are
able to debond and slip, thereby promoting the arrest and/or
diversion of propagating matrix cracks at/or around the reinforcing
fiber. Since crack inhibition/fracture toughness enhancement is the
primary advantage of fiber-reinforced ceramic matrix composites,
properly engineered fiber coating systems are thus essential for
improving the structural performance of these materials. Control of
interfacial bonding between the fiber and matrix following
manufacture and during service is typically provided by the use of
applied fiber coatings.
[0006] Fiber-reinforced ceramic matrix composites produced by the
chemical vapor infiltration (CVI) process are a particularly
promising class of engineered high-temperature structural
materials, which are now commercially available. The principal
advantage of the CVI process approach for fabricating ceramic
composites as compared to other manufacturing methods (e.g.,
reaction bonding, hot-pressing, melt infiltration, or polymer
impregnation/pyrolysis) is the ability to infiltrate and densify
geometrically complex, multidirectional fiber preforms to
near-net-shape with a ceramic matrix of high purity and
controllable stoichiometry without chemically, thermally or
mechanically damaging the relatively fragile reinforcing fibers. In
addition, because it is a relatively low temperature manufacturing
process, high purity refractory matrix materials can be formed
(deposited) at a small fraction of their melting temperature
(.about.T.sub.m/4). Despite the many possible high-temperature
ceramic matrix composite systems, however, the number of practical
systems is limited by the currently available reinforcing fibers.
To date, the majority of high performance ceramic matrix composites
produced have primarily been fabricated using carbon and
polymer-derived SiC (Nicalon and Hi-Nicalon) fiber reinforcement
and CVI-derived SiC matrices.
[0007] Carbon fibers offer the highest temperature capability of
all current commercially available refractory fibers. Carbon
fiber-reinforced SiC ceramics (C/SiC), however, are susceptible to
severe strength degradation when exposed to high-temperature
oxidizing environments for prolonged periods. This limitation is
due to the extensive process-induced matrix microcracking resulting
from the relatively large thermal expansion mismatch between the
carbon fiber reinforcement and the surrounding SiC matrix. The
resultant matrix cracks provide access to environmental intrusion,
particularly oxidation, which accelerates the degradation of the
compliant fiber coating (e.g., pyrolytic carbon and boron nitride)
and the reinforcing fiber. Commercially available small diameter
(.about.15 .mu.m) ceramic fibers such as Nicalon and Hi-Nicalon
microcrystalline SiC, although having limited elevated temperature
capability (<1200.degree. C. and <1400.degree. C.,
respectively) as compared to carbon fiber, exhibit excellent
thermomechanical compatibility with SiC matrices. These fibers thus
produce composites which are not initially microcracked. Although
these ceramic fibers are more oxidation resistant than carbon
fibers, the resultant composites also experience irreversible
oxidation embrittlement and strength degradation when stressed at
or beyond their matrix cracking strength and subsequently exposed
to high-temperature oxidation.
[0008] Unlike the near-linear tensile stress-strain behavior of the
microcracked C/SiC material system, SiC fiber-reinforced SiC matrix
composites (SiC/SiC) exhibit highly nonlinear stress-strain
characteristics; controlled by the low matrix failure strain
relative to the reinforcing fiber. As the composite is loaded in
tension, it deforms linear-elastically up to the onset of matrix
cracking. The tension threshold at which the onset of matrix
cracking occurs designates the "proportional limit" of the
material. That is, when the applied tensile strain reaches the
failure strain of the unreinforced matrix, ideally assuming
negligible residual thermoelastic effects, transverse matrix cracks
initiate and propagate rapidly across the composite, leaving the
fibers to bridge the cracks while supporting the far-field applied
load. Continued loading beyond the onset of matrix cracking results
in the formation of many regularly spaced matrix cracks (i.e.,
multiple matrix cracking) typically accompanied by a nonlinear
decrease in composite stiffness. This strain-induced compliance
behavior is the result of the diminishing contribution of the
matrix modulus with increased multiple matrix cracking and
fiber/matrix debonding. The diminishing stiffness behavior becomes
more significant with increasing applied strain to the point where
the elastic modulus of the composite is primarily dominated by the
reinforcing fibers. As the composite strain approaches the fiber
failure strain, the fibers progressively fracture, designating the
ultimate strength of the composite. For most practical structural
applications, however, the linear-elastic region represents the
"useful" design stress-strain region due to the negative effects of
hysteresis and environmental degradation of the fiber coating
and/or fiber reinforcement at elevated temperatures occurring after
matrix microcracking. Matrix microcracking is therefore a
fundamental life-limiting issue for ceramic matrix composites being
considered for use in extended-life thermostructural
applications.
[0009] From an engineering mechanics standpoint, the high elastic
modulus of the CVI-derived SiC matrix relative to the reinforcing
fiber is a disadvantage for load transfer. For a Nicalon SiC
fiber-reinforced/CVI SiC matrix composite with a 40 volume-percent
fiber loading, nominally 70% of the applied load is carried by the
matrix prior to the onset of matrix cracking. For an equal
volumetric loading of higher modulus Hi-Nicalon SiC fiber, about
60% of the applied load is initially carried by the matrix. The
relatively high matrix stiffness is thus a disadvantage from the
standpoint of matrix microcracking. The tensile strain at which the
onset of matrix cracking occurs in both Nicalon and Hi-Nicalon
reinforced SiC matrix composites is typically .about.0.04%, and
rarely exceeds 0.05%, as this is an intrinsic property of the
CVI-derived brittle matrix. The corresponding tensile matrix
cracking strengths typically range between 60 and 80 MPa,
respectively. In turn, the useful design strengths for both Nicalon
and Hi-Nicalon SiC fiber-reinforced/CVI SiC composites are only
about 30% of their respective fiber-dominated ultimate strengths.
Thus, the early onset of matrix microcracking and subsequent
oxidation embrittlement and strength degradation is a primary
performance limitation of current state-of-the-art materials.
Fundamentally, it would be desirable for a fiber-reinforced ceramic
composite material to have a useful design strength significantly
greater than 30% of its ultimate strength.
[0010] Although the oxidation embrittlement problem in
fiber-reinforced ceramic matrix composite materials may eventually
be controlled via advanced fiber coating and/or matrix oxidation
inhibition approaches, it will nevertheless still be desirable to
increase the elastic limit of the composite to reduce potential
fatigue, hysteresis and other complex nonlinear material behavioral
effects. Once the composite elastic limit is exceeded, the
structural designer is faced with using a nonlinear and potentially
time-dependent microcracked material system. One, however,
nontrivial approach towards increasing the matrix cracking strength
in ceramic matrix composites is by increasing the mechanical
properties (e.g., strength and fracture toughness) of the matrix
constituent. This approach has been successfully demonstrated by
the current inventor via microstructural engineering of the SiC
matrix into a strong/tough nanolayered composite constituent.
[0011] A nanolayered composite comprises a compositionally
modulated microlaminate consisting of periodically alternating
layers (lamellae) of two or more material constituents. The
thickness of each layer range from about one molecular monolayer
(.about.1 nm) to a thickness approaching the upper limit of very
fine grain refinement achievable from current state-of-the-art
materials processing techniques (>150 nm). These materials can
be engineered to exhibit remarkable mechanical, tribological,
thermal, and/or electrical properties that are uniquely different
from those of the individual constituents. In particular, strength
can be enhanced over currently available courser grained materials
by an order of magnitude or more. Also of importance is that a
conceivably wide range of refractory metal and ceramic materials
can be engineered into such nanostructural composites suitable for
extreme environmental structural applications.
[0012] Early efforts (over two decades ago) by researchers at the
Chemetal/San Fernando Laboratories (SFL) led to the discovery of a
unique form of chemical vapor deposited SiC. While attempting to
deposit "massive" bodies of SiC, unanticipated thermochemical
process instabilities (i.e., chugging) within the "cold-wall"
chemical vapor deposition (CVD) reactor resulted in producing a
material with an unusual layered microstructure. This material was
found to be composed of alternating lamellae of SiC and elemental
silicon (Si), ranging in thickness from 10 to 20 nm and 1 to 2 nm,
respectively. Reported properties for this SiC/Si material included
flexural strengths, elastic moduli, fracture toughness, and
hardness which exceeded 4000 MPa, 450 GPa, 6-12 MPa{square root}m,
and 45 GPa, respectively. Numerous subsequent evaluations by
government laboratory scientists Dutta, Graham, Rice, and
Mendiratta confirmed these astonishing results. Dutta, S., R. Rice,
H. Graham, and M. Mendiratta, Characterization and Properties of
Controlled Nucleation Thermochemical Deposited (CNTD) Silicon
Carbide, NASA TECH. MEMO. 79277, presented at the 80th Annual
Meeting of the American Ceramic Society. This work resulted in the
issuance of a number of domestic and foreign patents for which the
process was coined "Controlled Nucleation Thermochemical
Deposition", or CNTD.
[0013] Despite the extraordinary mechanical and physical properties
of this termed "ultra-structured" material, however,
commercialization was hindered by problems of reproducibility. In
short, processing difficulties associated with the inability to
control the naturally occurring chemical instability within the
cold-wall reactor during deposition prevented this product from
becoming commercially successful. Specifically, the uncontrollable,
and not well-understood "cyclic" instability phenomenon was not
easily scaled to larger or hot-wall reactors, resulting in low
yield, poor reproducibility and poor process economics.
Accordingly, it would be desirable to be able to artificially
reproduce the beneficial effects of this CNTD process such that it
could be effectively used in hot-wall CVD reactors (which can
accommodate large batch quantities of dissimilar parts) by
relatively simple and controllable mechanical means.
[0014] It is also known that layered fiber coatings can be applied
to fibrous preforms in such a way to increase the oxidation
resistance of the resultant ceramic matrix composite, while
preserving desirable strength and toughness. Particularly, by
controlling the flow of the chemical precursors (i.e., chemically
reactant precursor gases or gasified liquids) during application of
a ceramic coating onto a preform of refractory fibers, wherein two
or more independent chemical precursors are periodically turned on
and off at prescribed intervals, it has been shown that the
resultant microlayered coating produced creates an inherently
oxidation-resistant fiber coating material. FIG. 1 shows a
backscattered electron image (BEI) of an advanced "multilayered SiC
fiber coating" developed by the current inventor and applied using
the technique of cyclic "throttling" of the chemical precursors
during the deposition of the fiber coating. Details of this
multilayered fiber coating and processing method is more fully
described in U.S. Pat. Nos. 5,455,106 and 5,545,435, the
disclosures of which are incorporated herein by reference. The
microlayered fiber coating system (deposited on a .about.15 .mu.m
Hi-Nicalon SiC fiber) shown in FIG. 1 was engineered to mitigate
the inherent problems of oxidation resistance plaguing currently
available PyC and BN fiber coatings for structural ceramic matrix
composites. This was achieved by successfully tailoring the desired
mechanical characteristics (e.g., interfacial shear strength and
compliance) of the multilayered SiC coating system via
microstructural engineering necessary to enhanced strength and
toughness of the resultant composite.
[0015] These patent disclosures, however, did not address the
problem of increasing the strength and/or toughness of the matrix
constituent itself and were directed instead to the product and
method of depositing the oxidation-resistant multilayered ceramic
fiber coating material. These patent disclosures describe
depositing microlayers having a primary layer thickness of between
500 and 5000 nanometers, which is considered too thick to increase
the inherent strength of the material produced as will be described
later. Accordingly, it would be desirable to increase the strength
and toughness of the ceramic matrix constituent in order to enhance
the matrix cracking resistance in the resulting composite.
[0016] To better understand the unique behavior of CVI/CVD
nanostructures, it is useful to briefly describe the morphology of
these engineered materials. Nanolayered composites are produced by
depositing a layer of the primary, or major constituent species
with a thickness of on the order of a few tens of nanometers
(10-100 nm), followed by a layer of minor species with a thickness
of about an order of magnitude less (1-10 nm). Deposition durations
for each layer are very short, ranging from a few seconds to a few
tenths of a second. The exact deposition durations are dependent on
the deposition rates of the respective major and minor species
derived from a given process. The process is then cyclically
repeated until the desired thickness of the body is achieved.
Although there may appear to be great flexibility in the selection
of the secondary nanolayering constituent(s), they must be
carefully selected based on their (1) known ability to effectively
interrupt the deposition epitaxy of the major constituent, thereby
increasing strength and thermal shock resistance by controlling
grain refinement; and (2) propensity to provide beneficial elastic
modulus mismatch, thereby further increasing fracture toughness by
limiting dislocation motion.
[0017] In conventional CVI/CVD-deposited materials, high-purity
crystallites nucleate on the heated substrate surface (e.g., part
to be coated or fibers) and then grow epitaxially in a direction
perpendicular (i.e., normal) to the heated substrate; most often
through the entire thickness of the deposit. The crystallites thus
coarsen and weaken with increasing thickness during the growth
process. In the current invention (e.g., CVI/CVD nanolayering
process), the major disadvantages of conventional CVI/CVD are
eliminated. The first crystallites nucleate and start to grow,
competing for a preferred minimum energy orientation. Before they
ever have a chance to become oriented, the growth is interrupted by
the deposition of the second material. This secondary layer is
deposited so thinly that its crystallites do not have a chance to
grow, and thus do not achieve any preferred orientation. When the
cycle is repeated, the crystallites of the primary material must
re-nucleate and the process of nucleation/interruption is repeated;
thus, the grains in the deposited material never have a chance to
coarsen. When engineered successfully, this process has been shown
to result in producing materials with significantly increased
strength and hardness; beyond that predicted by the teachings of
Hall-Petch. Crystals of about 5 .mu.m are considered fine by most
materials scientists and engineers; and those of 1 .mu.m, extremely
fine. Virtually no structural components have ever been produced by
conventional methods with grains less than 0.4 .mu.m (400 nm) in
size. The CVI/CVD nanolayering process developed in the current
invention provides the ability to produce highly uniform
microstructures with grain sizes of between 1 and 100 nm. Although
engineering properties are improved by grain refinement in general,
it is not until the crystallites are maintained to less than
.about.150 nm that dramatic improvements to near-theoretical
strengths are observed.
[0018] FIG. 2 shows a microstructural example of a nanolayered CVD
SiC material developed by the current inventor. A tensile strength
enhancement of nearly one (1) order of magnitude (.about.8.times.)
over conventional CVD SiC has been experimentally demonstrated.
SUMMARY OF THE INVENTION
[0019] The present invention comprises a novel material and
manufacturing process for fabricating the same. The materials and
process described herein results in producing a fiber-reinforced
ceramic matrix composite material with increased resistance to
matrix cracking. According to the invention, a matrix constituent
with a nanolayered microstructure exhibiting increased strength and
toughness is produced by sequentially depositing a plurality of
thin layers of a primary ceramic matrix phase, interposed by a
plurality of very thin intermediate layers of a secondary matrix
phase onto the reinforcing fibers upon their consolidation.
[0020] Fiber-reinforcement is defined as any refractory fibers,
continuous or discontinuous, used for producing a fibrous preform
texture, which are capable of withstanding a use temperature of at
least 800.degree. C. in an atmosphere which is thermochemically
compatible with that fiber without suffering fundamental chemical,
physical or mechanical degradation. Examples include carbon fibers,
silicon carbide fibers, silicon nitride fibers, aluminum oxide
fibers, etc.
[0021] A fiber preform is a fibrous texture defined as any
assemblage of one or more reinforcing fiber types produced by
weaving, braiding, filament winding, fiber placement, felting,
needling, or other textile fabrication process.
[0022] Fiber preforming is a textile fabrication process by which
the collimated multifilamentary fiber bundles (i.e., tows) are
placed and maintained in position for purposes of controlling both
their orientation and content within a given volumetric space. As
such, we will refer to this spatial arrangement of fibers as a
preform architecture.
[0023] Fiber coating is defined as any refractory composition of
either carbon, metal carbide, metal nitride, metal boride, metal
silicide, metal oxide, or combinations thereof which is (are)
deposited (for example by chemical vapor infiltration) onto the
refractory fibers either before or after fiber preforming for
purposes of controlling the fiber/matrix interfacial bonding
characteristics in the resultant composite. The resultant fiber
coating thus encapsulates the reinforcing fibers. Examples include
pyrolytic carbon, silicon carbide, silicon nitride, boron carbide,
boron nitride, etc.; either as a single-layer phase, multilayered
phase or as a phase of mixed composition. In one embodiment, the
fiber coating system has a thickness of 0.05-5.0 micrometers.
[0024] Ceramic matrix is defined as any refractory composition of
either carbon, metal carbide, metal nitride, metal boride, metal
silicide, metal oxide, or combinations thereof which is
subsequently deposited (for example by chemical vapor infiltration)
onto the previously coated refractory fibers within the fibrous
preform thereby encapsulating the fibers and consolidating the
preform into the resultant densified composite. The reinforcing
fibers of the fibrous preform thus become imbedded within and
supported by the surrounding matrix. Examples include pyrolytic
carbon, silicon carbide, boron carbide, silicon nitride, boron
nitride, silicon boride, etc.; either as a single-phase or as a
phase of mixed composition.
[0025] Nanolayered ceramic matrix is defined as any ceramic matrix
with a compositionally modulated microstructure consisting of a
plurality of very thin periodically alternating layers (lamellae)
composed of two or more material constituents which is subsequently
deposited (for example by chemical vapor infiltration) onto the
previously coated refractory fibers within the fibrous preform
thereby encapsulating the fibers and consolidating the preform into
the resultant densified composite.
[0026] Chemical precursors are defined as any intermediate
chemical(s) or presursory material(s) used in a process to form a
resulting solid material upon thermochemical reaction. For
processing methods of chemical vapor infiltration (CVI) and
chemical vapor deposition (CVD), the chemical precursors used to
deposit ceramic materials are usually gases or vaporized
liquids.
BRIEF DESCRIPTION OF THE DRAWINGS
[0027] FIG. 1 shows a high magnification backscattered electron
image of an advanced multilayered SiC fiber coating produced by
chemical vapor infiltration.
[0028] FIG. 2 shows a high magnification scanning electron
micrograph of a nanolayered silicon carbide material produced by
chemical vapor deposition.
[0029] FIG. 3 shows a high magnification scanning electron
micrograph of a fiber-reinforced nanolayered silicon carbide matrix
composite material microstructure.
[0030] FIG. 4 shows experimental uniaxial tensile stress-strain
curves acquired from a baseline non-layered Hi-Nicalon/CVI SiC
composite material system with a PyC fiber coating.
[0031] FIG. 5 shows experimental uniaxial tensile stress-strain
curves acquired from a baseline non-layered fi-Nicalon/CVI SiC
composite material system with a BN fiber coating.
[0032] FIG. 6 shows the corresponding tensile fracture surface for
a baseline non-layered Hi-Nicalon/CVI SiC composite material system
with a PyC fiber coating.
[0033] FIG. 7 shows the corresponding tensile fracture surface for
a baseline non-layered Hi-Nicalon/CVI SiC composite material system
with a BN fiber coating.
[0034] FIG. 8 shows experimental uniaxial tensile stress-strain
curves acquired from a Hi-Nicalon/CVI SiC composite material system
having a nanolayered matrix and a PyC fiber coating.
[0035] FIG. 9 shows experimental uniaxial tensile stress-strain
curves acquired from a Hi-Nicalon/CVI SiC composite material system
having a nanolayered matrix and a BN fiber coating.
[0036] FIG. 10 shows the corresponding tensile fracture surface for
a Hi-Nicalon/CVI SiC composite material system having nanolayered
matrix and a PyC fiber coating.
[0037] FIG. 11 show the corresponding tensile fracture surface for
a Hi-Nicalon/CVI SiC composite material system having nanolayered
matrix and a BN fiber coating.
DETAILED DESCRIPTION OF THE INVENTION
[0038] In the present invention, a fiber-reinforced ceramic matrix
composite material with increased matrix cracking strength over
currently available materials and a method for producing the same
is described in detail with reference to the attached Figures.
According to the present invention, a high-temperature
fiber-reinforced ceramic matrix composite material exhibiting
increased resistance to the onset of matrix cracking is produced by
engineering the microstructure of the ceramic matrix
constituent.
[0039] One embodiment of the present invention is a ceramic matrix
composite material comprising a fibrous preform of refractory
fibers, a fiber coating and a matrix with a nanolayered
microstructure. The microstructure of the nanolayered ceramic
matrix material comprises a plurality of primary phase layers of a
first material and a plurality of secondary phase layers of a
second material, wherein the secondary phase layers are interposed
between the primary phase layers. Further, the nanolayered ceramic
matrix encapsulates the coated refractory fibers and consolidates
the fibrous preform into a densified composite.
[0040] Another embodiment of the present invention is a method for
manufacturing a fiber-reinforced nanolayered ceramic matrix
composite material comprising providing a fibrous reinforcing
preform composed of refractory fibers, depositing (for example by
chemical vapor infiltration) a fiber coating material onto the
refractory fibers of the fibrous preform, and depositing (for
example by chemical vapor infiltration) a nanolayered ceramic
matrix material comprising a plurality of primary phase layers of a
first material and a plurality of secondary phase layers of a
second material, wherein the secondary phase layers are interposed
between the primary phase layers. Further, the nanolayered ceramic
matrix material is deposited such that it encapsulates the coated
refractory fibers of the fibrous preform and consolidates the
preform into a densified composite.
[0041] According to a manufacturing process of the current
invention, a fibrous reinforcing preform is produced by weaving,
braiding, filament winding, fiber placement, felting, needling, or
other textile fabrication process using refractory fibers such as
carbon fibers, silicon carbide fibers, silicon nitride fibers,
aluminum oxide fibers, etc.
[0042] The fibrous preform is then generally fixtured in a suitable
holding tool. Tooling serves to provide structural support to the
otherwise flexible fiber preform thereby controlling and
maintaining the desired geometry, dimensional tolerance(s) and/or
fiber volume fraction of the preform prior to and following matrix
consolidation. Suitable tooling materials should be selected on the
basis of several considerations: (1) thermochemical stability in
the fiber coating and matrix densification processes; (2)
thermochemical and thermomechanical compatibility with the preform
reinforcing fiber; (3) thermophysical stability for maintaining
desired component geometry during and following materials
processing; and (4) cost. Examples include graphite, molybdenum and
stainless steel.
[0043] Following fabrication and fixturing of the fibrous preform,
a fiber coating or fiber coating system is deposited (for example
by chemical vapor infiltration) onto the preform. A method of
depositing a suitable a multilayered fiber coating system onto a
refractory fiber preform is disclosed in U.S. Pat. No. 5,455,106,
the disclosure of which has been previously incorporated herein. If
the reinforcing fibers have been coated prior to fibrous preforming
(e.g., fiber-level coating), this processing step may not be
necessary. The deposited fiber coating serves to control the
desired fiber/matrix interfacial compliance and bonding
characteristics in the resultant ceramic matrix composite. Examples
of suitable fiber coatings include pyrolytic carbon, silicon
carbide, silicon nitride, boron carbide, boron nitride, etc.,
either as a single-layer phase, multilayered phase or as a phase of
mixed composition.
[0044] In a preferred method of the present invention the
previously coated fibrous preform is consolidated with a
nanolayered ceramic matrix. The microstructure of the nanolayered
matrix consists of numerous very thin, periodically alternating
layers (i.e., lamellae) of two or more material species, or phases.
The nanolayered matrix constituent consists fundamentally of a
primary phase material and a secondary phase material. Suitable
examples for the primary phase material include silicon carbide,
silicon nitride, boron carbide, etc., either as a single-phase, or
as a phase of mixed composition. Suitable examples for the
secondary phase material include pyrolytic carbon, silicon, silicon
carbide, silicon nitride, boron carbide, boron nitride, etc.,
either as a single-layer phase, or as a phase of mixed
composition.
[0045] In accordance with the teachings of the present invention,
the nanolayered matrix is produced by depositing (for example by
chemical vapor infiltration) a plurality of layers of a primary
phase material with a thickness of on the order of a few tens of
nanometers, followed by depositing (for example by chemical vapor
infiltration) a plurality of layers of a secondary phase material
with a thickness of about an order of magnitude less than that of
the primary phase layers. The nanolayered matrix is thereby
deposited such that the secondary phase material layers are
interposed between the primary phase material layers. In a
preferred embodiment, the thicknesses of the primary phase layers
comprising the nanolayered ceramic matrix are between 5 and 500 nm,
and the thicknesses of the secondary phase layers comprising the
nanolayered ceramic matrix are between 1 and 100 nm. Thus, by
interrupting grain growth and limiting the grain size to the
thicknesses of the primary and secondary material layers, the
present invention dramatically increases the resistance to the
onset of matrix cracking in the resultant composite over that
exhibited by an unlayered or microlayered matrix constituent.
EXAMPLE
[0046] A preferred method and resultant ceramic composite material
of the present invention will now be described by way of example
using a Hi-Nicalon silicon carbide (SiC) fiber-reinforced SiC
matrix composite (SiC/SiC) material system produced by methods of
chemical vapor infiltration (CVI). The SiC/SiC composite material
system is selected for its desired nonlinear tensile stress-strain
behavior and well-defined matrix cracking strength, or proportional
limit. The CVI-based processing technique for fiber coating and
matrix densification is selected for its ability to produce very
high purity materials with outstanding control and uniformity of
deposited material stoichiometry, morphology and thickness.
[0047] A first step in producing ceramic matrix composite materials
according to the present invention is the fabrication of a suitable
fiber reinforcing preform. For this example, fibrous preforms are
fabricated by stacking 8 plies of plain weave Hi-Nicalon SiC woven
fabric in a balanced and symmetric cross-ply
(0.degree./90.degree.).sub.2s laminate orientation. Each dry (i.e.,
no resinous binders) laminated preform is then fixtured and
flat-wise compacted in a graphite holding tool to maintain a fiber
volume fraction of .about.40% by controlling the preform thickness
prior to subsequent composite processing.
[0048] An initial processing step in a preferred method of the
present invention is the application of a suitable fiber coating or
fiber coating system by methods of CVI onto the laminated fibrous
preforms prior to their consolidation with the matrix material
deposited by CVI or other suitable matrix formation process. As
previously discussed, a fiber coating is required to impart the
necessary fiber/matrix interfacial mechanical characteristics
(e.g., low interfacial shear strength) to promote high strength and
toughness in the resulting composite. For this example, two (2)
suitably different fiber-coating materials, namely pyrolytic carbon
(PyC) and pyrolytic boron nitride (BN), will be used independently
for demonstrating the benefits of the nanolayered SiC matrix
system. It has been found that PyC and BN coating thicknesses of
.about.0.4 .mu.m and .about.0.6 .mu.m, respectively, result in
near-optimum mechanical performance for CVI-based SiC/SiC
composites. Thus for this example, PyC and BN fiber coatings will
be applied to the laminated fabric preforms by CVI with thicknesses
of .about.0.4 .mu.m and .about.0.6 .mu.m, respectively.
[0049] In the current example, the CVI-derived PyC fiber coating is
produced in a high-temperature, low-pressure chemical vapor
deposition (CVD) reactor by the thermal decomposition of a
hydrocarbon-containing gas in the presence of hydrogen according to
the following chemical reaction:
1/nC.sub.nH.sub.m+.alpha.H.sub.2.fwdarw.C+(m/2n+.alpha.)H.sub.2,
[0050] where C.sub.nH.sub.m is the gaseous hydrocarbon reactant
(i.e., methane, propane, propylene etc.) and .alpha. is defined as
the molar ratio of H.sub.2 to C.sub.nH.sub.m. At deposition
temperatures between 1000-1400.degree. C., the deposit is typically
smooth laminar PyC, which has a hexagonal structure and, depending
on the deposition temperature, has a density ranging between
1.8-2.0 g/cm.sup.3.
[0051] Further, the CVI-derived BN fiber coating is produced in a
high-temperature, low-pressure CVD reactor by the hydrogen
reduction of gaseous boron trichloride in the presence of ammonia
according to the following chemical reaction:
BCl.sub.3+.eta.NH.sub.3.fwdarw.BN+(.eta.-1)NH.sub.3+3HCl,
[0052] where .eta. is defined as the molar ratio of NH.sub.3 to
BCl.sub.3. At deposition temperatures between 800-1200.degree. C.,
the deposit is typically amorphous and, depending on the deposition
temperature, has a density ranging between 1.8-2.0 g/cm.sup.3. At
temperatures above 1200.degree. C., an increasingly crystalline
deposit is obtained, until at around 1400.degree. C. a crystalline
BN is formed which has a hexagonal structure and a theoretical
density of 2.2 g/cm.sup.3.
[0053] Following the application of the PyC and BN fiber coatings,
the fixtured fiber preforms are prepared for matrix consolidation
and densification processing. Baseline (i.e., unlayered matrix)
SiC/SiC composites incorporating each the PyC and BN fiber coatings
are also fabricated as experimental control in order to quantify
the matrix cracking test results obtained on the respective
nanolayered matrix composites of the present invention. The
baseline unlayered SiC matrix is produced in a CVD reactor by the
thermal decomposition of vaporized methyltrichlorosilane (MTS)
using hydrogen as a carrier gas at elevated temperature and reduced
pressure according to the following chemical reaction:
CH.sub.3SiCl.sub.3+.alpha.H.sub.2.fwdarw.SiC+3HCl+.alpha.H.sub.2,
[0054] where .alpha. is defined as the molar ratio of H.sub.2 to
CH.sub.3SiCl.sub.3. At deposition temperatures between
900-1300.degree. C., the deposit is typically crystalline beta-SiC,
which has a cubic structure and a theoretical density of 3.21
g/cm.sup.3. The laminated preforms should remain fixtured until an
initial level of CVI SiC matrix is deposited to adequately rigidize
the preforms. Upon rigidization, the external graphite tooling can
be removed to enable matrix consolidation to continue on the
freestanding preforms. CVI SiC densification processing is thus
continued until a zero residual open porosity level is obtained in
the resultant composite material.
[0055] In a preferred method, the nanolayered SiC matrix composite
utilizes SiC as the primary or major matrix constituent phase and
PyC as the secondary or minor matrix constituent phase. PyC is
selected as the minor nanolayering constituent because of its known
ability to effectively interrupt the epitaxial growth of vapor
deposited SiC, while combining the benefits of extreme elastic
modulus mismatch with SiC for reasons previously discussed. A
suitable thickness of the major SiC constituent is less than 100
nm. A suitable selected thickness of the minor PyC constituent is
about 5 nm. The layer thickness of the primary phase material will
be substantially thicker than that of the secondary phase,
generally to an order of magnitude or greater. The CVI SiC primary
layers are deposited from vaporized methyltrichlorosilane (MTS) as
previously described above, and the CVI PyC secondary, or
interrupter layers, are produced by the thermal decomposition of
methane, also as previously described. Both the CVI process
temperature and pressure can generally be maintained constant
throughout the deposition of the nanolayered matrix in order to
provide overall simplicity and time-related economics to the
process. In a preferred method of the present invention, primary
and secondary constituent layer thicknesses can be precisely
controlled by microprocessor-based cyclic throttling of the
respective chemical precursors at prescribed time intervals. As
discussed above, the appropriate time intervals will be dependant
on the respective deposition rates of the primary and secondary
phase matrix materials, and the nanolayer thicknesses desired and
are readily determinable by one skilled in the art.
[0056] The laminated fiber preforms should remain fixtured until a
level of CVI nanolayered SiC matrix is deposited to adequately
rigidize the preforms. The external graphite tooling can then be
removed upon rigidization to enable matrix consolidation to
continue on the freestanding preforms. With the tooling fixtures
removed, cyclic application of the primary and secondary matrix
constituent materials is continued on the freestanding preforms
until a zero residual open porosity level is obtained. FIG. 3 shows
a typical scanning electron micrograph of a nanolayered SiC matrix
composite microstructure.
[0057] Mechanical test specimens are then machined from each of the
two (2) densified baseline and two (2) nanolayered SiC/SiC ceramic
composite plates produced for experimental evaluation. Nine (9)
replicate tensile test specimens from each of the four (4) SiC/SiC
composite variants (e.g., two (2) fiber coatings and two (2)
matrices) are prepared and instrumented with resistive foil gages
for longitudinal strain measurement and adhesive-bonded end tabs
for load introduction. In order to quantify the beneficial results
of nanolayering the matrix microstructure on the matrix cracking
strength, axial monotonic loading is performed at ambient
conditions under an enforced displacement rate of 1.27 mm/min as
per ASTM testing standards. An acoustic emission technique can be
utilized during tensile testing to better establish and
characterize the initial on-set of matrix cracking and multiple
matrix fracture process signatures for the composite specimens
tested. Data acquisition for this example demonstration included
continuous tensile stress-strain curves and continuous
strain-dependent acoustic matrix cracking signatures.
[0058] FIGS. 4 and 5 show experimental uniaxial tensile
stress-strain curves (9 replicates each) acquired from the baseline
Hi-Nicalon/non-layered SiC matrix composites with PyC and BN fiber
coatings, respectively. As can be seen from these graphical
figures, the overall stress-strain characteristics appear to be
insensitive to the particular fiber coating system as both exhibit
"classical" brittle-matrix composite stress-strain behavior. That
is, both graphs exhibit a three (3) regime characteristic
including: (1) an initial linear-elastic region, followed by (2) an
erratic "knee" during early multiple matrix fracture, which then
(3) terminates to a continuous nonlinear trajectory during the
transition between latent matrix cracking, crack saturation and
fiber debonding. For both conventional fiber-coating systems,
deviation from elastic linearity, designating the onset of matrix
cracking, consistently occurred at stress levels of about 80 MPa.
The region immediately following the onset of matrix cracking is of
particular interest because of the erratic behavior in the
stress-strain characteristics. Each apparent stress "jumping" event
is due to the formation and dynamic propagation (i.e., energy
release) of relatively large matrix macrocracks which travel
rapidly across the test specimen cross-section.
[0059] FIGS. 6 and 7 show the corresponding tensile fracture
surfaces for the baseline PyC and BN fiber-coated SiC matrix
composites, respectively. These high-magnification images depict a
classical fracture morphology typical of successfully manufactured
continuous fiber-reinforced ceramic composites as exhibited by
limited matrix crack tortuosity and liberal fiber pullout.
[0060] In contrast, FIGS. 8 and 9 show the uniaxial tensile
stress-strain curves (9 replicates each) acquired from
Hi-Nicalon/nanolayered SiC matrix composites with PyC and BN fiber
coatings, respectively, made in accordance with the teachings of
the present invention. Deviation from linearity (e.g., onset of
matrix cracking) in these cases occur at stress levels consistently
above 100 Mpa. Of particular interest is that following the onset
of matrix cracking, these composites exhibit a very smooth and
rapid nonlinear stress-strain transition to a secondary linear
region. This is believed to be due to the tortuosity of matrix
crack propagation in the nanolayered matrix following the onset of
cracking. That is, throughout the multiple matrix fracturing
process, cracks which initiate at free surfaces (e.g., external and
internal voids) must propagate towards the fiber through individual
matrix layers via the shear lag mechanism. Because of combined
compliance mismatch and low interfacial bond strength, these layers
impede rapid co-planer crack propagation through extensive
interlayer microcracking, debonding and displacement slip.
Propagating internal matrix microcracks are thus effectively
diverted and/or blunted prior to their approach towards the
reinforcing fiber, serving to increase the apparent toughness of
the matrix constituent. This mechanism effectively dampens
energetic macrocrack propagation and resulting erratic
stress-strain behavior typical in conventionally processed SiC
matrix composites wherein matrix cracks are usually deflected at
the fiber/matrix interphase region. The smooth continuous
transition during early multiple matrix cracking shown in FIGS. 8
and 9 is likely due to the occurrence of a greater frequency of
lower fracture energy release events through the thin matrix
layers. These lower energy events aid in dampening the discrete
jumps in stress by producing a greater density of fine matrix
microcracks as opposed to fewer, more energetic macrocracks.
Following the region of intense multiple matrix fracture, the
stress-strain trajectories quickly recover, resulting in
quasi-bilinear elastic behavior.
[0061] FIGS. 10 and 11 show the corresponding tensile fracture
surfaces for PyC and BN fiber coated nanolayered SiC matrix
composites, respectively, produced in accordance with the teachings
of the present invention described above. These images clearly show
the unique and complex failure mechanisms afforded by the matrix
nanolayering process of the present invention. The "stair-stepping"
fracture morphology with extensive layer debonding and crack
branching within the matrix layers can clearly be seen. Examination
of the high-magnification scanning electron micrographs also shows
that matrix microcrack propagation appears to "jump" between four
and six discrete primary matrix layers. This is likely due to the
magnitude of the strain energy that must be dissipated (i.e.,
released) during the propagation of internal fractures upon the
formation of new surfaces. With similar reasoning, this interlayer
crack propagation phenomenon may be a function of the applied
strain rate wherein higher rates may result in greater discrete
jumps. Unlike that of the baseline materials, these composites
consistently exhibited highly irregular "cup-like" fracture
surfaces. This apparent matrix crack tortuosity is indicative of a
high level of matrix material toughness.
[0062] Table 1 summarizes the comparative tensile test results for
the baseline non-layered Hi-Nicalon/CVI SiC composites as compared
with the Hi-Nicalon/CVI SiC composites having a nanolayered matrix
microstructure produced in accordance with the teachings of the
present invention. As indicated by these results, the nanolayered
SiC matrix composite of the present invention exhibits a consistent
matrix cracking strength increase of over 30% for both PyC and BN
fiber coatings as compared to that of the conventional un-layered
SiC matrix composite baseline.
[0063] A nanolayered matrix composite material and a method for
producing the same has been herein shown and described. While the
preferred embodiments of the devices and methods have been
described, they are merely illustrative of the principles of the
invention. Other embodiments and configurations may be devised
without departing from the spirit of the inventions and the scope
of the appended claims.
1 TABLE 1 Matrix System Baseline SiC** Nanolayered SiC** Fiber
Coating System Fiber Coating System Material Properties* Units PyC
BN PyC BN Fabric Areal Weight g/m.sup.2 280 280 280 280 Fiber
Density g/cm.sup.3 2.74 2.74 2.74 2.74 Fiber Volume Fraction % 29.0
(1.4) 31.5 (1.5) 29.8 (1.0) 33.0 (1.1) Fiber Coating Thickness
.mu.m .about.0.4 .about.0.6 .about.0.4 .about.0.6 Fiber Coating
Fraction % 3.2 (0.2) 5.2 (0.3) 3.3 (0.1) 5.5 (0.2) Theoretical
Density g/cm.sup.3 3.06 (0.01) 3.06 (0.01) 3.06 (0.01) 3.05 (0.01)
Measured Bulk Density g/cm.sup.3 2.71 (0.04) 2.68 (0.02) 2.67
(0.05) 2.54 (0.04) Residual Porosity % 11.7 (1.3) 12.4 (0.8) 12.8
(1.7) 12.5 (1.3) Matrix Layer Thickness nm -- -- .about.100
.about.100 Matrix Cracking Strength MPa 81.2 (10.3) 80.0 (11.2)
106.1 (8.9) 104.9 (11.9) Cracking Strength Increase % -- -- 30.7
30.1 Ultimate Tensile Strength MPa 318.6 (26.2) 332.3 (17.2) 298.6
(20.5) 307.0 (14.1) Tensile Failure Strain % 0.68 (0.09) 0.79
(0.09) 0.57 (0.06) 0.80 (0.03) Tensile Elastic Modulus GPa 322.8
(85.3) 271.1 (19.3) 277.0 (61.8) 293.5 (67.4) Number of Replicate
Tests -- 9 9 9 9 *Average Value (Standard Deviation) **Plain-Weave
Fabric, 8-Ply (0/90) Laminate
* * * * *