U.S. patent application number 09/790118 was filed with the patent office on 2002-05-16 for composite materials and methods for making the same.
Invention is credited to Fareed, Ali Syed, Garnier, John Edward, Kennedy, Christopher Robin, Schiroky, Gerhard Hans, Sonuparlak, Birol.
Application Number | 20020058107 09/790118 |
Document ID | / |
Family ID | 25501192 |
Filed Date | 2002-05-16 |
United States Patent
Application |
20020058107 |
Kind Code |
A1 |
Fareed, Ali Syed ; et
al. |
May 16, 2002 |
Composite materials and methods for making the same
Abstract
The present invention generally relates to mechanisms for
preventing undesirable oxidation (i.e., oxidation protection
mechanisms) in composite bodies. The oxidation protection
mechanisms include getterer materials which are added to the
composite body which gather or scavenge undesirable oxidants which
may enter the composite body. The getterer materials may be placed
into at least a portion of the composite body such that any
undesirable oxidant approaching, for example, a fiber
reinforcement, would be scavenged by (e.g., reacted with) the
getterer. The getterer material(s) may form at least one compound
which acts as a passivation layer, and/or is able to move by bulk
transport (e.g., by viscous flow as a glassy material) to a crack,
and sealing the crack, thereby further enhancing the oxidation
protection of the composite body. One or more ceramic filler
materials which serve as reinforcements may have a plurality of
super-imposed coatings thereon, at least one of which coatings may
function as or contain an oxidation protection mechanism.
Specifically, a coating comprising boron nitride which has been
engineered or modified to contain some silicon exhibits improved
corrosion resistance, specifically to oxygen and moisture. The
coated materials may be useful as reinforcing materials in high
performance composites to provide improved mechanical properties
such as fracture toughness. The present invention also relates to
improved composites which incorporate these materials, and to their
methods of manufacture.
Inventors: |
Fareed, Ali Syed; (Newark,
DE) ; Garnier, John Edward; (Newark, DE) ;
Schiroky, Gerhard Hans; (Newark, DE) ; Kennedy,
Christopher Robin; (Newark, DE) ; Sonuparlak,
Birol; (Longmont, CO) |
Correspondence
Address: |
Larry J. Palguta
Honeywell Law Department
3520 Westmoor Street
South Bend
IN
46628
US
|
Family ID: |
25501192 |
Appl. No.: |
09/790118 |
Filed: |
February 21, 2001 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
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09790118 |
Feb 21, 2001 |
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08958685 |
Oct 27, 1997 |
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Current U.S.
Class: |
427/255.39 |
Current CPC
Class: |
Y10T 428/15 20150115;
C04B 2235/6567 20130101; C04B 2235/405 20130101; C04B 2235/5436
20130101; Y10T 428/12035 20150115; C04B 35/62894 20130101; Y10T
428/2991 20150115; C04B 35/62884 20130101; C04B 35/117 20130101;
C04B 2235/5244 20130101; C04B 2237/38 20130101; C04B 35/62897
20130101; Y10T 428/12056 20150115; C04B 2235/407 20130101; C04B
2235/526 20130101; C04B 2235/614 20130101; C04B 2235/96 20130101;
C04B 2237/586 20130101; Y10T 428/12458 20150115; C04B 2235/5256
20130101; C04B 2237/343 20130101; C04B 35/62852 20130101; C04B
2235/401 20130101; C04B 35/62871 20130101; C04B 35/80 20130101;
C04B 2235/6586 20130101; C04B 2235/5264 20130101; C04B 35/62863
20130101; C04B 35/652 20130101; C04B 41/52 20130101; C04B 2235/6587
20130101; C04B 2235/6562 20130101; C04B 2235/40 20130101; C04B
35/581 20130101; C04B 2235/3454 20130101; C04B 2235/5268 20130101;
B32B 18/00 20130101; C04B 2235/662 20130101; C04B 35/62868
20130101; C04B 35/653 20130101; C04B 2235/46 20130101; C04B
2235/6028 20130101; C04B 2235/428 20130101; C04B 35/62844 20130101;
C04B 41/009 20130101; C04B 2235/3279 20130101; C04B 2235/402
20130101; C04B 41/52 20130101; C04B 41/4531 20130101; C04B 41/5064
20130101; C04B 41/52 20130101; C04B 41/4531 20130101; C04B 41/5059
20130101; C04B 41/52 20130101; C04B 41/4531 20130101; C04B 41/5066
20130101; C04B 41/522 20130101; C04B 41/52 20130101; C04B 41/4531
20130101; C04B 41/5031 20130101; C04B 41/522 20130101; C04B 41/009
20130101; C04B 35/62844 20130101; C04B 41/009 20130101; C04B 35/80
20130101; C04B 41/009 20130101; C04B 14/4693 20130101 |
Class at
Publication: |
427/255.39 |
International
Class: |
C23C 016/08; C23C
016/00 |
Goverment Interests
[0002] This invention was made with Government support under
Contract No. DE-FC02-92CE40994 awarded by the Department of Energy.
The Government has certain rights to this invention.
Claims
What is claimed is:
1. A reinforcement for a composite material, comprising: (a) a
permeable mass or preform comprising a plurality of bodies of at
least one filler material; and (b) a layer covering at least the
majority of surface area presented by said bodies of filler
material, said layer comprising boron, silicon and nitrogen.
2. A composite material, comprising: (a) a permeable mass or
preform comprising a plurality of bodies of at least one filler
material; (b) a matrix embedding said permeable mass or preform;
and (c) at least one coating disposed between said bodies of at
least one filler material and said matrix, at least one of said at
least onecoating comprising boron, silicon and nitrogen.
3. A method for making a reinforcement for a composite material,
comprising: providing at least one filler material; communicating a
local atmosphere comprising a halogenated boron source, a
halogenated silicon source, and ammonia, to said at least one
filler material; and heating said at least one filler material and
said local atmosphere to a temperature in the range of about
700.degree. C. to 1200.degree. C., thereby depositing on said at
least one filler material a coating comprising boron, silicon and
nitrogen.
4. A coated fiber comprising a substrate fiber or filament and a
plurality of coatings disposed coextensively with said substrate,
said plurality comprising at least one protective coating and at
least one debond coating disposed between said substrate and said
at least one protective coating, and further at least one of said
at least one debond coating comprises boron, silicon and
nitrogen.
5. A coated filler material for use as a reinforcement in a
composite body, said coated filler comprising: a substrate body;
and a layer no greater than about 0.5 micron in thickness covering
at least a majority of surface presented by said substrate body,
said layer comprising boron, silicon, nitrogen and oxygen.
6. The reinforcement of claim 1, wherein said layer comprises boron
nitride.
7. The reinforcement of claim 1, wherein at least one of said
plurality of bodies of at least one filler material comprises a
filament or fiber, said layer being coextensive with a longitudinal
axis of said filament or fiber.
8. The coated fiber of claim 4, wherein said at least one
protective coating comprises at least one material selected from
the group consisting of silicon carbide, silicon nitride, and
aluminum oxide.
9. The coated fiber of claim 4, wherein said debond coating
comprises boron nitride.
10. The coated fiber of claim 9, wherein said debond coating is at
least partially amorphous.
11. The coated fiber of claim 4, wherein said debond coating
exhibis limited crystallinity.
12. The coated fiber of claim 4, wherein said debond coating
comprises a plurality of regions or domains each about 5 to 20
nanometers in size, wherein the debond coating material within a
region or domain exhibits a lamellar structure.
13. The coated fiber of claim 30, wherein a lamellar crystal
structure within a given region or domain essentially is randomized
in orientation with respect to a lamellar crystal structure in a
different region or domain.
14. The composite material of claim 2, further comprising at least
one oxide glass network-former.
15. The composite material of claim 2, wherein said matrix
comprises a material selected from the group consisting of silicon,
silicon carbide and aluminum oxide.
16. The composite material of claim 2, further comprising at least
one oxygen getterer.
17. The composite material of claim 2, further comprising at least
two zonal junctions, at least one of said zonal junctions being
weak relative to the remaining zonal junction(s) to permit
debonding and pull-out of said at least one filler material with
respect to said matrix upon application of stress sufficient to
cause fracture of said composite material.
18. The composite material of claim 17, wherein said debonding
occurs at an interface between a coating and (a) said matrix, (b)
said filler material or (c) another coating, and not within a
coating.
19. The composite material of claim 2, wherein said matrix is
produced by a method selected from the group consisting of directed
metal oxidation and melt infiltration.
20. The method of claim 3, wherein said halogenated boron source
comprises boron trichloride, and said halogenated silicon source
comprises silicon tetrachloride.
21. The method of claim 3, wherein said local atmosphere is
communicated to said at least one filler material at a pressure of
about 1 Torr to about 10 Torr, and said temperature is in the range
of about 700.degree. C. to 800.degree. C.
22. The method of claim 3, wherein an atomic ratio of said
halogenated silicon source to said halogenated boron source in said
local atmosphere ranges from about 0.25 to about 7.7.
23. The coated filler material of claim 5, wherein at least about
0.5 atom percent of said layer comprises said silicon.
24. The coated filler material of claim 5, wherein said layer
further comprises carbon.
25. The coated filler material of claim 5, wherein said substrate
body comprises a fiber comprising silicon carbide.
26. The coated filler material of claim 5, wherein about 1 percent
to about 3 percent of atoms making up said layer comprise silicon
atoms.
Description
CROSS REFERENCE TO RELATED APPLICATIONS
[0001] The present patent disclosure is a continuation-in-part of
U.S. patent application Ser. No. 08/472.613, filed on Jun. 7, 1995,
now U.S. Pat. No. 5,682,594, which issued on Oct. 28, 1997, in the
names of Christopher R. Kennedy et al., and entitled "Composite
Materials and Methods For Making The Same".
TECHNICAL FIELD
[0003] The present invention generally relates to filler materials
which are adapted for use as the reinforcement phases(s) in
composite bodies. Coated ceramic filler materials comprised of
ceramic particles, fibers, whiskers, etc. having at least two
substantially continuous coatings thereon are provided. The
coatings are selected so that the interfacial shear strength
between the ceramic filler material and the first coating, between
coatings, or between the outer coating and the surrounding matrix
material, are not equal so as to permit debonding and pull-out when
fracture occurs. The resultant, multicoated ceramic filler
materials may be employed to provide composites, especially ceramic
matrix composites with increased fracture toughness.
[0004] The ceramic filler materials are designed to be particularly
compatible with ceramic matrices formed by directed oxidation of
precursor metals, but such ceramic filler materials are also
adaptable for use in many other composite material systems.
[0005] The present invention also relates to techniques for
increasing the corrosion resistance of composite materials,
particularly of ceramic fiber reinforced composites exposed to
oxygen and water vapor at elevated temperatures. One approach to
inhibiting corrosion in a ceramic matrix composite body is to
reduce the number and/or size of microcracks in the body, thereby
reducing access of corrodants to the interior of the body. Another
broad approach is to provide chemical additives to the body which
are capable of gettering a corrodant or interfering with its
corrosion mechanisms.
BACKGROUND ART
[0006] A ceramic composite is a heterogeneous material or article
comprising a ceramic matrix and filler such as ceramic particles,
fibers or whiskers, which are intimately combined to achieve
desired properties. These composites are produced by such
conventional methods as hot pressing, cold pressing and firing, hot
isostatic pressing, and the like. However, these composites
typically do not exhibit a sufficiently high fracture toughness to
allow for use in very high stress environments such as those
encountered by gas turbine engine blades.
[0007] A novel and useful method for producing self-supporting
ceramic composites by the directed oxidation of a molten precursor
metal is disclosed in Commonly Owned U.S. Pat. No. 4,851,375, which
issued on Jul. 25, 1989, described below in greater detail.
However, the processing environment is relatively severe, and there
is a need, therefore, to protect certain fillers from the strong
oxidation environment. Also, certain fillers may be reduced at
least partially by molten metal, and therefore, it may be desirable
to protect the filler from this local reducing environment.
Further, the protective means should be conducive to the metal
oxidation process, yet not degrade the properties of the resulting
composite, and even more desirably provide enhancement to the
properties. Still further, in some instances it may be desirable
for the means or mechanisms for protecting the filler during matrix
or composite formation to also protect the fillers against
undesirable attack of oxidants diffusing through the matrix during
actual service of the composite.
[0008] It is known in the art that certain types of ceramic fillers
serve as reinforcing materials for ceramic composites, and the
selection or choice of fillers can influence the mechanical
properties of the composite. For example, the fracture toughness of
the composite can be increased by incorporating certain high
strength filler materials, such as fibers or whiskers, into the
ceramic matrix. When a fracture initiates in the matrix, the filler
at least partially debonds from the matrix and spans the fracture,
thereby resisting or impeding the progress of the fracture through
the matrix. Upon the application of additional stress, the fracture
propagates through the matrix, and the filler begins to fracture in
a plane different from that of the matrix, pulling out of the
matrix and absorbing energy in the process. Pull-out is believed to
increase certain mechanical properties such as work-of-fracture by
releasing the stored elastic strain energy in a controlled manner
through friction generated between the material and the surrounding
matrix.
[0009] Debonding and pull-out have been achieved in the prior art
by applying a suitable coating to the ceramic filler material. The
coating is selected so as to have a lower bonding strength with the
surrounding matrix than the filler, per se, would have with the
matrix. For example, a boron nitride coating on silicon carbide
fibers has been found to be useful to enhance pull-out of the
fibers. Representative boron nitride coatings on fibers are
disclosed in U.S. Pat. No. 4,642,271, which issued on Feb. 10,
1987, in the name of Roy W. Rice, and are further disclosed in U.S.
Pat. No. 5,026,604, which issued on Jun. 25, 1991, in the name of
Jacques Thebault. However, the use of boron nitride coated fibers
in composites may present significant processing disadvantages. For
example, the production of ceramic matrix composites containing
boron nitride coated materials requires the use of reducing
atmospheres since a thin layer of boron nitride readily oxidizes
(e.g., converts to boron oxide in an oxygen-containing atmosphere)
at temperatures above 800-900.degree. C. A reducing atmosphere,
however, may often times not be compatible with the directed
oxidation of molten parent metal for fabricating ceramic
composites. Further, in the directed oxidation process the coating
desirably is compatible with the molten metal in that the molten
metal wets the coated filler under the process conditions, for
otherwise the oxidation process and matrix growth may be impeded by
the filler.
[0010] Another drawback of boron nitride is that, upon oxidation,
the boria reaction product can dissolve or further react with water
to form boric acid, which can be a vapor under the local oxidizing
conditions. Thus, the boria is not a passive layer, but can be
continually removed through volatilization. U.S. Pat. No. 5,593,728
to Moore et al. addresses this shortcoming of boron nitride.
Specifically, by producing a pyrolytic BN coating containing from 2
to 42 wt % silicon, with substantially no free silicon present,
Moore et al. observe greatly reduced rates of oxidative weight
loss. The coating is formed by CVD using reactant vapors of ammonia
and a gaseous source of both boron and silicon. The gases are
flowed into a reaction chamber between a temperature of
1300.degree. C. and 1750.degree. C. and within a pressure range of
0.1 Torr to 1.5 Torr.
[0011] It is not clear, however, whether the modified BN layer of
Moore et al. permits molten parent metal to wet the coating (for
infiltration) and yet resist any adverse reaction therewith.
Further, the modified BN coatings of Moore et al. were deposited
onto single filaments. Due to the high deposition rates resulting
from the deposition conditions, it is unclear whether the Moore et
al. technique could be applied to coat a plurality of fibers, e.g.
a stack of fabrics making up a preform.
[0012] Also, in order to prevent or minimize filler degradation,
certain limits may be imposed on the conventional fabrication
processes, such as using low processing temperatures or short times
at processing temperature. For example, certain fillers may react
with the matrix of the composite above a certain temperature.
Coatings have been utilized to overcome degradation, but as
explained above, the coating can limit the choice of processing
conditions. In addition, the coating should be compatible with the
filler and with the ceramic matrix.
[0013] A need therefore exists to provide coated ceramic filler
materials which are capable of debonding and pull-out from a
surrounding ceramic matrix. A further need exists to provide coated
ceramic filler materials which may be incorporated into the ceramic
matrix at elevated temperatures under oxidizing conditions to
provide composites exhibiting improved mechanical properties such
as increased fracture toughness.
[0014] In order to meet one or more of these needs, the prior art
shows filler materials bearing one or more coatings. Carbon is a
useful reinforcing filler but typically is reactive with the matrix
material. It therefore is well known in the art to provide the
carbon fibers with a protective coating. U.S. Pat. 4,397,901, which
issued on Aug. 9, 1983, in the name of James W. Warren, teaches
first coating carbon fibers with carbon as by chemical vapor
deposition, and then with a reaction-formed coating of a metallic
carbide, oxide, or nitride. Due to a mismatch in thermal expansion
between the fiber and the coating, the fiber is capable of moving
relative to the coating to relieve stress. A duplex coating on
carbon fibers is taught by U.S. Pat. No. 4,405,685, which issued on
Sep. 20, 1983, in the names of Honjo et al. The coating comprises a
first or inner coating of a mixture of carbon and a metal carbide
and then an outer coating of a metal carbide. The outer coatings
prevent degradation of the fiber due to reaction of unprotected
fiber with the matrix material, and the inner coating inhibits the
propagation of cracks initiated in the outer layer. U.S. Pat. No.
3,811,920, which issued on May 21, 1974, in the names of Galasso et
al. relating to metal matrix composites, discloses coated fibers as
a reinforcing filler, such as boron filaments having a silicon
carbide surface layer and an additional outer coating of titanium
carbide. This reference teaches that the additional coating of
titanium carbide improves oxidation resistance as well as provides
a diffusion barrier between the filament and metal matrix.
[0015] However, the prior art fails to teach or suggest filler
materials with a duplex coating for protection from potentially
corrosive environments during manufacture or operation of the
composite body and yet in the composite material permit the filler
to debond and pull-out from the surrounding matrix. Moreover, the
prior art does not recognize certain other oxidation protection
mechanisms which can be employed jointly. Specifically, the prior
art fails to appreciate certain important aspects of utilizing
getterer materials which function to scavenge undesirable oxidants,
and optionally after such scavenging has occurred, forming
desirable compounds or materials (e.g., one or more glassy
compounds) which assist in protecting the reinforcement materials
from undesirable oxidation.
Description of Commonly Owned U.S. Patents and Patent
Applications
[0016] The filler materials utilized in this invention may be
protected by a number of different mechanisms in a number of
different composite bodies. Filler materials containing a coating
or plurality of coatings, in accordance with the teachings of this
invention, are particularly applicable or useful in the production
of ceramic composites disclosed and claimed in Commonly Owned U.S.
Pat. No. 4,851,375, entitled "Methods of Making Composite Ceramic
Articles Having Embedded Filler," which issued on Jul. 25, 1989,
from U.S. patent application Ser. No. 819,397, filed Jan. 17, 1986,
which is a continuation-in-part of Ser. No. 697,876, filed Feb. 4,
1985 (now abandoned), both in the names of Marc S. Newkirk et al.
and entitled "Composite Ceramic Articles and Methods of Making
Same". This Commonly Owned Patent discloses a novel method for
producing a self-supporting ceramic composite by growing an
oxidation reaction product from a precursor metal or parent metal
into a permeable mass of filler.
[0017] The method of growing a ceramic product by an oxidation
reaction of a parent metal is disclosed generically in Commonly
Owned U.S. Pat. No. 4,713,360, which issued on Dec. 15, 1987, in
the names of Marc S. Newkirk et al. and entitled "Novel Ceramic
Materials and Methods of Making Same"; and in U.S. Pat. No.
4,853,352, which issued on Aug. 1, 1989, in the names of Marc S.
Newkirk et al. and entitled "Methods of Making Self-Supporting
Ceramic Materials".
[0018] Commonly Owned U.S. Pat. No. 4,713,360 discloses a novel
method for producing a self-supporting ceramic body by oxidation of
a parent metal (as defined below) to form an oxidation reaction
product which then comprises the ceramic body. More specifically,
the parent metal is heated to an elevated temperature above its
melting point but below the melting point of the oxidation reaction
product in order to form a body of molten parent metal which reacts
upon contact with a vapor-phase oxidant to form an oxidation
reaction product. The oxidation reaction product, or at least a
portion thereof which is in contact with and extends between the
body of molten parent metal and the oxidant, is maintained at the
elevated temperature, and molten metal is drawn through the
polycrystalline oxidation reaction product and towards the oxidant,
and the sported molten metal forms oxidation reaction product upon
contact with the oxidant. As the process continues, additional
metal is transported through the polycrystalline oxidation reaction
product formation thereby continually "growing" a ceramic structure
of interconnected crystallites. Usually, the resulting ceramic body
will contain therein inclusions of nonoxidized constituents of the
parent metal drawn through the polycrystalline material and
solidified therein as the ceramic body cooled after termination of
the growth process. As explained in these Commonly Owned Patents
and Patent Applications, resultant novel ceramic materials are
produced by the oxidation reaction between a parent metal and a
vapor phase oxidant, i.e., a vaporized or normally gaseous
material, which provides an oxidizing atmosphere. In the case of an
oxide as the oxidation reaction product, oxygen or gas mixtures
containing oxygen (including air) are suitable oxidants, with air
usually being preferred for obvious reasons of economy. However,
oxidation is used in its broad sense in the Commonly Owned Patents
and in this application, and refers to the loss or sharing of
electrons by a metal to an oxidant which may be one or more
elements and/or compounds. Accordingly, elements other than oxygen
may serve as the oxidant. In certain cases, the parent metal may
require the presence of one or more dopants in order to influence
favorably or to facilitate growth of the ceramic body, and the
dopants are provided as alloying constituents of the parent metal.
For example, in the case of aluminum as the parent metal and air as
the oxidant, dopants such as magnesium and silicon, to name but two
of a larger class of dopant materials, are alloyed with the
aluminum alloy utilized as the parent metal.
[0019] The aforementioned Commonly Owned U.S. Pat. No. 4,853,352
discloses a further development based on the discovery that
appropriate growth conditions as described above, for parent metals
requiring dopants, can be induced by externally applying one or
more dopant materials to the surface or surfaces of the parent
metal, thus avoiding the necessity of alloying the parent metal
with dopant materials, e.g. metals such as magnesium, zinc and
silicon, in the case where aluminum is the parent metal and air is
the oxidant. External application of a layer of dopant material
permits locally inducing metal transport through the oxidation
reaction product and resulting ceramic growth from the parent metal
surface or portions thereof which are selectively doped. This
discovery offers a number of advantages, including the advantage
that ceramic growth can be achieved in one or more selected areas
of the parent metal's surface rather than indiscriminately, thereby
making the process more efficiently applied, for example, to the
growth of the ceramic plates by doping only one surface or only
portions of a surface of a parent metal plate. This improvement
invention also offers the advantage of being able to cause or
promote oxidation reaction product growth in parent metals without
the necessity of alloying the dopant material into the parent
metal, thereby rendering the process feasible, for example, for
application to commercially available metals and alloys which
otherwise would not contain or have appropriately doped
compositions.
[0020] In forming a ceramic composite body, as described in the
aforesaid Commonly Owned U.S. Pat. No. 4,851,375, the parent metal
is placed adjacent a permeable mass of filler material, and the
developing oxidation reaction product infiltrates the mass of
filler material in the direction and towards the oxidant and
boundary of the mass. The result of this phenomenon is the
progressive development of an interconnected ceramic matrix,
optionally containing some nonoxidized parent metal constituents
distributed throughout the growth structure, and an embedded
filler.
[0021] In producing the ceramic composite, any suitable oxidant may
be employed, whether solid, liquid, or gaseous, or a combination
thereof. If a gas or vapor oxidant, i.e. a vapor-phase oxidant, is
used the filler is permeable to the vapor-phase oxidant so that
upon exposure of the bed of filler to the oxidant, the gas
permeates the bed of filler to contact the molten parent metal
therein. When a solid or liquid oxidant is used, it is usually
dispersed through a portion of the bed of filler adjacent the
parent metal or through the entire bed, typically in the form of
particulates admixed with the filler or as coatings on the filler
particles.
[0022] Polycrystalline bodies comprising a metal boride are
produced in accordance with Commonly Owned U.S. Pat. No. 4,777,014,
which issued on Oct. 11, 1988, in the names of Marc S. Newkirk, et
al., and entitled "Process for Preparing Self-Supporting Bodies and
Products Made Thereby". In accordance with this invention, boron or
a reducible metal boride is admixed with a suitable inert filler
material, and the molten parent metal infiltrates and reacts with
the boron source. This reactive infiltration process produces a
boride-containing composite, and the relative amounts of reactants
and process conditions may be altered or controlled to yield a
polycrystalline body containing varying volume percents of ceramic,
metal, reinforcing filler, and/or porosity.
[0023] U.S. Pat. No. 5,202,059 to Kennedy et al. teaches ceramic
filler materials having a plurality of superimposed coatings
thereon. The coated materials are useful as reinforcing materials
in ceramic matrix composites to provide improved mechanical
properties such as fracture toughness. The coatings are selected so
that the interfacial shear strength between the ceramic filler
material and the first coating, between coatings, or between the
outer coating and the surrounding matrix material, are not equal so
as to permit debonding and pull-out when fracture occurs. By reason
of this invention, the coated ceramic filler materials not only
provide improved mechanical properties, but also the filler is
protected from severe oxidizing environments and yet amenable to
the process conditions for the manufacture of the ceramic
composite.
[0024] The entire disclosures of each of the Commonly Owned Patents
and Patent Applications are incorporated herein by reference.
SUMMARY OF THE INVENTION
[0025] In accordance with the present invention, there is disclosed
a plurality of distinct, but combinable, mechanisms for preventing
undesirable oxidation (i.e., oxidation protection mechanisms) of
reinforcement materials (e.g., fibers) in composite bodies. These
oxidation protection mechanisms include the use of getterer
materials which are present in at least a portion of the composite
body (e.g., in at least a portion of the matrix; in at least a
portion of one or more interfacial coatings; or in, on or adjacent
to at least a portion of the reinforcing materials, etc.). These
getterer materials tend to scavenge (e.g., react with) undesirable
oxidants which enter the composite body such as, for example, by
diffusion mechanisms, through microcracks, etc. These oxidation
protection mechanisms may, in certain embodiments, also include
techniques for reducing the number and/or size of such microcracks
in a portion of or throughout the composite body. The reduction in
the amount and/or size of microcracks limits the transport of
undesirable oxidants into and out of the composite body.
[0026] When a composite body is put into service in an oxidizing
environment, and assuming that the oxidizing environment would have
an adverse effect upon the reinforcing material, some type of
oxidation protection mechanism should be utilized to prevent the
reinforcement from oxidizing undesirably. If a getterer material
was placed on, or at least in close proximity to, the reinforcing
material, then an oxidant which came into contact with the getterer
material, such as by diffusion mechanisms, through microcracks,
etc., could be gettered (e.g., reacted) by the getterer materials,
thereby ameliorating undesirable reaction(s) with the reinforcing
material. Further, if the getterer material forms a compound, such
as for example, a glass, the compound could provide even further
oxidation protection to the reinforcing material. In this regard,
if a glass so formed had an appropriate viscosity, then the formed
glass could flow into any microcracks which may be present near the
glass, thereby permitting the glass to function as a crack sealant.
Such desirable compounds are often termed "trap sealants". In this
regard, the formed glass ideally has an oxidant permeability which
is sufficiently low to provide for suitable oxidation protection at
the intended operation temperatures of the composite body for the
desired amount of time.
[0027] In another embodiment of the invention, a glassy material or
a glass-forming material is provided to the composite body during
fabrication.
[0028] The composite body can be engineered so that one or more
getterer materials are included in the composite body such that one
or more desirable compounds (e.g., glasses) are formed. Each of the
getterer materials could react with one or more oxidants at
different temperatures and form one or more desirable compounds
(e.g., one or more desirable glasses) which may provide for
differing amounts of oxidation protection at different
temperatures. In addition, the formed compounds could further react
with other species contained in the composite body to produce
additional desirable compounds. Further, such a formed compound
could react or interact (e.g., alloy) with a glass or glass-former
material which may have been provided to the composite body during
fabrication. Accordingly, a composite body could be produced which
contained a plurality of different oxidation protection mechanisms,
wherein each oxidation protection mechanism was included to provide
for desirable oxidation protection at different service
temperatures of the composite body.
[0029] One exemplary manner of placing an oxidant getterer material
onto a reinforcing material would be to dip, paint or spray an
appropriate material onto at least a portion of the reinforcing
material prior to matrix formation. Alternatively, chemical vapor
deposition (CVD) or chemical vapor infiltration (CVI) techniques
could be utilized to obtain one or more coatings on at least a
portion of, or in a preferred embodiment, substantially all of, a
reinforcing material. It would be desirable for such coatings to be
capable of surviving any matrix formation steps in addition to
providing in-service oxidation protection. Moreover, such coating
could contain, or be modified to contain as, for example, one or
more additional species.
[0030] Such species might function as an oxygen getterer or may
provide oxidation resistance by some other mechanism.
[0031] In a preferred embodiment of the invention, a coated ceramic
filler material, adaptable for use as a reinforcing component in a
ceramic matrix or metal matrix composite, is provided with a
plurality of superimposed coatings. The filler or reinforcing
material useful for this embodiment includes materials where the
length exceeds the diameter, typically in a ratio of at least about
2:1 and more preferably at least about 3:1, and includes such
filler materials as whiskers, fibers, and staple. The coating
system includes a first coating in substantially continuous contact
with the ceramic filler material, and one or more additional or
outer coatings superimposed over the underlying coating, and in
substantially continuous contact therewith. Zonal junctions are
formed between the filler and first coating, between superimposed
coatings, and between the outer coating and the ceramic matrix. The
coatings are selected so that the interfacial shear strength of at
least one of these several zones is weak relative to the other
zones. As used herein and in the appended claims, a zonal junction
is not limited to an interface, per se, between the surfaces but
also includes regions of the coatings in proximity to the
interfaces, and shear, therefore, is zonal in that it may occur at
an interface or within a coating. Further, it is understood that
the zonal junction between adjacent surfaces may be minimal or
negligible and exhibit essentially no bonding or adhesion, or the
adjacent surfaces may exhibit appreciable bonding or a strong bond.
Upon the application of fracture stress to the composite, the weak
zone allows for debonding of the filler before the filler
fractures, and pull-out or shear of the filler upon fracture of the
filler. This debonding and friction pull-out enhances certain
mechanical properties of the composite, and in particular debonding
improves the fracture toughness. Thus, in a duplex coating system,
for example, having a first coating and a second, outer coating
superimposed on the first coating, the coatings are chosen to
facilitate debonding and pull-out such that junction between one of
the three interfaces (i.e. the interface between the filler and the
inner coating, the interface between the inner coating and the
outer coating, the interface between the outer coating and the
surrounding matrix, or the strength of a coating) is weak relative
to the other zonal junctions and allows for debonding and
pull-out.
[0032] By reason of this embodiment of the invention, the coated
ceramic filler materials not only provide improved mechanical
properties, but also the filler is protected from severe corrosive
environments during use and yet amenable to the processing
conditions for making a composite (e.g., matrix formation). For
example, in developing a ceramic matrix by directed metal
oxidation, certain fillers and/or coatings thereon may be at least
partially reduced by the molten parent metal upon contact, and thus
the outer coating protects the filler and inner coating against
this local reducing environment. Thus, duplex coated fillers are
adaptable for use as a reinforcing component in a ceramic matrix
composite formed by the directed oxidation reaction of a molten
precursor metal or parent metal with an oxidant. for many of the
same reasons, such duplex coated filler materials are adaptable for
use in metal matrix composite systems in which the metallic matrix
is formed by infiltration.
[0033] Coated fillers also find utility in composite materials
formed by an infiltration process where the filler material is not
wetted by the infiltrant. In composite systems featuring the melt
infiltration of silicon based metals, the infiltrating silicon
alloy will wet silicon carbide fillers, for example, but does not
readily wet other useful fillers such as ceramic oxides. Moreover,
duplex (or higher order) coated fillers may find utility in
composite systems where the matrix is formed by an infiltration
process but where an inner coating on the filler (provided, for
example, for de-bonding the filler from the matrix) may not be
wetted by the infiltrant material. Boron nitride, for example,
makes a desirable debond coating, but boron nitride is not readily
wet by metals such as aluminum or silicon.
[0034] In another preferred embodiment of the invention, the
coatings may protect the fibers by a means different from, but
possibly in addition to, the above-described mechanisms.
Specifically, under in-service conditions (e.g., at elevated
temperatures), the coatings may help to preserve the "original" or
"as-fabricated" strength of the fibers by preserving the original
character of the fibers, in particular, the fiber chemistry and/or
crystal structural (or lack thereof). Without wishing to be bound
by any particular theory or explanation, it is possible that the
fiber coating serves to prevent or at least retard thermal
decomposition, specifically by preventing, or at least retarding,
outgassing from the fiber, which outgassing may, in some
circumstances, be accompanied by a change in the character of the
crystals making up the fiber such as, for example, through growth
of certain of the crystals or, in the case of an originally
amorphous fiber, by crystallization or devitrification of this
amorphous structure. Specifically, from the perspective of
concentration gradients, a coating on a fiber containing the same
elemental species as the fiber might be expected to retard
diffusion of that species out of the fiber. For example, a carbon
doped boron nitride coating on a silicon carbide based fiber also
containing some oxygen and nitrogen might be expected to reduce the
diffusional loss of carbon and nitrogen from the fiber.
[0035] In general, coated filler materials of this invention may be
utilized in the manufacture of composite materials (e.g., ceramic
matrix composites) that provide improved mechanical properties,
especially increased fracture toughness. When so employed, the
thickness of the coatings should be sufficient to protect the
ceramic filler material against corrosive environments such as
those of molten metals. However, the coatings should not be so
thick as to hinder matrix formation or to interfere with the
function of the filler.
[0036] When relatively thick preforms are to be coated by means of
CVI, it can be a challenge sometimes to adequately coat the filler
in the center of the preform without closing off the pore space
between bodies of filler material residing toward the preform
exterior (e.g., "canning"), thereby rendering the preform
impermeable. Where a preform comprises an assemblage of units, it
has been discovered that arranging the units such that the more
porous, higher permeability units are situated closer to the
preform exterior, and likewise the less porous, lower permeability
units being situated closer to the center of the preform provides
for a more uniform deposit (thickness-wise) throughout the preform
of reaction product from the reactant gases. Thus, if a preform is
to consist of a plurality of woven fabric plies superimposed on top
of one another, it would be desirable from the CVI coating
uniformity perspective to place the fabric plies having the more
"open" weaves on the outside of those fabric plies having a
tighter, less permeable weave.
[0037] It is noted that particular emphasis is herein placed upon
matrices formed by the directed oxidation of a molten metal,
however, certain aspects of the coating composition and/or coating
thickness may be transferable to other matrices (e.g., glass
matrices, etc.) and/or other matrix formation conditions (e.g.,
melt infiltration, chemical vapor infiltration, etc.).
[0038] Moreover, the coatings can be selected so that one or more
of the coatings themselves serves as an oxidant getterer when the
composite is put into service. In a further preferred embodiment,
once the oxidant getterer has formed a compound (e.g., at least one
glassy compound) due to a reaction between the getterer and the
oxidant, the formed compound provides further protection due to,
for example, flowing into a crack to function as a crack sealant.
Still further, the formed compound may interact with (e.g., react,
alloy or modify) a glass to form a different glass which could then
provided oxidation protection in a different temperature
regime.
[0039] In yet another achievement of the invention, an approach for
reducing the number and/or size of microcracks formed during
composite formation and/or formed during composite service or use,
is discussed. Microcracks may be undesirable because such
microcracks may permit ready access of undesirable oxidants to the
reinforcement materials which can result in degradation of some
properties of the composite body. Specifically, microcracking of a
matrix material located between adjacent plies of fiber
tows/bundles (e.g., silicon carbide fibers) can be reduced or
possibly even eliminated by introducing into the matrix one or more
materials having a relatively low coefficient of thermal expansion
such as, for example, silicon carbide particulate. Thus, to
practice this embodiment of the invention, an appropriate material
or combination of materials could be inserted between one or more
fiber tows or between fiber layers to form a preform from a
combination of fibers and particulate. After formation of the
preform, a ceramic matrix comprising, for example, an oxidation
reaction product, could be formed.
[0040] The composite bodies of the present invention do not require
a seal coat applied over the exterior of the bulk body.
Accordingly, the composite bodies of the present invention are
adaptable to finishing operations such as machining, polishing,
grinding, etc. The resultant composites are intended to include,
without limitation, industrial, structural, and technical ceramic
bodies for applications where improved strength, toughness and wear
resistance are important or beneficial.
[0041] While this disclosure focuses primarily on ceramic matrix
composite bodies having a matrix formed by the directed oxidation
of a molten metal, it should be understood that the coating
techniques of the invention are by themselves novel and useful and
have industrial applicability in many other composite body
formation processes (e.g., other ceramic matrix composite formation
techniques, glass matrix formation techniques, polymer matrix
formation techniques, metal matrix formation techniques, etc.).
Accordingly, this invention also relates to the specific techniques
for forming such coatings.
Definitions
[0042] The following terms, as used herein and in the claims, have
the stated meanings as defined below:
[0043] The term "oxidation reaction product" in conjunction with
both oxidation reaction product growth and gettering means one or
more metals in any oxidized state wherein the metal(s) have given
up electrons to or shared electrons with another element, compound,
or combination thereof. Accordingly, an "oxidation reaction
product" under this definition includes the product of the reaction
of one or more metals (e.g., a parent metal comprising aluminum,
silicon, tin, titanium, zirconium, etc.) with an oxidant such as
oxygen or air, nitrogen, a halogen, sulfur, phosphorous, arsenic,
carbon, boron, selenium, tellurium; compounds such as silica (as a
source of oxygen), and methane, ethane, propane, acetylene,
ethylene, and propylene (as a source of carbon); and mixtures such
as H.sub.2/H.sub.2O and CO/C).sub.2 which are useful in reducing
the oxygen activity of the environment.
[0044] The term "oxidant" means one or more suitable electron
acceptors or electron sharers and may be a solid, liquid, or gas
(vapor) or some combination of these. Thus, oxygen (including air)
is a suitable vapor-phase gaseous oxidant for the formation of
oxidation reaction product, with air being preferred for reasons of
economy. Boron, boron carbide and carbon are examples of solid
oxidants for the formation of oxidation reaction product under this
definition.
[0045] The term "parent metal" refers to that metal, e.g. aluminum,
which is the precursor of a polycrystalline oxidation reaction
product such as alumina, and includes that metal or a relatively
pure metal, a commercially available metal having impurities and/or
alloying constituents therein, and an alloy in which that metal
precursor is the major constituent; and when a specified metal is
mentioned as the parent metal, e.g. aluminum, the metal identified
should be read with this definition in mind unless indicated
otherwise by the context.
[0046] The term "ceramic" is not limited to a ceramic body in the
classical sense, that is, in the sense that it consists entirely of
non-metallic, inorganic materials, but rather, it refers to a body
which is predominantly ceramic with respect to either composition
or dominant properties, although the body may contain substantial
amounts of one or more metallic constituents such as derived from
the parent metal, most typically within a range of from about 1-40%
by volume, but may include still more metal.
[0047] The term "glass" or "glassy compound" as used in this
disclosure broadly refers to inorganic materials exhibiting only
short-range order, e.g., non-crystalline character. The term thus
includes the traditional "glass-forming oxides" such as silica and
boria, but also includes materials which normally exhibit
long-range order, e.g., crystallinity, but which order has been
disrupted through rapid solidification or the presence of defects
such as impurity atoms.
[0048] The term "inorganic polymer" or "preceramic polymer" refers
to that class of polymeric materials which upon pyrolysis convert
to ceramic materials. Such polymers may be solid or liquid at
ambient temperature. Examples of these polymers include the
polysilazanes and polycarbosilazanes which can be pyrolyzed to
yield ceramic materials comprising silicon nitride and silicon
carbide, respectively.
[0049] The term "melt infiltration" refers to a technique for
producing composite materials by infiltration whereby a molten
metal comprising silicon is placed into contact with a permeable
mass which is wettable by the molten metal, and the molten metal
infiltrates the mass without the requirement for the application of
pressure or vacuum. The infiltration may occur with significant
reaction or with substantially no reaction of infiltrating metal
and one or more components of the permeable mass.
[0050] The term "reaction-formed" in the context of silicon carbide
refers to a melt infiltration process whereby silicon in the
infiltrant metal reacts with a carbon source in the permeable mass
to produce silicon carbide in the matrx phase of the resulting
composite body. This in-situ formed silicon carbide may or may not
be interconnected.
[0051] The term "trap sealant", as used herein, refers to a
chemical species which is capable of gettering oxygen, and upon so
doing, forms or contributes to oxide glass formation, which glass
is capable of interfering with oxygen gas transport.
BRIEF DESCRIPTION OF THE DRAWINGS
[0052] FIG. 1 is a scanning electron micrograph taken at about
350.times.magnification of a coated ceramic filler material in a
ceramic matrix and made according to the invention.
[0053] FIG. 2 is a scanning electron micrograph taken at about
850.times.magnification of ceramic matrix composite having a coated
NICALON.RTM. ceramic fiber as filler material and made according to
the Example below.
[0054] FIG. 3 is a scanning electron micrograph taken at
250.times.magnification of a fractured surface of the composite
made with the coated fibers according to the Example below showing
extensive pull-out of the fibers.
[0055] FIG. 4 is a scanning electron micrograph taken at
800.times.magnification of a fractured surface of the composite
made with uncoated fibers according to the Example below showing no
pull-out of the fibers.
[0056] FIG. 5a is a schematic of the top view of harness satin
weave fabric in the as is position as discussed in Example 2.
[0057] FIG. 5b is a schematic cross-sectional representation of a
harness satin weave fabric in the as-is position as discussed in
Example 2.
[0058] FIG. 5c is an isometric schematic view illustrating the axes
of rotation for a harness satin weave fabric in the as-is position
as discussed in Example 2.
[0059] FIG. 5d is a schematic cross-sectional representation of a
fabric preform comprised of harness satin fabric as discussed in
Example 2.
[0060] FIG. 5e is an isometric schematic representation of a
graphite containment fixture for effecting the coating of a fabric
preform as discussed in Example 2.
[0061] FIG. 5f is a isometric schematic representation of a
cantilever graphite fixture for holding a boron nitride coated
fabric preform to enable coating of the preform with a second
coating as discussed in Example 2.
[0062] FIG. 5g is a schematic cross-sectional representation of a
growth lay-up for forming a fiber reinforced ceramic composite body
as discussed in Example 2.
[0063] FIG. 5h is a schematic cross-sectional representation of a
lay-up for removing the metallic component of the formed fiber
reinforced ceramic composite body discussed in Example 2.
[0064] FIG. 6 is a schematic cross-sectional representation of a
typical lay-up for removing at least one metallic constituent of a
metallic component from substantially all surfaces of a composite
body.
[0065] FIG. 7 is an orthoscopic view of tensile and stress rupture
test specimens.
[0066] FIG. 8 is a typical stress-strain curve for a
fiber-reinforced ceramic composite tensile test specimen.
[0067] FIG. 9 is a SEM photograph at about 50.times.magnification
of the fracture surface of a tensile test specimen.
[0068] FIG. 10 shows tensile strength of a fiber-reinforced ceramic
matrix composite vs. T(.degree. C.) in air.
[0069] FIG. 11 shows tensile strength vs. temperature for thermally
cycled and non-thermally cycled fiber ceramic matrix composite test
specimens.
[0070] FIG. 12 shows results of stress rupture testing of
NICALON.RTM. fiber reinforced Al.sub.2O.sub.3 at 1000, 1100 and
1200.degree. C. in air.
[0071] FIG. 13 is a SEM photograph at about 50.times.magnification
of the fracture surface of a stress rupture tested specimen.
[0072] FIG. 14 is a scanning electron micrograph taken at about
25000.times.magnification of a polished cross-section of Sample H
near the rupture surface.
[0073] FIG. 15 is a scanning electron micrograph taken at about
10,000.times.magnification of a polished cross-section of Sample H
near the rupture surface.
[0074] FIG. 16 shows total strain vs. time for a 1100.degree. C.
stress rupture specimen at about 70 MPa tensile load in air.
[0075] FIG. 17a is an approximately 50.times.magnification optical
photomicrograph of a polished cross-section of a NICALON.RTM. fiber
reinforced ceramic matrix composite revealing the presence of
several microcracks in the matrix material between adjacent fiber
tows.
[0076] FIG. 17b is an approximately 50.times.magnification optical
photomicrograph of a polished cross-section of a fiber reinforced
ceramic composite body which shows how additions of silicon carbide
particulate placed between adjacent plies of NICALON.RTM. fiber
virtually eliminates these matrix microcracks.
[0077] FIGS. 18a and 18b are isometric drawings of the graphite
support fixture of Example 15 loaded with fabric preforms and in
the unloaded condition, respectively.
[0078] FIG. 19 is an S-N plot showing the life of a fiber
reinforced ceramic composite body as a function of temperature and
the maximum applied cyclical tensile stress.
[0079] FIG. 20 is a plot of sample strain versus time for a fiber
reinforced ceramic composite body subjected to thermal cycling
under an applied tensile dead load.
[0080] FIG. 21 shows the four point flexural strength of a fiber
reinforced ceramic composite body as a function of the atomic
percentage (ratio) of silicon to boron in the reactant gases used
to deposit by CVD a modified boron nitride coating layer on the
fibrous reinforcement of the composite body.
Detailed Description of the Invention and Preferred Embodiments
[0081] By way of review, in composite material systems,
particularly ceramic matrix composite systems, frequently it is
desirable for the reinforcement phase to debond and pull away or
pull out of the matrix, at least partially. Such debonding and pull
out absorbs mechanical energy which might otherwise have gone into
fracturing the composite body. Typically one or more coatings are
applied to the reinforcement material, e.g., the fibers, to
accomplish the debonding under applied load. Not only are the
composite fabrication conditions (e.g., matrix development) harsh
from a chemical corrositivity point of view, so are the end use
conditions, generally. Corrosion of the reinforcement or the debond
coating(s) becomes a concern because chemical reaction ordinarily
renders the reinforcement or the debond coating(s) less effective.
Thus, the concept of the duplex coating was developed: a debond
coating on a fiber itself coated with (and thereby protected by) a
refractory material.
[0082] Suitable ceramic filler materials which may be used in the
invention include metal oxides, borides, carbides, nitrides,
silicides, and mixtures or combinations thereof, and may be
relatively pure or contain one or more impurities or additional
phases, including composites of these materials. The metal oxides
include, for example, alumina, magnesia, calcia, ceria, hafnia,
lanthanum oxide, neodymium oxide, samaria, praseodymium oxide,
thoria, urania, yttria, beryllium oxide, tungsten oxide and
zirconia. In addition, a large number of binary, ternary, and
higher order metallic compounds such as magnesium-aluminate spinel,
silicon aluminum oxynitride, borosilicate glasses, and barium
titanate are useful as refractory fillers. Additional ceramic
filler materials may include, for example, silicon carbide, silica,
boron carbide, titanium carbide, zirconium carbide, boron nitride,
silicon nitride, aluminum nitride, titanium nitride, zirconium
nitride, zirconium boride, titanium diboride, aluminum
dodecaboride, and such materials as Si--C--O--N compounds,
including composites of these materials. The ceramic filler may be
in any of a number of forms, shapes or sizes depending largely on
the matrix material, the geometry of the composite product, and the
desired properties sought for the end product, and most typically
are in the form of whiskers and fibers. The fibers can be
discontinuous (in chopped form as staple) or in the form of a
single continuous filament or as continuous multifilament tows.
They also can be in the form of two- or three-dimensional woven
continuous fiber mats or structures. Further, the ceramic mass may
be homogeneous or heterogeneous.
[0083] In a major aspect of the present invention, the oxidation
protection mechanisms of the invention include the use of getterer
materials which are present in at least a portion of the composite
body (e.g., in at least a portion of the matrix; in at least a
portion of one or more interfacial coatings; or in, on or adjacent
to at least a portion of the reinforcing materials, etc.). These
getterer materials tend to scavenge (e.g., react with) undesirable
oxidants which enter the composite body such as, for example, by
diffusion mechanisms, through microcracks, etc. These oxidation
protection mechanisms may, in certain embodiments, also include
techniques for reducing the number and/or size of such microcracks
in a portion of or throughout the composite body. The reduction in
the amount and/or size of microcracks may limit the ability of
undesirable oxidants to negatively impact the reinforcement
material(s) in the composite body.
[0084] When a composite body is put into service in an oxidizing
environment, and assuming that the oxidizing environment would have
an adverse effect upon the reinforcing material, some type of
oxidation protection mechanism should be utilized to prevent the
reinforcement from oxidizing undesirably. If a getterer material
was placed on, or at least in close proximity to, the reinforcing
material, then an oxidant which came into contact with the getterer
material such as, for example, by diffusion mechanisms through
microcracks, etc., could be gettered (e.g., reacted) by the
getterer materials, thereby ameliorating undesirable reaction(s)
with the reinforcing material. Further, if the getterer material
forms a particular compound, such as for example, a glass, the
compound could provide even further oxidation protection to the
reinforcing material. In this regard, if a formed glass had an
appropriate viscosity, then the formed glass could flow into any
microcracks which may be present near the formed glass, thereby
permitting the formed glass to function as a crack sealant. Such a
desirable compound is sometimes referred to as a "trap sealant." In
this regard, the formed glass should have an oxidant permeability
which is low enough to provide for suitable oxidation protection of
the composite body at the intended operation temperatures for a
desirable amount of time (e.g., the intended lifetime of the
composite body).
[0085] In another embodiment of the invention, a glassy material,
glass-network-forming material or glass modifier material is
provided to a composite body during fabrication. For example, one
can envision coating woven ceramic fiber plies with a particulate
slurry comprising a glass-former such as silica and, optionally,
one or more structural modifiers such as alumina, zirconia, calcia,
etc.
[0086] A number of candidate getterer materials useful in
combination with various matrices and reinforcements will become
apparent to an artisan of ordinary skill upon review of this
disclosure. Specifically, in a preferred embodiment of the
invention, many reinforcement materials (e.g., fibers) are
susceptible to oxidation by oxidants such as oxygen. Accordingly,
it often is vitally important to prevent oxygen from contacting the
reinforcing fibers so as to prevent any negative effects upon the
fibers. In this regard, oxygen typically is transported to a fiber
surface by a combination of different mechanisms. In general,
oxygen usually enters the surface of a composite body due to some
flaw present on the surface (e.g., machining marks, a broken or
cracked outer protective skin, etc.). Once the oxygen has permeated
the surface of a composite body, oxygen may then ingress further
into the composite body by various channels present in the
composite body due to microcracking from processing, thermal shock,
physical shock, etc. In addition, molecular or atomic oxygen
diffusion may also occur in combination with the physical ingress
of oxygen into the composite body. If an appropriate oxygen
getterer material was positioned such that the oxygen which
ingressed into the composite could be gettered (e.g., reacted with)
by the oxygen getterer, then further ingress of that particular
oxygen molecule would be inhibited. However, if additional oxygen
ingressed into approximately the same area in the composite, at
some point substantially all of the oxygen gettering material will
eventually react with the ingressing oxygen. At that point, it
would be desirable for another oxidation protection mechanism to
occur. In this regard, if the oxygen gettering material were chosen
so that one or more desirable compounds (e.g., oxides or glasses)
were formed upon a reaction with the oxygen, then such glasses or
other oxides could block (e.g., flow into) any cracks, channels,
microcracks, etc., to inhibit the physical transport of oxygen
further into the composite body.
[0087] Examples of materials which function as suitable oxygen
getterers and glass formers are boron, silicon, and the carbides
and nitrides of boron and silicon. When reacted with oxygen the
boron containing species may form a boria based glass and the
silicon containing species may form a silica based glass. Moreover,
it is possible that when both boron oxide glass and silicon dioxide
glass are present, the glasses may exist independently and/or may
form a borosilicate glass. Still further, if additional materials
are present in the vicinity of the forming glasses, such as
aluminum (e.g., as a metal or an oxidized compound such as
Al.sub.2O.sub.3) and/or zirconium, in various forms both oxidized
and non-oxidized, etc., it is possible to form in addition to those
glasses mentioned above, glasses such as zirconium borosilicates,
aluminum borosilicates, etc.
[0088] Thus, it should be apparent that one or more oxygen getterer
materials can be included in a composite body to form a number of
desirable compounds, such as those glasses discussed immediately
above. In this regard, it is possible to design a composite body so
that when a composite body is subjected to use in an oxidizing
environment, a first glass, such as a low melting boron oxide or
borosilicate glass, will form and protect the reinforcing material
of the composite at low temperatures. As the temperature of the
composite body is increased, it is possible to form more refractory
or higher softening point glasses which may result in oxidation
protection at even higher service temperatures. For example, a high
viscosity or high softening point glass such as a zirconium
borosilicate may extend the service life of a composite body to
heretofore believed to be impossible times at elevated
temperatures. It also may be necessary to provide oxidation
protection at intermediate temperatures. In this regard, it may be
desirable to form a glass such as an aluminum borosilicate which
would bridge the gap in service temperature between, for example,
the lower viscosity boron oxide glasses and the higher viscosity
glasses such as zirconium borosilicate. As is apparent from the
above discussion, the number of combinations of oxygen gettering
materials which can form desirable glasses, which may or may not
react with other materials in the composite body, is quite
large.
[0089] Further, it should be apparent to an artisan of ordinary
skill that desirable glasses need not be formed entirely from the
action of oxygen getterers. Instead, the desired glass or its
components (glass-formers, modifiers, etc.) can be incorporated
into the composite body during composite fabrication. In service,
glassy particulates may fuse to one another and flow into cracks.
Glass formers and modifiers may alloy and/or react to form the
desired glass. The compounds formed by "spent" oxygen getterers may
also particiapte to modify these glasses originally formed without
oxygen getterer involvement.
[0090] Further, an important criterion in selecting materials which
function to getter oxygen is the viscosity and oxygen permeability
of the glassy material which is to be formed or modified due to
reaction or alloying with an oxidized or "spent" gettering
material. For example, in a silicon carbide fiber reinforced
aluminum oxide material, an oxygen gettering material which could
be coated onto the fibers and subsequently form a glass may need to
be such that the glass so formed has an oxygen permeability of
about 1.times.10.sup.-9 g-O.sub.2/cm.sup.2 sec in order for the
composite body to survive a few hours. However, if it is important
for the composite body to survive thousands of hours, the oxygen
permeability may need to be even lower; for example, about
1.times.10.sup.-12 g-O.sub.2/cm.sup.2 sec may be necessary. By way
of comparison, a microcrack may exhibit an effective permeability
of 1.times.10.sup.-6 g-O.sub.2/cm.sup.2 sec or more. It is of
course apparent that oxygen permeability is a function of
temperature and an artisan of ordinary skill would need to
determine the precise service temperature or temperatures that a
composite body would be exposed to during service to determine the
best combination of oxygen gettering and glass forming materials to
be used to extend the useful life of the composite body.
[0091] Another factor to consider in designing an oxygen getterer
system which possesses glass sealing characteristics is the effect
of moisture, particularly at elevated temperatures.
Specifically,boron oxide (e.g., B.sub.20.sub.3) glass dissolves in
water according to the formula
B.sub.20.sub.3+3H.sub.20<->2H.sub.3B0.sub.3
[0092] At elevated temperature (e.g., 900.degree. C.)
H.sub.3B0.sub.3 is in the vapor phase. Thus, exposure of
B.sub.20.sub.3 glass to water vapor at such temperatures causes the
volatilization of the former. The reactivity/solubility of
borosilicate glasses with water is much less than that of straight
boron oxide glass. Thus, all other things being equal, it may be
better to design a materials system to produce borosilicate glasses
than unmodified boria glass.
[0093] In general, oxygen gettering materials which form boron
oxide or borosilicate glasses provide for relatively low
temperature oxidation protection (e.g., less than about 600.degree.
C.); however, oxygen gettering materials which form a calcium
aluminosilicate glass may provide intermediate temperature
oxidation protection (e.g., about 600.degree. C.-1200.degree. C.);
oxygen getterers that form silicate glasses may provide
intermediate to high temperature oxidation protection (e.g., about
600.degree. C.-1800.degree. C.); oxygen gettering materials which
form a zirconium silicate glass or zircon structure may provide
high temperature oxidation protection (e.g., about 1200.degree.
C.-1800.degree. C.); and oxygen gettering materials which form
zirconia and silica glasses may provide for very high temperature
oxidation protection (e.g., about 1800.degree. C.-2200.degree.
C.).
[0094] Accordingly, it is apparent that a composite body can be
engineered so that one or more getterer materials are included in
the composite body such that one or more desirable compounds (e.g.,
glasses) are formed. Each of the getterer materials could react
with one or more oxidants at different temperatures and form one or
more desirable compounds (e.g., one or more desirable glasses)
which may provide for differing amounts of oxidation protection at
different temperatures. Accordingly, a composite body could be
produced which contained a plurality of different oxidation
protection mechanisms, wherein each oxidation protection mechanism
was included to provide for desirable oxidation protection at
different service temperatures of the composite body.
[0095] One exemplary manner of placing an oxidant getterer onto a
reinforcing material or at least in close proximity thereto would
be to dip, paint or spray an appropriate material onto at least a
portion of the reinforcing material prior to matrix formation or
onto at least a portion of another material which is in contact
with the reinforcing material. Alternatively, or in conjunction
with such coating by dipping, painting or spraying, chemical vapor
deposition (CVD) or chemical vapor infiltration (CVI) techniques
could be utilized to obtain one or more coatings on at least a
portion of, or in a preferred embodiment, substantially all of, a
reinforcing material. For example, a first coating comprising boron
nitride could be deposited onto a reinforcing material by CVI. One
or more oxidant getterer materials might then be applied to the
boron nitride coating by dip coating, for example, the coated
reinforcing material into a solution comprising, for example, the
nitrates or acetates of silicon, aluminum, zirconium and/or
yttrium, which dip coated reinforcing material could then be heated
in a nitrogen atmosphere, for example, to convert the nitrates to
nitrides. It would be desirable for such coatings to be capable of
surviving any matrix formation steps as well as providing
in-service oxidation protection. If necessary, one or more
additional coatings comprising, for example, silicon carbide could
then be applied, for example, by CVI to protect the underlying
coatings and reinforcing material from chemical degradation during
subsequent processing.
[0096] CVD or CVI is a particularly desirable means of placing one
or more oxidant getterers in proximity to a reinforcing material.
Specifically, the precise location of the oxidant getterer in
relation to the reinforcing material may be highly controlled. For
example, if a first coating comprising boron nitride is to be
applied, an oxidant getterer comprising aluminum or silicon could
be applied as aluminum nitride or silicon nitride, respectively,
using CVI. Further, the oxidant getterer could be applied before,
during and/or after the deposition of one or more coatings to the
reinforcing material to produce a coated reinforcing material
having an oxidant getterer underneath, mixed within (e.g.,
intermixed) and/or on top of (e.g., exterior to) the coatings.
Moreover, it is possible to apply different oxidant getterer
materials at different locations relative to the reinforcing
material. For example, one may choose to deposit, for example, an
aluminum nitride oxidant getterer beneath a first coating
comprising boron nitride and a zirconium nitride oxidant getterer
on top of this first coating and/or on top of a second coating
comprising silicon carbide. Still further, two or more oxidant
getterer materials may be simultaneously co-deposited using CVD or
CVI, such as, for example, simultaneous depositions of oxidant
getterers comprising aluminum and zirconium as their respective
nitrides. Depending upon conditions and choice of coating materials
to be deposited, it is even possible to simultaneously deposit one
or more of the oxidant getterer materials with the coatings for the
filler material. For example, oxidant getterers (e.g., aluminum,
silicon, yttrium, zirconium, etc.) may be co-deposited during
deposition of the boron nitride debond coating onto the reinforcing
filler material. Finally, through careful control of the reactant
gas concentrations, a graded or tailored concentration of one or
more oxidant getterer materials can be achieved within a
coating.
[0097] Without wishing to be bound by any particular theory or
explanation, it has been observed that boron nitride doped with
silicon exhibits increased oxidation resistance, particularly where
moisture is also present. A convenient technique for producing such
silicon doped boron nitride is by CVD, specifically by providing
boron, nitrogen and silicon sources. Seemingly the silicon would
chemically react with the nitrogen source to produce silicon
nitride. Silicon nitride and boron nitride co-deposition is not
thermodynamically favorable at low temperatures, so to achieve a
significant presence of silicon in the boron nitride deposit, the
co-deposition may need to be conducted at high temperatures, for
example at or above 1200.degree. C. Because reaction rates tend to
increase with increasing temperature, the precipitation of solid
reaction product is rapid at such temperatures. The rapid
deposition rates may not pose a problem for coating single
filaments or fiber tows. However, it may be difficult or impossible
to uniformly coat a fabric or stack of fabrics or a
three-dimensionally woven fiber preform under such conditions
without the bulk of the deposit residing on the exterior of the
preform and potentially sealing off the interior regions.
[0098] From a uniformity of coating deposition standpoint, slow
deposition is better than rapid deposition. Low reaction
temperatures are conducive to slower deposition rates. As stated
above, low deposition temperatures are not conducive
thermodynamically to co-depositing silicon nitride with boron
nitride. Fortunately, it is still possible to co-deposit a few
percent of silicon along with the balance of boron nitride at the
low temperatures, (e.g., about 700.degree. C. to 800.degree. C.)
although it is not clear if silicon nitride is in solid solution
with boron nitride or even if the silicon is present as silicon
nitride. More fortunate still has been the discovery that even
small amounts of silicon co-deposited with boron nitride can have a
large beneficial effect on the resistance to environmental
degradation of this silicon modified boron nitride coating.
[0099] It should be understood that the thickness of any coating
which may be applied to a reinforcing material influences a number
of different properties, including the mechanical properties of a
composite body, at both ambient temperature and elevated
temperatures, as well as the amount of oxidation protection
afforded the reinforcing material. In general, the thickness of
coatings on fibers in ceramic matrix composite bodies, where the
ceramic matrix composite bodies are to be subjected to elevated
temperature environments, should be from a few tenths of a micron
thick to a few tens of microns in thickness and even more
preferably about 0.2 to about 20 microns in thickness.
Specifically, if a fibrous reinforcing material is chosen, the
thickness of the coating on the fiber should be sufficient to
permit fiber pull-out to occur. Thicknesses greater than a few tens
of microns may result in adverse degradation of mechanical
properties (e.g., a coating which is too thick may cause a failure
mode to change from one which is predominantly fiber pull-out to a
different failure mode which could have an overall weakening effect
on the composite body), whereas thicknesses less than a few tenths
of a micron may not provide for adequate oxidation protection of
the underlying fibers and/or not permit fiber/matrix debonding to
occur (e.g., if a thickness of coating was too thin, fibers may be
bonded too strongly to the matrix thus inhibiting fiber pull-out
mechanisms from occurring). Accordingly, numerous considerations
need to be taken into account when selecting the thickness of one
or more coatings to be placed upon a fiber reinforcement in a
composite body.
[0100] In another aspect of the present invention, and particularly
in regard to forming a ceramic matrix composite body by a directed
metal oxidation of a parent metal, it has been discovered that a
useful filler material or strengthening component for the ceramic
matrix composite body should be provided with two or more coatings.
The first or inner coating is applied to the filler as a continuous
film or layer, and preferably forms a bond with the filler. The
second and any subsequent coatings are superimposed over an
underlying coating and become attached or bonded therewith as
additional layers or stratum. Each coating is applied as a
substantially continuous layer, and each is in substantially
continuous contact with the underlying coating or filler in the
case of the first coating. The bond formed between adjacent
surfaces may be weak or negligible in that there may be little or
no adhesion or connection, but in the preferred embodiment there is
a measurable or appreciable bonding or union between surfaces.
[0101] In the embodiment of the invention in which multiple
coatings are called for, two coatings applied to the filler
material are normally sufficient. In such a system utilizing a
duplex coating, the coatings are selected to provide adequate
mismatch in bonding strengths so as to allow for debonding and
pull-out upon application of stress. Also, the duplex coating is
selected to provide protection against degradation of the filler,
and the outer coating is selected to exhibit wettability of molten
parent metal and to protect the inner coating from degradation or
corrosion in high temperature, oxidizing environments under the
conditions of the matrix formation process. Also, a system using
two coatings rather than three or more, may be somewhat more
advantageous from an economic standpoint.
[0102] Thus, the coatings are selected so as to be compatible with
the filler material, and to the process conditions for the
manufacture of the composites. Also, the coatings should complement
each other in achieving the desired characteristics or properties.
In a ceramic composite system having incorporated therein a filler
with a duplex coating, for example, the first and outer coatings
are selected to provide an adequate mismatch in interfacial shear
strength so that one of the three zonal junctions is weak relative
to the remaining zonal junctions to provide relative movement
between the inner coating and the filler, or between coatings, or
between the outer coating and the adjacent ceramic matrix. In this
manner, debonding and pull-out should occur, thereby improving or
enhancing the fracture toughness of the ceramic composite body.
[0103] Debonding and pull-out is especially beneficial for filler
materials having a relatively high length to diameter ratio, such
as fibers, typically at least about 2:1 and more particularly at
least 3:1. Filler material with a low length to diameter ratio such
as particles or spheres, characteristically exhibits crack
deflection toughening.
[0104] In applying the coatings to the filler material, the
thickness of each coating and the cumulative thickness of all
coatings can vary over a wide range. This thickness can depend on
such factors as the composition of each coating and their
interaction, the type and geometry of the filler, and the process
conditions and, for example, the parent metal used in the
manufacture of the composite. Generally, the cumulative thickness
for the coatings should be sufficient to completely cover the
ceramic filler material and protect it from, for example, oxidation
degradation, attack from molten metal, and other corrosive
environments which may be encountered in employment of the finished
composite. In the preferred embodiment, the inner coating is
compatible with the filler material so as not to degrade its
integrity, and further the inner coating can be selected to allow
for debonding and pull-out or shear. The coating system is selected
to be compatible with the matrix material, especially the precursor
for the matrix, and further the coating system is selected so as to
be capable of withstanding the process conditions used in the
manufacture of the composites. While the inner coating may afford
adequate protection against degradation of the filler or allow for
shear between this first coating and the filler, a second or outer
coating is selected to be compatible with the process conditions
employed in the manufacture of the ceramic composite body, in that
it should be substantially inert and not degrade, and further
should exhibit wettability to molten parent metal when serving as a
precursor to the ceramic matrix. Also, if the first coating or
fiber is susceptible to attack and degradation by the process
environment during composite manufacture or by attack of oxidants
diffusing through the matrix during actual service, the second or
outer coating is chosen to protect the inner coating or fiber from
exposure to processing conditions and/or end use conditions (e.g.,
the inner coating may function as an oxygen getterer material alone
or in combination with other components of the composite body such
as other coatings or other materials in the composite body). Thus,
the coating system protects the fibers from degradation, as does
one coating superimposed on another, and concomitantly provides for
compatibility for matrix formation and use, and for relative
movement to allow for shear. By reason of this coating system,
structural degradation of the composite components is mitigated
thereby prolonging the useful life and performance of the
composite, and the fracture toughness of the composite is
improved.
[0105] If the surface of a fibrous filler material is very
irregular and exhibits nodules, barbs, fibrils, projections, or
protuberances, the fiber can mechanically interlock or bond with
the adjacent surface including the adjacent coating or adjacent
fiber thereby impeding or preventing debonding and pull-out, which
can be deleterious to the properties of the composite. It therefore
is desirable to provide a coating system which is sufficiently
thick to completely cover the irregularities in the fibers. Again,
when large numbers of fibers or filaments are being coated at the
same time, the coating cannot be so thick as to isolate the fibers
in the middle of a bundle from those near the exterior.
[0106] The thickness and properties of the coatings may vary
depending on the deposition process and the filler material. In a
duplex coating system, the thickness for each coating, as measured
from the center of a filler material body out normal to the surface
of the body, typically may range from about 0.05 to about 25
microns, preferably to about 10 microns, but the innermost coating
can be as thin as a single monolayer in order to separate the
second coating from the filler particle. The cumulative thickness
for a coating system may be to about 25 microns, and more
preferably 2-10 microns. Usually a coating system having a
thickness within this range can be applied to the filler by
conventional or known means and will provide the desired properties
described above.
[0107] It has been found that a number of coating compositions can
be employed in the coating system of this invention. These
compositions include the metal oxides, nitrides, borides and
carbides, alkaline metal salts, alkaline earth metal salts, carbon,
silicon, and the like. The choice of coating compositions will
depend on the filler material, the compatibility of coatings to
each other, and the process conditions for the manufacture of the
ceramic composite. For example, silicon carbide fibers are a
popular choice as filler in composites intended for use at elevated
temperatures. In order to provide for debonding and pull-out, the
silicon carbide fibers may be coated with boron nitride which
prevents a relatively strong bond between the coated fiber and the
surrounding matrix. However, boron nitride may be degraded by the
oxidation reaction conditions associated with a directed metal
oxidation process. Further, boron nitride may not be wet by certain
metals, such as aluminum or silicon, under the conditions of the
matrix formation process by infiltration (e.g., directed metal
oxidation, melt infiltration, etc.), and therefore as an outer
coating would tend to interfere with the matrix formation. However,
an inner coating exhibiting little or no wettability by the
infiltrant metal under process conditions can be advantageous. For
example, the coating system may have pores or flaws, but the
contact angle of the molten infiltrant metal with the inner coating
may preclude transport of the metal through any pores or flaws in
the inner coating and thereby protect the filler from attack by
molten metal. The presence of an additional wettable outer coating
on the filler would then avoid impedance to the matrix formation
process. Therefore, a suitable outer coating such as silicon
carbide is applied to the boron nitride coating to achieve
compatibility with the forming process and to protect the boron
nitride from degradation, such as by oxidation. Silicon carbide is,
for example, wet by doped aluminum and relatively
oxidation-resistant in an air environment at 1000.degree. C.,
whereas boron nitride is typically not wet by aluminum, and is
oxidation-prone, at this temperature. Further, the bond between the
two coatings is weak relative to the other bonds thereby
facilitating debonding and pull-out of the fibers during fracture.
Other useful coating compositions include, for example, titanium
carbide, silicon, calcium silicate, calcium sulfate, and carbon as
the inner coating, and silicon, silica, alumina, zirconia,
zirconium nitride, titanium nitride, aluminum nitride, and silicon
nitride as an outer coating. Other suitable compositions for the
first and outer coatings may be selected for use with the ceramic
filler material provided these coatings complement each other as in
the manner described above.
[0108] A typical cross-sectional representation of the coated
ceramic filler material is shown in FIG. 1 (discussed below in
greater detail). In this typical example, the ceramic filler
material comprising silicon carbide bears a first inner coating of
boron nitride and an additional outer coating of silicon carbide,
thus a duplex coating. One or more additional outer coatings may be
provided depending on need. For example, an additional outer
coating of titanium carbide may be applied to the coating of
silicon carbide.
[0109] Moreover, it may be desirable to provide dual or multiple
duplex coatings such as boron nitride/silicon carbide/boron
nitride/silicon carbide. This multiple coating scheme may result in
desirable internal oxidation protection mechanisms. Specifically,
as discussed above, the interface between boron nitride and silicon
carbide may function as a zonal debond junction, thus increasing
the fracture toughness of a material, as well as providing for
oxidation protection. As discussed above, the precise composition
and combination of coatings depends on a number of factors
including the processing or manufacturing environment for the
composite body as well as the environment into which the composite
body will be placed.
[0110] Non-oxide ceramic materials tend to decompose in the
presence of oxygen at elevated temperatures. This problem can be
particularly acute for high surface-to-volume geometries such as
that of a fiber. Whether by choice of design or by circumstances,
many commercially available non-oxide ceramic fibers contain at
least minor amounts of impurity materials. For example, in the case
of a stabilized silicon carbide fiber such as NICALON.RTM. fiber,
the impurity materials comprise oxygen and nitrogen. These
impurities can stabilize or have the effect of potentially
stabilizing the NICALON.RTM. fiber as manifested by preserving a
substantial fraction of such a fiber's ambient temperature strength
up to elevated temperatures. Specifically, the oxygen and/or
nitrogen atoms occupy positions between microcrystalline silicon
carbide grains. The effect of these impurities is to increase high
temperature tensile strength and reduce high temperature creep.
Nevertheless, left unprotected at elevated temperatures,
NICALON.RTM. silicon carbide fiber eventually loses a substantial
portion of its original or as-fabricated tensile strength, even
when the exposure is conducted in a non-oxidizing environment such
as, for example, in an argon atmosphere. Concurrent with this
strength loss, a mass loss from the fiber and, in particular, the
oxygen and/or nitrogen impurities, has also been observed, such
mass being lost through, for example, volatilization. Without
wishing to be bound by any particular theory or explanation, it
seems that in at least one instance this loss of the "stabilizing"
impurities permits the further crystallization and growth of
crystals within the fiber. The fiber then consists of an assemblage
of crystallites or grains of, for example, silicon carbide, having
distinct grain boundaries. If such grains continue to grow upon
continued high temperature exposure of the fiber, the tensile
strength of such a treated fiber may decrease with respect to the
same fiber in its original form. This strength loss can be
attributed, at least in part, to the development and growth of
grain boundaries which may have or exhibit a strength limiting
defect.
[0111] In certain circumstances, the presence of impurity atoms has
the effect of stabilizing a crystal structure or stabilizing an
amorphous structure under conditions in which such a structure
would not normally be stable. In other circumstances, this
stabilizing effect manifests itself by ameliorating the growth of
the grains making up the reinforcing material which can occur when
a crystalline material is maintained at elevated temperatures,
specifically, temperatures which are at a substantial fraction of
the material's melting point. Such grain growth typically has a
deleterious effect on the strength of the fiber reinforcement and
thus the component material itself because the size of
strength-limiting flaws is often proportional to the grain size of
a material. Thus, by stabilizing the as-fabricated small grain size
of a reinforcing material at elevated temperature, strength losses
may be reduced or eliminated. Impurity materials (which are
typically located at grain boundaries) may serve to "pin" the grain
boundaries of the grains making up the reinforcing material, thus
stabilizing the reinforcing material against grain growth, and thus
strength loss, at elevated temperatures. In such circumstances, it
is therefore desirable to maintain the presence of such impurity
materials within a reinforcing fiber. Furthermore, because some of
these impurities may tend to volatilize out of certain reinforcing
fibers at elevated temperatures, it may be desirable to have in
place around each reinforcing fiber a coating which may serve to
prevent the stabilizing impurity materials from such
volatilizing.
[0112] In the case of a stabilized silicon carbide fiber such as
NICALON.RTM. fiber, a boron nitride coating has been observed to
help maintain the oxygen and/or nitrogen stabilizing impurity atoms
within the fiber during elevated temperature exposure of the fiber.
Without wishing to be bound by any particular theory or
explanation, it has been hypothesized that the nitrogen component
in the boron nitride coating effectively establishes, at elevated
temperature, a localized, fiber-external nitrogen atmosphere or
nitrogen partial pressure. This nitrogen atmosphere or nitrogen
partial pressure represents a steep concentration gradient of
nitrogen across the fiber/coating interface. This concentration
gradient is biased against diffusion of the nitrogen impurity out
of the fiber and thereby tends to maintain the nitrogen stabilizing
impurity within the fiber. Moreover, oxygen in the NICALON.RTM.
fiber, if such a fiber is left unprotected, would likewise diffuse
out of the fiber, similarly resulting in a degration of the fiber's
strength. With an adjacent boron nitride coating, however, the
diffusing oxygen contacts the boron nitride and reacts to form a
very thin (e.g., nanometers thick) coating of boron oxide at the
fiber/boron nitride interface. This thin coating of boria thereby
appears to inhibit further diffusion of oxygen from the fiber to
the external environment for the same reason that the boron nitride
coating suppresses nitrogen diffusion from the fiber.
[0113] Accordingly, it may be possible to protect similarly other
non-oxide ceramic reinforcing fibers from the effects of prolonged
exposure at elevated temperatures by coating the fibers with
coating materials having at least one element in common with either
a constituent of the basic fiber material or a necessary impurity
material located within the basic fiber material. Further, it may
be possible through application of the above-described concepts to
improve the high temperature stability of fibers based upon oxide
systems. For example, a fiber reinforcing material may comprise
grains of aluminum oxide and small amounts of one or more impurity
substances located at aluminum oxide grain boundaries to stabilize
the aluminum oxide grains against elevated temperature grain
growth. If the stabilizing impurity material tends to volatilize at
these temperatures, a coating comprising at least one elemental
component of the impurity material which is in sufficient proximity
to such a reinforcing material may help to maintain the stabilizing
impurity material within the aluminum oxide fiber.
[0114] The amount or thickness of coating material applied to the
filler material is also of importance, especially in composite
materials whose matrices are formed by infiltration. Coatings which
are too thick may tend to hinder infiltration of matrix material by
sealing or isolating regions of the permeable mass or preform to be
infiltrated, particularly those regions toward the center of the
body. Conversely a debond coating which is too thin may not provide
sufficient debonding or pull-out of the filler material
reinforcement from the matrix. Similarly, a protective coating
which is too thin may not be sufficiently protective.
[0115] It has been discovered for the case of coating reinforcement
materials such as fibers and, in particular, fibers comprising
silicon carbide and occupying about 35-38 percent of the bulk
volume of a preform, the thickness of the boron nitride coating
which optimizes both the ambient and elevated temperature flexural
strength of the resulting aluminum oxide matrix composite body is
between about 0.2 micron and about 0.5 micron and preferably
averages about 0.3 micron. The flexural strength of the composite
has been observed to decrease for boron nitride coating thicknesses
below about 0.2 micron due to stressing of the composite body
beyond its yield point (i.e., proportional limit). In other words,
as the composite body is stressed, an insufficient number of fibers
pull out of the matrix to relieve the increasing elastic strain
energy. Likewise, the flexural strength has been observed to
decrease in some composite bodies where the boron nitride coating
thickness exceeds about 0.5 micron due to the onset of a new
failure mode, specifically, for example, that of interlaminar
shear. Thus, with regard to mechanical strength, there does not
appear to be any benefit to applying boron nitride coatings thicker
than about 0.5 micron.
[0116] There also exists an optimal range for the silicon carbide
overcoat thickness.
[0117] Specifically, for the above-described system comprising
boron nitride and silicon carbide coated onto fabric plies of woven
continuous fibers comprising ceramic grade NICALON.RTM. silicon
carbide and occupying about 35-38 percent of the bulk volume of a
preform, the nominal thickness of the silicon carbide coating for
optimizing the flexural strength of the alumina matrix composite
formed by directed metal oxidation has been found to be in the
range from about 2.0 to 2.3 microns. More particularly, nominal
silicon carbide coating thicknesses thinner than about 1.75 microns
yielded composites with fracture strengths which were significantly
below the fracture strength of composites having coatings of the
thicknesses discussed above. Without wishing to be bound by any
particular theory or explanation, this loss of strength may result
from the relatively thin silicon carbide coatings inadequately
protecting the underlying boron nitride and/or silicon carbide
fiber reinforcement materials from chemical attack. Such chemical
attack of the reinforcement materials may occur during the
formation of the matrix phase of the composite and/or during
subsequent exposure of the formed composite to undesirable
oxidant(s) at elevated temperatures. Likewise, nominal silicon
carbide coating thicknesses greater than about 2.3 microns have
also yielded flexural strength losses. Again, without wishing to be
bound by any particular theory or explanation, silicon carbide
coatings having nominal thicknesses which are greater than about
2.3 microns appear to "can" or seal-off the space within and/or
between the fiber plies. This "canning" could then result in the
creation of closed porosity which may prevent subsequent
infiltration of oxidation reaction product into such closed
porosity during the directed metal oxidation process, thereby
yielding weakly bonded fiber plies, and thus mechanically weakened,
composite body.
[0118] It should be understood that the fiber coating thicknesses
discussed herein have been calculated from the total weight gains
which the fiber preforms experience during the fiber coating
process which in most cases herein refers to a chemical vapor
infiltration (CVI) coating process. For boron nitride, the "actual"
coating thickness as measured from photomicrographs of fiber
cross-sections agree well with the calculated values, as each fiber
is coated with a relatively thin layer of boron nitride. For the
thicker silicon carbide coating however, the two values diverge.
The subsequent coating of silicon carbide, however, is relatively
much thicker, and as the silicon carbide coatings build up in
thickness, they may come into contact with one another,
particularly where individual filaments are in close proximity, as
is shown in FIG. 14a, for example. This merging of individual
silicon carbide coatings has the effect of isolating some parts of
the filaments from further coating deposition. Hence, the actual
thickness of a silicon carbide coating on those portions of a fiber
where such a coating actually exists is somewhat larger than the
"nominal" silicon carbide thickness calculated from weight gain
values. Further, as a result of the nature of the CVI coating
process, the silicon carbide coating thickness is somewhat greater
for fibers in the outer preform plies tan in the inner plies. Thus,
for the above-described preform system comprising about 35-38
volume percent of approximately 15 to 20 micron diameter fibers, a
nominal or calculated silicon carbide thickness of about 2.3
microns corresponds to an actual coating thickness of about 4-6
microns near the exterior of the preform. Similarly, a nominal
thickness of about 1.5 microns corresponds to an actual coating
thickness of about 2 microns near the preform exterior.
[0119] As mentioned immediately above, the nature of the CVD or CVI
coating process is to deposit a thicker-than-average layer on the
exterior regions of the fiber preform and a thinner-than-average
layer in those zones more towards the center of the preform. For
example, depending on the permeability or "openess" of the weave
and the thickness of the preform, an "average" SiC coating
thickness of 2-2.3 microns, based upon weight gain values may
correspond to actual depositions (as measured by microscopy)
ranging from about 4-6 microns at the preform exterior to only
about 0.5-1.0 micron at the very center of the preform.
[0120] A technique has been found, however, to at least partially
amelioate this effect. Specifically, and as discussed elsewhere,
when the preform is assembled as a stack of fabric plies, utilizing
plies having a more "open" weave as the plies at the closest to the
exterior of the preform provides the CVI or CVD reactant gases
greater access to the interior regions of the preform. For example,
a fiber preform may be assembled using "eight harness satin weave"
(8 HSW) and 12 HSW fabric plies. Because the 12 HSW plies exhibit a
"tighter" weave than do the 8HSW plies, to make the coating
thickness as uniform as possible through the preform, the 8 HSW
plies should be placed toward the exterior of the preform and the
12 HSW plies should be placed at the center. An artisan of ordinary
skill will appreciate the many possible combinations among 12 HSW,
8 HSW and plain weave fabrics and even three-dimensionally woven
fiber preforms. For example, a 3-D woven fiber preform may be
sandwiched in between at least one pair of plain weave fabrics
having a lower volumetric loading of reinforcement material. In
general, and for all other factors being equal, the permeability or
"openness" of woven fabric increases in the following order: 12
HSW, 8HSW, 5 HSW, plain weave. The permeability of 3-D woven
preforms depends upon too many other variables to allow a
generalization to be made;
[0121] however, for the same substance (e.g., NICALON.RTM. fiber),
the permeability of any given 3-D woven preform may be estimated
vis a vis that of a 2-D fabric weave by comparing the bulk
densities of the respective fiber forms, in other words, by
comparing the respective total porosities.
[0122] In yet another aspect of the invention, a technique for
reducing the number and/or size of microcracks formed during
composite formation and/or formed during composite service or use,
may be provided. Microcracks may be undesirable because such
microcracks may permit ready access of undesirable oxidants to the
reinforcement material(s) which can result in degradation of some
properties of the composite body. Specifically, microcracking of a
matrix material located between adjacent plies of fiber
tows/bundles (e.g., silicon carbide fiber) can be reduced or
possibly even eliminated by introducing into the matrix one or more
materials having a relatively low coefficient of thermal expansion
(e.g., lower than that of the matrix material) such as, for
example, silicon carbide particulate. Thus, to practice this
embodiment of the invention, an appropriate material or combination
of materials could be inserted between one or more fiber tows or
between fiber layers to form a preform from a combination of fibers
and particulate. After formation of the preform, a ceramic matrix
comprising, for example, an oxidation reaction product could be
formed.
[0123] When plies or sheets of woven silicon carbide fiber tows are
utilized in conjunction with an aluminum oxide matrix and more
particularly where the warp yarns of adjacent plies are oriented at
ninety degrees to one another, microcracks in the matrix may
result. In this right angle orientation especially, there are
(inevitable) regions between adjacent fiber plies substantially
unoccupied by reinforcement fibers. During the directed metal
oxidation process, these regions as well as any void spaces between
individual fibers within a fiber tow and between individual fiber
tows within a fiber ply, fill in with ceramic oxidation reaction
product. It has been observed that the ceramic matrix material
between adjacent fiber plies may be particularly susceptible to
microcracking. FIG. 17A, for example, is an approximately
50.times.magnification optical photomicrograph of a polished
cross-section of such a fiber reinforced ceramic matrix composite
which shows several such cracks 300 within the ceramic matrix
material 302 located between adjacent plies of woven tows of the
reinforcement fibers 304.
[0124] Without wishing to be bound by any particular theory or
explanation, microcracking of the matrix material, particularly,
those portions of the matrix occupying the space between adjacent
fiber plies, may result from a difference or mismatch in the local
thermal expansion coefficient of the composite material,
specifically between the region of the composite within and between
fiber plies, respectively. Accordingly, the observed matrix
microcracking may occur at some point during the cooling of the
infiltrated preform from the process temperatures to ambient
temperature. For example, the NICALON.RTM. silicon carbide fibers
employed in fabricating the fiber plies making up the preform have
a thermal expansion coefficient of about 4 ppm/.degree. C., whereas
the aluminum oxide and aluminum alloy making up a typical matrix
material of a composite body formed by directed metal oxidation
have thermal expansion coefficients of about 8 and 23,
respectively.
[0125] Several concepts have been advanced in an effort to
ameliorate the adverse consequences of these matrix microcracks. In
particular, it has been hypothesized that thermal expansion
mismatch stresses might be reduced if the composite body were made
more isotropic by, for example, reducing the difference in fiber
orientation between adjacent plies from ninety degrees to thirty or
forty-five degrees. Another idea has been to reduce the (largely
unoccupied) space between adjacent fiber plies in the preform by
using thinner plies or by clamping the assemblage of plies more
tightly together during chemical vapor infiltration.
[0126] In contrast to these concepts directed to reducing the
amount of space between adjacent plies is the concept of filling
this space with another material (e.g., a filler material) whose
thermal expansion coefficient is selected such that the local
thermal expansion coefficient of the composite material between
adjacent fiber plies is closer (in value) to that of the composite
material within the fiber plies after such plies have been embedded
with the ceramic matrix. According to this reasoning and in view of
the thermal expansion coefficient of the fiber reinforcement being
lower than that of the matrix, an appropriate filler material for
the space between the plies could include a filler which has a
thermal expansion coefficient lower than that of the matrix. One
such low thermal expansion coefficient filler material which has
been shown to be effective in this regard is silicon carbide. Not
only does adding silicon carbide particulate between the fabric
plies reduce the thermal expansion coefficient of the otherwise
unreinforced alumina matrix material by virtue of the rule of
mixtures, but additionally the lower expansion silicon carbide
bodies act to constrain the contraction (upon cooling from the
processing conditions) of the higher expansion alumina matrix.
Although the morphology of the added silicon carbide which has been
successfully employed (as discussed later herein) was in the form
of particulates, other forms such as platelets, whiskers or chopped
fibers could also be expected to work effectively.
[0127] Further, for many of the applications contemplated for the
fiber reinforced composite materials of the present invention, high
thermal conductivity of the composite is a desirable attribute. The
presence of flaws in a material such as cracks tends to reduce the
thermal conductivity of the material. Thus, a reduction in the
number and/or size of microcracks in a composite body also has the
desirable affect of increasing thermal conductivity. Moreover, it
is possible to further increase the thermal conductivity of the
body through selection of materials having relatively high thermal
conductivity and a low thermal expansion coefficient. In this
regard, the selection of silicon carbide for placement as a filler
material between plies of silicon carbide fiber tows is also an
excellent choice because of the relatively high thermal
conductivity of silicon carbide.
[0128] The filler material may be introduced between the fiber
plies either before or after the CVI coating of the fiber plies. In
a preferred embodiment, however, the filler materials which are
introduced to the fiber plies are, at some point, CVI coated. In a
particularly preferred embodiment, the filler materials are
introduced between the fiber plies prior to CVI coating by means
of, for example, vibration or slurry infiltration. Thus, the plies
are assembled into a preform in the presence of such filler
materials and the resulting preform assembly is then subsequently
CVI coated such that both the fibers comprising the tows, as well
as the filler material between adjacent plies of fiber tows are
simultaneously coated.
[0129] Because increasing the thermal conductivity of a ceramic
body reduces its susceptibility to cracking due to thermal shock,
one may choose to approach the matrix microcracking problem by
selecting filler materials for placement between the fiber plies,
not necessarily based upon low thermal expansion coefficient, but
based upon high thermal conducivity. Accordingly, even filler
materias having relative high thermal expansion coefficients such
as, for example, TiB2, may be gainfully employed in this way to
reduce that matrix microcracking which is due to thermal shock,
provided that such filler materials possess high thermal
conductivity.
[0130] To further expand on this fiber reinforcement embodiment,
one or more substantially non-reactive filler materials different
from the fibrous filler material may be added among the fibers
making up a preform (e.g., a 3-D woven fiber preform) or between
the layers of fabric making up a preform. Provided that sufficient
fiber loading remains to accomplish their purpose (e.g., composite
toughening), a wide variety of other filler materials may be added
to tailor a host of desired properties, for example, bulk density,
thermal conductivity, wear resistance, ballistic performance,
etc.
[0131] For example, once the required toughness has been met by
providing a certain volumetric loading of de-bondable fibers, one
or more other properties may be tailored through the addition of a
different filler material. For example, it may be desirable to
improve the hardness of the composite material through addition of
hard filler materials such as silicon carbide, boron carbide, the
transition metal carbides, titanium diboride and/or boron
carbide.
[0132] One convenient method for incorporating such additional
filler materials is by means of slurry infiltration or
impregnation. Specifically, the different filler material may be
provided in whisker or particulate form and sipersed in water, an
organic solvent or a preceramic polymer such as CERASET=SN
inorganic polymer (Lanxide Performance Materials, Newark, Del.) to
make a slurry. The slurry could then be painted, sprayed or poured
onto fabric plies of the fibrous reinforcement. Alternatively, the
fabric plies or 3-D fibrous preforms could be dipped into the
slurry. Optionally, pressure or vacuum could be administered to
assist the infiltration of the slurry into the void space between
the fibers. The slurry infiltrated fibrous preform or fabric plies
(once assembled into the desired shape of the final self-supporting
body) would then be heated to a modest temperature sufficient to
remove volatiles or cure any polymeric components.
[0133] It should be understood that while this disclosure relates
primarily to matrices which are formed by directed metal oxidation,
the concepts disclosed should be applicable to other matrix/fiber
combinations. Accordingly, both reduction of microcracking and the
increasing of thermal conductivity can be enhanced in other systems
as well.
[0134] The first and outer coatings, typically, are deposited onto
the ceramic filler material by conventional or known means such as
chemical vapor deposition, plasma spraying, physical vapor
deposition, plating techniques, sputtering or sol-gel processing.
Achievement of a substantially uniform coating system according to
these prior art techniques is within the level of skill in this
art. For example, chemical vapor deposition of a uniform coating of
boron nitride on ceramic filler materials can be achieved by using
boron trifluoride and ammonia at a temperature of about
1000-1500.degree. C. and a reduced pressure of 1-100 torr; boron
trichloride and ammonia at a temperature of 600-1200.degree. C. and
reduced pressure of 0.1-100 torr; borazine at a temperature of
300-650.degree. C. and a reduced pressure of 0.1-1 torr; or
diborane and ammonia at a temperature of 600-1250.degree. C. and a
reduced pressure of 0.1-1 torr. A coating of silicon carbide by
chemical vapor deposition can be accomplished, for example, by
using methyltrichlorosilane at a temperature of 800-1500.degree. C.
and a pressure of 1-760 torr; dimethyldichlorosilane at a
temperature of 600-1300.degree. C. and a reduced pressure of 1-100
torr; and silicon tetrachloride and methane at a temperature of
900-1400.degree. C. and a reduced pressure of 1-100 torr.
[0135] It should be understood that various combinations of ceramic
materials having one or more coatings may be produced depending on
the specific properties desired in the coated ceramic material and
its ultimate application. A possible combination includes silicon
carbide fiber with a first layer of titanium carbide and an
additional outer layer of silicon nitride. Another coating system
includes silicon carbide fiber with a first coating of boron
nitride and additional outer coatings of silicon carbide and
alumina.
[0136] In the manufacture of ceramic matrix composites according to
the directed metal oxidation embodiment of the invention, the
coated materials may be provided in the form of a loose mass or may
be laid up into a porous preform of any desired configuration. The
parent metal is placed adjacent the preform. The parent metal is
then heated in the presence of an oxidant to above its melting
point whereby the molten metal oxidizes to form and develop an
oxidation reaction product embedding the coated ceramic material.
During growth of the oxidation reaction product, the molten parent
metal is transported through its own otherwise impervious oxidation
reaction product, thus exposing free metal to the oxidizing
atmosphere to yield additional reaction product. The result of this
process is the progressive growth of an interconnected ceramic
oxidation reaction product which optionally may contain nonoxidized
parent metal.
[0137] A variety of ceramic matrices may be produced by the
oxidation reaction of parent metals depending upon the choice of
parent metal and oxidant. For example, ceramic matrices may include
oxides, nitrides, borides, or carbides of such parent or infiltrant
metals as aluminum, silicon, titanium, tin, zirconium or hafnium.
The ceramic matrix composites of the invention may comprise, by
volume, 5 to 85% of the coated ceramic filler materials and 95 to
15% of ceramic matrix. A useful composite comprises an alumina
matrix formed by the oxidation reaction of aluminum parent metal in
air, or an aluminum nitride matrix by oxidation reaction (i.e.,
nitridation) of aluminum in nitrogen, and incorporating as a
reinforcing filler such materials as alumina, silicon carbide,
silicon nitride, etc., bearing the coating system.
[0138] The choice of parent metal and oxidant will determine the
composition of the polycrystalline matrix, as explained in the
Commonly Owned Patents and Patent Applications. Thus a filler
bering the coating system may have admixed therewith a solid or
liquid oxidant, such as boron, silica, or glasses (e.g., low
melting glasses), or the oxidant may be gaseous, such as an
oxygen-containing gas (e.g. air) or a nitrogen-containing gas (e.g.
forming gas typically comprising, by volume, 96% nitrogen and 4%
hydrogen).
[0139] Another useful composite material system is that of melt
infiltration. Here, a silicon-based metal is melted and contacted
to a permeable mass. The permeable mass comprises a material which
can be wetted by molten silicon, such as silicon carbide. Under
wetting conditions, molten silicon-containing metal can infiltrate
such a permeable mass in a pressureless manner. The infiltration
typically is conducted under an inert atmosphere such as argon, or
in a vacuum. The permeable mass optionally may include a carbon
source, typically graphite, which may react with the infiltrating
silicon to form silicon carbide in the matrix. Depending upon the
amount of carbon source and the degree of reaction, the in-situ
formed silicon carbide may be interconnected or discrete,
discontinuous bodies.
[0140] The oxidation protection mechanisms of the present invention
can also be applied to composite systems whose matrices may be
formed by chemical vapor infiltration (e.g., CVI SiC) or by
repeated infiltration and pyrolysis of ceramic precursor polymers
such as the polysilazanes.
[0141] It should be understood that while this disclosure relates
primarily to matrices which are formed by directed metal oxidation,
the concepts disclosed should be applicable to other matrix
processing systems such as sintering, hot pressing or other
infiltration techniques employed to produce glass (e.g., "Black
Glass" or "Comp Glass"), metal, polymer or other ceramic
matrices.
[0142] The following examples illustrate certain aspects and
advantages of various embodiments of the invention.
EXAMPLE 1
[0143] Two fiber-reinforced alumina-matrix ceramic composite bodies
were fabricated in accordance with the present invention. The
fibers employed were NICALON.RTM. ceramic grade silicon carbide as
Si--C--O--N (from Nippon Carbon Co., Ltd., Japan) measuring
approximately 2 inches long and approximately 10-20 lm in diameter.
Each fiber was coated via chemical vapor deposition with a duplex
coating. The duplex coating comprised a 0.2-0.5 lm thick first
coating of boron nitride applied directly to the fiber, and a
1.5-2.0 lm thick second (outer) coating of silicon carbide applied
to the boron nitride coating.
[0144] The duplex coated fibers were gathered into bundles, each
containing 500 fibers tied with a single fiber tow. Two, 2 inch
square by 1/2 inch thick bars of aluminum alloy designated 380.1
(from Belmont Metals, having a nominally identified composition by
weight of 8-8.5% Si, 2-3% Zn, and 0.1% Mg as active dopants, and
3.5% Cu, as well as Fe, Mn, and Ni, but the actual Mg content was
sometimes higher as in the range of 0.17-0.18%) were placed into a
bed of Wollastonite (a mineral calcium silicate, FP grade, from
Nyco, Inc.) contained in a refractory crucible such that a 2 inch
square face of each bar was exposed to the atmosphere and
substantially flush with the bed, while the remainder of each bar
was submerged beneath the surface of the bed. A thin layer of
silica sand was dispersed over the exposed surface of each bar to
serve as an additional dopant. Three of the above-described bundles
of duplex-coated fibers were placed on top of each of the two
sand-layered metal surfaces, and these set-ups were covered with
Wollastonite.
[0145] The crucible with its contents was placed in a furnace which
was supplied with oxygen at a flow rate of 500 cc/min. The furnace
temperature was raised to 1000.degree. C. at a rate of 200.degree.
C./hour, and held at 1000.degree. C. for 54 hours.
[0146] The crucible was then removed while the furnace temperature
was at 1000.degree. C., and allowed to cool to room temperature.
The ceramic composite products were recovered. Examination of the
two ceramic composite products showed that an alumina ceramic
matrix, resulting from oxidation of aluminum, had infiltrated and
embedded the fiber bundles.
[0147] Two specimens were machined from each of the two ceramic
composite products. FIGS. 1 and 2 are scanning electron micrographs
at about 350.times.magnification and about 850.times.magnification,
respectively, showing this ceramic matrix composite. Referring to
the micrographs, there is shown the alumina matrix 2 incorporating
silicon carbide fibers 4 bearing a first inner coating 6 of boron
nitride and an outer coating 8 of silicon carbide. One machined
specimen from each composite product was tested for flexural
strength (Sintec strength testing machine, Model CITS 2000/6, from
Systems Integrated Technology Inc., Stoughton, Mass.) in 4 point
bend with a 12.67 mm upper span and a 28.55 mm lower span. The
values obtained were 448 and 279 MPa. The remaining specimen from
each product was tested for Chevron notch fracture toughness, and
the values obtained were 19 and 17 MPa-m.sup.1/2, respectively.
FIG. 3 is a scanning electron micrograph at 250.times.magnification
of the fractured surface of the ceramic composite showing extensive
pull-out of the fibers.
[0148] This nun was repeated with the exception that the
NICALON.RTM. fibers were not coated. FIG. 4 is a scanning electron
micrograph at 800.times.magnification of the fractured surface
showing essentially no pull-out of the fibers. Typical values for
strength ranged from 100-230 MPa, and for toughness ranged from 5-6
MPa-m.sup.1/2.
[0149] The utility of coated filler material made according to the
invention is clearly demonstrated by this Example and the
comparative data.
EXAMPLE 2
[0150] The following Example demonstrates a method for forming a
fiber reinforced ceramic composite body, and illustrates the
resultant mechanical properties of the body from about room
temperature to about 1400.degree. C. Specifically, this Example
demonstrates a method for forming a silicon carbide fiber
reinforced alumina composite body wherein the silicon carbide
fibers are coated with a first layer of boron nitride and a second
layer of silicon carbide to create a debond zone between the
silicon carbide fiber and the alumina matrix.
[0151] A fabric preform 103 was made by stacking a plurality of
layers of 8 harness satin weave (8 HSW) fabric and 12 harness satin
weave (12 HSW) fabric made from ceramic grade NICALON.RTM. silicon
carbide fiber (obtained from Dow Corning Corporation, Midland,
Mich.) on top of each other. FIGS. 5a and 5b are schematics
depicting a top view and a cross-sectional view respectively of the
as-is position for a HSW fabric. In reference to FIG. 5a and 5b, a
HSW fabric is designated to be in the "as-is position" when, as
viewed in cross-section, the axes of the warp yarns 92 of the
fabric 90 are in the plane of the cross-sectional view and are
located at the bottom (i.e., as shown in the cross-sectional view)
of the fabric 90 and the axes of the fill yarns 91 are
perpendicular to the plane of the cross-sectional view and are
located at the top of the fabric 90. The orientation of additional
fabric layers can be described in reference to the as-is position.
For example, as depicted in FIG. 5c, additional fabric layers can
be (1) rotated about an axis 93 perpendicular to the plane of the
fabric 90 and/or (2) rotated about an axis 94 perpendicular to the
plane of the cross-section of the fabric 90 and then subsequently
contacted or layered upon a fiber layer positioned in the as-is
configuration. Thus, for example, as schematically depicted in
cross-section in FIG. 5d, a substantially square fabric preform 103
can be made from 8 pieces of HSW fabric, stacked in the following
sequence:
[0152] A first fabric layer 95 comprising an 8 HSW fabric was
placed on a supporting surface in the as-is position to start the
fabric preform 103;
[0153] A second fabric layer 96 comprising a 12 HSW fabric, was
rotated about 90.degree. in the counterclockwise direction from the
as-is position about an axis 93 perpendicular to the plane of the
fabric and was placed on the first fabric layer 95 so that the
edges of the second fabric layer 96 were substantially aligned with
the edges of the first fabric layer 95;
[0154] A third fabric layer 97 comprising a 12 HSW fabric, in the
as-is position, was placed on the second fabric layer 96 so the
edges of the third fabric layer 97 were substantially aligned with
the edges of the second fabric layer 96;
[0155] A fourth fabric layer 98 comprising a 12 HSW fabric, was
rotated about 90.degree. in the counterclockwise direction from the
as-is position about an axis 93 perpendicular to the plane of the
fabric and was placed on the third fabric layer 97 so that the
edges of the fourth fabric layer 98 were substantially aligned with
the edges of the third fabric layer 97;
[0156] A fifth fabric layer 99 comprising a 12 HSW fabric, was
rotated about 90.degree. in the counterclockwise direction from the
as-is position about an axis 93 perpendicular to the plane of the
fabric and then rotated about 180.degree. in the clockwise
direction about an axis 94 perpendicular to the plane of the
cross-sectional view of the fabric and was placed on the fourth
fabric layer 98 so that the edges of the fifth fabric layer 99
substantially aligned with the edges of the fourth fabric layer
98;
[0157] A sixth fabric layer 100 comprising a 12 HSW fabric, was
rotated about 180.degree. in the clockwise direction from the as-is
position about an axis 94 perpendicular to the plane of the
cross-sectional view of the fabric and was placed on the fifth
fabric layer 99 so that the edges of the sixth fabric layer 100
were substantially aligned with the edges of the fifth fabric layer
99;
[0158] A seventh fabric layer 101 comprising a 12 HSW fabric, was
rotated about 90.degree. in the counterclockwise direction from the
as-is position about an axis 93 perpendicular to the plane of the
fabric and then rotated about 180.degree. in the clockwise
direction about an axis 94 perpendicular to the plane of the
cross-sectional view of the fabric and was placed on the sixth
fabric layer 100 so that the edges of the seventh fabric layer 101
were substantially aligned with the edges of the sixth fabric layer
100; and
[0159] Finally, an eighth fabric layer 102 comprising an 8 HSW
fabric, was rotated about 180.degree. in the clockwise direction
from the as-is position about an axis perpendicular 94 to the plane
of the cross-sectional view of the fabric and was placed on the
seventh fabric layer 101 so that the edges of the eighth fabric
layer 102 were substantially aligned with the edges of the seventh
fabric layer.
[0160] In reference to FIG. 5e, the fabric preform 103 comprising
two 8 HSW outer fabric layers and six 12 HSW inner fabric layers
and measuring about 6.75 inch (171 mm) square and about 0.125 inch
(3.2 mm) thick was placed on a perforated graphite plate 104
machined from Grade AXF-5Q graphite (Poco Graphite, Inc., Decatur,
Tex.) which measured about 7.75 inches (197 mm) square and about
0.5 inch (13 mm) thick. The inner perforated region 105 of the
perforated plate measured about 6.25 inches (159 mm) square. The
holes 106 of the perforated region 105 had a diameter of about 0.25
inch (6.4 mm) and a center-to-center spacing of about 0.375 inch
(9.5 mm) and comprised a 17 hole.times.17 hole array which was
bordered by an about 1 inch (25 mm) unperforated region. After the
fabric preform 103 had been placed on the first graphite plate 104,
a second graphite plate 104, substantially the same as the first,
was placed over the fabric preform 103 and the plates were clamped
using C-clamps to compress the fabric preform 103. Two graphite
channel members 107 machined from Grade AXF-5Q graphite (Poco
Graphite, Inc., Decatur, Tex.) and measuring about 7.75 inches (197
mm) long were placed over common ends of both perforated graphite
plates 104 so as to contact opposite ends of the first and second
perforated graphite plates 104 thereby creating a preform
containment fixture 108. FIG. 5e is an isometric schematic view of
the preform containment fixture 108. After the graphite channels
107 were secured to the perforated plates 104, the C-clamps were
removed from the perforated plates 104 and the elastic force
exerted by the compressed fabric preform 103 biased the perforated
graphite plates 104 against the graphite channel members 107 to
form a relatively rigid preform containment fixture 108. The warp
yarns 92 of the eighth layer 102 of the fabric preform 103 within
the graphite containment fixture 108 were positioned so as to be
parallel to the length of the graphite channel members 107 of the
preform containment fixture 108.
[0161] The graphite containment fixture 108 containing the fabric
preform 103 was placed into a reactor chamber of a chemical vapor
infiltration apparatus having an outer diameter of about 12 inches
(305 mm). The inner diameter of the reactor chamber measured about
9.45 inches (240 mm) after being lined with a quartz tube having a
wall thickness of about 0.5 inch (13 mm) and lined with a graphite
tube having a wall thickness of about 0.25 inch (6.4 mm). The warp
yarns 92 of the eighth layer 102 of the fabric preform 103 were
parallel to the gas flow direction within the chamber as well as
being parallel to the longitudinal axis of the reactor chamber. The
reactor chamber was closed and evacuated to about 0.004 inch (0.1
mm) of mercury (Hg). Then the reactor chamber was heated to about
800.degree. C. at about 10.degree. C. per minute so that the
contents of the reactor chamber were at about 730.degree. C., as
indicated by a thermocouple contained therein. When the temperature
within the reactor chamber reached about 730.degree. C., a gas
mixture comprised of ammonia (NH.sub.3) flowing at about 1200
standard cubic centimeters per minute (sccm) and boron chloride
(BCl.sub.3) flowing at about 800 sccm was introduced into the
reactor chamber while maintaining a total operating pressure of
from about 0.047 to about 0.051 inches of mercury (about 1.2 to
about 1.3 mm Hg). After about 6.5 hours at about 730.degree. C.,
the gas mixture flowing into the reactor chamber was interrupted,
the power to the furnace heating the reactor chamber was
interrupted, and the furnace and its contents were naturally cooled
to about 200.degree. C. At about 200.degree. C., the reactor
chamber door was opened and the graphite containment fixture 108
was removed, cooled and disassembled to reveal that the fibers of
the fabric layers of the fabric preform 103 were coated and that
the fabric layers comprising the fabric preform 103 were bonded
together by a boron nitride coating formed during the process at
about 730.degree. C., thereby forming a coated and bonded fabric
preform 109. The boron nitride coating had a thickness of about 0.4
microns.
[0162] The boron nitride coated and bonded fabric preform 109 was
then suspended from a graphite cantilever support fixture 110 made
from Grade AXF-5Q graphite (Poco Graphite, Inc., Decatur, Tex.) by
wires 111 comprised of a Kanthal.backslash.iron-chromium-aluminum
alloy all of which are depicted schematically in FIG. 5f. The
graphite cantilever support fixture 110 and the boron nitride
bonded fabric preform 109 were then replaced into the reactor
chamber of the chemical vapor infiltration apparatus discussed
above such that the warp yarns 92 of the eighth layer 102 comprised
of the 8 harness satin weave fabric were parallel to the gas flow
direction within the chamber as well as being parallel to the
longitudinal axis of the reactor chamber. After the reactor chamber
door was closed, the reactor chamber and its contents were
evacuated to about 0.591 inches (15 mm Hg) and hydrogen gas flowing
at about 2500 sccm was introduced into the reactor chamber. The
reactor chamber was heated at about 10.degree. C. per minute so
that the contents of the reactor chamber were at about 925.degree.
C. as indicated by a thermocouple therein. When the reactor chamber
contents were at about 925.degree. C., additional hydrogen, flowing
at about 2500 sccm, was introduced into the reactor chamber to give
a total hydrogen gas flow rate of about 5000 sccm. Once the
temperature of the contents of the reactor chamber had
substantially completely stabilized at about 925.degree. C., about
2500 sccm hydrogen were diverted away from direct entry into the
reactor chamber, and were first bubbled through a bath of
trichloromethylsilane (CH.sub.3SiCl.sub.3) also known as
methyltrichlorolsilane (MTS) (Hulls/Petrarch System, Bristol, Pa.),
maintained at about 25.degree. C., before entering the reactor
chamber. After about 26 hours at about 925.degree. C., the power to
the furnace heating the reactor chamber was interrupted and the
about 2500 sccm hydrogen that was being directed through the MTS
bath was again permitted to flow directly into the reactor chamber
to re-establish a direct hydrogen gas flow rate of about 5000 sccm
into the reactor chamber. It was noted that about 4.75 liters of
MTS had been consumed during the 26 hour of the run at about
925.degree. C. After about a half hour during which a hydrogen gas
flow rate at about 5000 sccm was maintained, the hydrogen flow rate
was interrupted and the furnace and its contents were evacuated to
about 0.039 inches 0.1 mm of mercury (Hg). The pressure within the
reactor chamber was then allowed to increase to about atmospheric
pressure while argon was introduced at a flow rate of about 14
liters per minute. After the reaction chamber had cooled to a
temperature of about 200.degree. C., the argon flow rate was
interrupted and the reaction chamber door was opened. The graphite
cantilever support fixture 110 and the fabric preform were removed
from the reactor chamber to reveal that the boron nitride bonded
fabric preform 109 had been coated with a second layer of silicon
carbide thereby forming a silicon carbide (SiC)/boron nitride
(BN)-coated fabric preform 112. The silicon carbide had an overall
average thickness of about 2.3 microns, as calculated from the
weight gain of the preform during the silicon carbide coating
procedure, as alluded to previously.
[0163] A wax box pattern having a closed end and outer dimensions
of about 7 inches (178 mm) square by about 2 inches (51 mm) tall
and a wall thickness of about 0.25 inches (6.5 mm) was assembled
from high temperature wax sheet (Kit Collins Company, Cleveland,
Ohio) which contained adhesive backing on one side thereof. The wax
box pattern was assembled by using a hot wax knife. The closed end
of the wax pattern was beveled at an angle of about 22.degree.. A
slurry mixture comprised by weight of about 5 parts
BLUONIC.backslash.A colloidal alumina (Buntrock Industries, Lively,
Va.) and about 2 parts -325 mesh (average particle diameter less
than about 45 .mu.m) wollastonite (a calcium silicate mineral) was
made by hand mixing the materials together. The slurry mixture was
then painted onto the outer surface of the wax box pattern with a
one inch sponge brush and wollastonite powder (-10,+100 mesh)
having substantially all particles between about 150 and 2000
microns in diameter was sprinkled liberally onto the slurry mixture
coating to prevent runoff and to form a first precursor layer of a
shell 120. This procedure was repeated to build additional layers
of coating with an about 0.5 hour drying period between the
formation of the precursor layers. When enough precursor layers of
slurry mixture/coarse wollastonite were formed to produce a
thickness of about 0.25 inch (6.4 mm), the coated wax box pattern
was set aside to dry at about room temperature for about 24 hours.
The about 0.25 inch (6.4 mm) thick coating nominally comprised
about 12 slurry mixture/coarse wollastonite layers. After the
coated wax box pattern had substantially completely dried at about
room temperature, the wax box pattern was placed into an air
atmosphere furnace maintained under an exhaust hood and the furnace
and its contents were held at a temperature of about 120.degree. C.
for about 6 hours, during which time the wax melted leaving behind
an unfired precursor to an alumina bonded wollastonite shell 120.
The furnace and its contents were then heated to about 950.degree.
C. in about 2 hours and held at about 950.degree. for about 4 hours
to substantially completely remove any residual wax and ensure the
sintering of the alumina bonded wollastonite shell. The furnace and
its contents were then cooled to about room temperature.
[0164] About 40 grams of VASELINE.backslash.petroleum jelly vehicle
(Cheseborough Ponds, Inc., Greenwich, Conn.) were melted in a small
aluminum weighing dish on a hot plate set at about medium heat
until the jelly turned to a liquid. A clean sable brush was then
used to substantially completely coat one of the 6.75 inch (171 mm)
square surfaces of the SiC/BN-coated fabric preform 112 to provide
an interface for the application of a nickel oxide powder. A
mixture comprising about 8 grams of -325 mesh (particle diameter
less than about 45 .mu.m) nickel oxide powder and about 16 grams of
ethanol was applied with a sponge brush to substantially completely
cover the petroleum jelly coated surface of the SiC/BN-coated
fabric preform. After the ethanol had substantially completely
evaporated, the SiC/BN-coated fabric preform 112 was inserted into
the alumina bonded wollastonite shell 120 such that the uncoated
side of the SiC/BN-coated preform 112 not coated with the nickel
oxide powder contacted the bottom of the shell 120, as shown in
FIG. 5g. The spaces between the perimeter of the SiC/BN-coated
fabric preform 112 and the walls of the alumina bonded wollastonite
shell 120 were filled with coarse (-10, +100 mesh) wollastonite
until the surface of the wollastonite powder was substantially
flush with the nickel oxide powder-coated surface of the
SiC/BN-coated fabric preform 112. The alumina bonded wollastonite
shell 120 containing the SiC/BN-coated fabric preform 112 was then
placed onto stilts 122, which were made from fire brick, and was
thereafter surrounded by wollastonite powder 123 which was
contained in a refractory boat 124.
[0165] The SiC/BN-coated fabric preform 112 was then leveled. About
1600 grams of a parent metal was distributed into four 30 gram clay
crucibles (obtained from J.H. Berge, Inc., South Plainfield, N.J.)
in amounts of about 400 grams per crucible. The parent metal
comprised by weight of about 8.5 to 11.0 percent silicon, 3.0 to
4.0 percent copper, 2.7 to 3.5 percent zinc, 0.2 to 0.3 percent
magnesium, .ltoreq.0.01 percent calcium, .ltoreq.0.10 percent
titanium, 0.7 to 1.0 percent iron, .ltoreq.05 percent nickel,
.ltoreq.0.5 percent manganese, .ltoreq.0.35 percent tin,
.ltoreq.0.001 percent beryllium, .ltoreq.0.15 percent lead and the
balance aluminum. The refractory boat 124 and its contents, as well
as the four 30 gram clay crucibles containing the parent metal,
were placed into an air atmosphere furnace and the furnace door was
closed. The furnace and its contents were then heated from about
room temperature to about 700.degree. C. at about 400.degree. C.
per hour, during which time the VASELINE.backslash.petroleum jelly
volatilized and the nickel oxide powder 125 fell onto the surface
of the SiC/BN-coated fabric preform 112. After about an hour at
about 700.degree., during which time the parent metal 126 had
substantially completely melted, the parent metal 126 was then
poured into the alumina bonded wollastonite shell 120 and onto the
nickel oxide powder-coated side of the SiC/BN-coated fabric preform
112, thereby covering the surface of the preform 112. Wollastonite
powder 127 was then poured onto the surface of the molten parent
metal 126 within the alumina bonded wollastonite shell 120 to
substantially completely cover the surface of the molten parent
metal. This assembly formed the lay-up for growth of a ceramic
matrix composite body. The furnace and its contents comprising the
lay-up were then heated to about 950.degree. C. in about an hour.
After about 90 hours at about 950.degree. C., the furnace and its
contents were cooled to about 700.degree. C. in about 2 hours. At
about 700.degree. C., the lay-up was removed from the furnace and
residual molten parent metal was decanted from the alumina bonded
wollastonite shell 120, the shell 120 was quickly broken away from
the SiC/BN-coated fabric preform 112 and the preform 112 was buried
in a silica sand bed to cool the preform 112 to about room
temperature. At about room temperature, it was observed that an
oxidation reaction product had grown into and substantially
completely embedded the SiC/BN-coated fabric preform 112, thereby
forming a fiber reinforced ceramic composite body 130 having a
plurality of fabric layers comprised of harness satin weaves.
Specifically, the fiber reinforced ceramic composite body 130
comprised two outer layers of 8 harness satin weave silicon carbide
fabric and six inner layers of 12 harness satin weave silicon
carbide fabric embedded by an aluminum oxide oxidation product. The
composite body also comprised a metallic constituent comprising
residual unreacted parent metal.
[0166] Once the ceramic composite body had been manufactured, a
metal removal process was begun for the purpose of removing this
residual parent metal within the composite body. The first step of
the metal removal process was to form a filler material mixture for
infiltration by metal contained in the formed ceramic matrix
composite body.
[0167] Specifically, filler material mixture comprising by weight
of about 90 percent E67 1000 grit (average particle diameter of
about 5 .mu.m) alumina (Norton Co., Worcester, Mass.) and about 10
percent -325 mesh (particle diameter less than about 45 .mu.m)
magnesium powder (Reade Manufacturing Company, Lakehurst, N.J.) was
prepared in a one gallon NALGENE.RTM. wide mouth plastic container
(Nalge Co., Rochester, N.Y.). Alumina milling balls were added to
the filler material mixture in the plastic container and the
container lid was closed. The plastic container and its contents
were placed on a jar mill for about 4 hours to mix the alumina and
magnesium powders together. After the alumina mixing balls had been
separated from the alumina-magnesium filler material mixture 131,
the filler material mixture 131 was complete.
[0168] A stainless steel boat 132 measuring about 7 inches (179 mm)
square by about 2 inches (50.8 mm) deep and having a wall thickness
of about 0.063 inches (1.6 mm) was lined with a graphite foil box
133 made from a piece of GRAFOIL{cube root} graphite foil (Union
Carbide Corp., Carbon Products Division, Cleveland, Ohio). About 1
inch (25 mm) of the filler material mixture 131 was hand packed
into the bottom of the graphite foil lined stainless steel boat
132. The fiber reinforced ceramic composite body 130 was then
placed onto and forced into the filler material mixture 131.
Additional filler material mixture 131 was then poured over the
fiber reinforced ceramic composite body 130 to substantially
completely cover it. The filler material mixture 131 was then hand
packed to ensure good contact between the filler material mixture
131 and the fiber reinforced ceramic composite body 130, thereby
forming a metal removal lay-up as depicted schematically in
cross-section in FIG. 5h.
[0169] The metal removal lay-up comprising the stainless steel boat
132 and its contents was then placed into a resistance heated
controlled atmosphere furnace and the furnace chamber door was
closed. The furnace chamber and its contents were first evacuated
to at least 30 inches (762 mm) of mercury (Hg) vacuum, then the
vacuum pump was disconnected from the furnace chamber and nitrogen
was introduced into the chamber to establish about atmospheric
pressure of nitrogen in the chamber. This operation was repeated.
After the pressure in the furnace chamber reached about atmospheric
pressure, the furnace chamber and its contents were heated from
about room temperature to about 750.degree. C. at a rate of about
250.degree. C. per hour and held at about 750.degree. C. for about
5 hours and cooled from about 750.degree. C. to about 300.degree.
C. at about 200.degree. C. per hour with a nitrogen gas flow rate
of about 4000 sccm being maintained throughout the heating and
cooling. At about 300.degree. C., the nitrogen flow was
interrupted, the furnace door was opened, and the stainless steel
boat and its contents were removed and cooled by forced convection.
At about room temperature, the filler material 131 was separated
from the fiber reinforced ceramic composite body 130 and it was
noted that the metallic constituent of the fiber reinforced ceramic
composite body 130 had been substantially completely removed. The
fiber reinforced ceramic composite body 130 was then subjected to
grit blasting by a sand blaster which operated with a working
pressure of about 75 pounds per square inch to remove any excess
filler material that had adhered to the surface of the composite
body 130. The fiber reinforced ceramic composite body was then cut
with a diamond saw and machined into mechanical test specimens
measuring about 2.4 inches (60 mm) long by about 0.2 inch (6 mm)
wide by about 0.11 inch (3 mm) thick for mechanical properties
measurements, specifically flexeral strength testing.
[0170] Several of the machined mechanical test specimens were then
subjected to additional heat treatments. Except as otherwise noted,
these heat treatments were limited to the fiber reinforced ceramic
composite material of the present Example. Specifically, a first
group of samples was heat treated at about 1200.degree. C. for
about 24 hours and a second group of samples was heated treated at
about 1200.degree. C. for about 100 hours. The heat treatments were
effected by placing the mechanical test specimens onto alumina
trays with the tensile side of the test specimen facing away from
the alumina trays. The alumina trays and their contents were then
placed into air atmosphere furnaces and heated to about
1200.degree. C. at a rate of about 200.degree. C. per hour. After
about 24 hours at about 1200.degree. C., the furnace containing the
first group of samples was cooled to about room temperature at a
rate of about 200.degree. C. per hour, whereas after about 100
hours at about 1200.degree. C., the furnace containing a second
group of samples, was cooled to about room temperature at a rate of
about 200.degree. C. per hour.
[0171] The flexural strengths of the fiber reinforced ceramic
composite test specimens were measured using the procedure defined
by the Department of the Army's proposed MIL-STD-1942A (Nov. 21,
1983). This test was specifically designed for strength
measurements of high-performance ceramic materials. The flexural
strength is defined in this standard as the maximum outer fiber
stress at the time of failure. A four-point-1/4-point flexural test
was used. The height and width of the test bars were measured with
a precision of about 390 microinch (0.01 mm). The test bars were
subjected to a stress which was applied at four points by two lower
span bearing points and two upper span bearing points. The lower
span bearing points were about 1.6 inches (40 mm) apart, and the
upper span bearing points were about 0.79 inch (20 mm) apart. The
upper span was centered over the lower span, so that the load was
applied substantially symmetrically on the test bar. The flexural
strength measurements were made with a Sintec Model CITS-2000/6
universal testing machine (Systems Integrated Technology, Inc.,
Stoughton, Mass.). The crosshead speed during testing was about
0.02 inch per minute (0.55.degree. C., about 1300.degree. C. and
about 1400.degree. C. were performed with another universal testing
machine equipped with an air atmosphere resistance heated furnace
(Advanced Test Systems, Butler, Pa.).
[0172] Table I contains a summary of the four point flexural
strengths for NICALON.RTM. silicon carbide reinforced alumina
oxidation reaction product composite bodies. Specifically, Table I
summarizes the sample condition, the test temperature, the number
of samples tested, the average flexural strength and standard
deviation, the maximum flexural strength and the minimum flexural
strength. These data suggest that the flexural strength of fiber
reinforced ceramic composite bodies subjected to the methods of the
instant invention are substantially unaffected by test temperature
between about room temperature and about 1200.degree. C. Moreover,
these data suggest that the flexural strengths of fiber reinforced
ceramic composite bodies subjected to the methods of the instant
invention are only slightly degraded at test temperatures greater
than 1200.degree. C. and by extended exposure times at 1200.degree.
C.
EXAMPLE 3
[0173] This Example illustrates that fiber reinforced ceramic
composite bodies having varying ceramic matrix composition can be
formed. Specifically, Sample A of this Example comprised a silicon
carbide fiber reinforced alumina composite body; and Sample B of
this Example comprised a silicon carbide fiber reinforced aluminum
nitride composite body.
Sample A
[0174] A SiC/BN-coated fabric preform measuring about 3.0 inches
(76 mm) long by about 3.0 inches (76 mm) wide by about 0.125 inch
(3.2 mm) thick was prepared by stacking eight layers of 12-harness
satin weave (12 RSW) fabric comprising silicon carbide fibers
(ceramic grade NICALON.RTM. fibers obtained from Dow Coming
Corporation, Midland, Mich.) the fibers having a diameter ranging
from about 394 microinch (10 .mu.m) to about 787 microinch (20
.mu.m). The 12 HSW silicon carbide fabrics were stacked such that
each succeeding fabric layer was placed with its fill yarns being
rotated about 90.degree. with respect to the fill yarns of the
previous fabric layer. The fabric preform comprising the stacked
layers were then placed into a chemical-vapor-infiltration (CVI)
reactor and the fibers were coated with a first layer of boron
nitride (BN) substantially in accordance with the methods of
Example 2. Thereafter, the reaction conditions in the CVI reactor
were modified such that a CVI coating of silicon carbide (SiC) was
placed on top of the BN coating substantially in accordance with
the method of Example 2. The CVI coatings held the stacked fabric
layers together, thereby forming the SiC/BN-coated fabric
preform.
[0175] The SiC/BN-coated fabric preform comprising the eight
stacked layers of 12 HSW fabric coated with a first layer of BN and
a second layer of SiC was placed into the bottom of a porous
castable refractory boat having boles at the bottom to facilitate
air flow to the composite during composite growth, thereby forming
a lay-up. Specifically, the porous castable refractory boat having
an inner cavity measuring about 3.25 inches (83 mm) square by about
3.0 inches (76 mm) deep and having a wall thickness of about 0.125
inch (3.2 mm) was cast from a mixture comprised by weight of about
56.3% plaster of Paris (BONDEX-, Bondex International), about 28.1%
water and about 15.6% 90 grit alumina (E1 ALUNDUM.backslash.,
Norton Company, Worcester, Mass.). After the SiC/BN-coated fabric
preform was placed into the porous castable refractory boat, -325
mesh (particle diameter less than about 45 .mu.m) wollastonite
particulate (a calcium silicate obtained from Peltz-Rowley Chemical
Co., Philadelphia, Pa.) was placed into the void space between the
SiC/BN-coated fabric preform and the porous castable refractory
boat until the level of the wollastonite was substantially flush
with the top surface of the preform. A thin layer of molten
petroleum jelly
1TABLE I Number of Average Max. Min. Test Samples Strength Strength
Strength Sample Condition Temp. Tested (MPa) (MPa) (MPa) Metallic
constituent Room 8 461 .+-. 28 511 438 removed temp. Metallic
constituent 1200.degree. C. 10 488 .+-. 22 517 440 removed Metallic
constituent 1300.degree. C. 4 400 .+-. 12 412 386 removed Metallic
constituent 1400.degree. C. 4 340 .+-. 11 348 325 removed Metallic
constituent Room 3 288 .+-. 21 302 264 removed and heat temp.
treated at 1200.degree. C. in air for 24 h. Metallic constituent
1200.degree. C. 3 397 .+-. 9 404 387 removed and heat treated at
1200.degree. C. in air for 24 h. Metallic constituent Room 3 265
.+-. 12 275 253 removed and heat temp. treated at 1200.degree. C.
in air for 100 h. Metallic constituent 1200.degree. C. 3 401 .+-.
28 433 379 removed and heat treated at 1200.degree. C. in air for
100 h.
[0176] (VASELINE.TM., Cheesebrough-Ponds, Inc., Greenwich, Conn.)
was first applied to the top surface of the SiC/BN-coated fabric
preform and then covered with nickel oxide (NiO) powder
substantially in accordance of the methods of Example 2.
[0177] The porous castable refractory boat, having stilts at its
corners, was placed into a resistance heated air atmosphere furnace
and heated to about 700.degree. C. at a rate of about 400.degree.
C. per hour. A parent metal, comprising by weight about 7.5-9.5%
Si, 3.04.0% Cu, .ltoreq.2.9% Zn, 0.2-0.3% Mg, .ltoreq.1.5% Fe,
.ltoreq.0.5% Mn, .ltoreq.0.35% Sn, and the balance aluminum and
weighing about 420 grams, was also placed in a refractory container
in the resistance heated air atmosphere furnace and heated to about
700.degree. C. When parent metal was molten, the furnace door was
opened and the parent metal was poured into the heated porous
castable refractory boat and onto the NiO powder coated preform,
thereby covering the surface of the SiC/BN-coated fabric preform.
Wollastonite powder was then placed onto the surface of the molten
parent metal within the porous boat to substantially completely
cover the surface of the molten parent metal, thereby forming a
lay-up. Then the furnace and its contents comprising the lay-up
were heated to about 1000.degree. C. in about an hour. After about
60 hours at about 1000.degree. C., the furnace and its contents
were cooled to about 700.degree. C. in about 2 hours. At about
700.degree. C., the lay-up was removed from the furnace and
residual molten parent metal was decanted from the porous castable
refractory boat. The refractory boat was rapidly broken away from
the formed composite, and the formed composite was buried in silica
sand to permit the composite to cool to about room temperature. At
about room temperature, the composite was removed from the silica
sand and it was observed that an oxidation reaction product
comprising alumina had grown into and substantially completely
embedded the SiC/BN-coated fabric preform, thereby forming the
ceramic matrix composite body having a plurality of fabric layers
of 12 HSW ceramic grade NICALON.RTM. fibers silicon carbide as a
reinforcement. The ceramic matrix also comprised some residual
unreacted parent metal. The silicon carbide fiber reinforced
alumina composite body was then cut into bars measuring about 2.4
inches (60 mm) long by about 0.2 inch (6 mm) wide by about 0.11
Inch (3 mm) thick in preparation for the removal of at least a
portion of the metallic constituent of the formed fiber reinforced
ceramic composite body.
Sample B
[0178] A graphite foil box having an inner cavity measuring about
4.0 inches (102 mm) long by about 4.0 inches (102 mm) wide by about
3.0 inches (96 mm) deep was made from a piece of graphite foil
(GRAFOIL.TM., Union Carbide, Carbon Products Division, Cleveland,
Ohio) measuring about 10.0 inches (254 mm) long by about 10.0
inches (254 mm) wide by about 0.015 inch (0.38 mm) thick. Four
parallel cuts, 3.0 inches (76 mm) from the side and about 3.0
inches (76 mm) long were made into the graphite foil. The cut
graphite foil was then folded and stapled to form the graphite foil
box.
[0179] A parent metal, comprising by weight about 3 percent
strontium and the balance aluminum and measuring about 4.0 inches
(102 mm) long by about 4.0 inches (102 mm) wide by about 1.0 inch
(25 mm) thick was coated on one side thereof measuring about 4.0
inches (102 mm) long by about 4.0 inches (102 mm) wide with a
slurry comprising by weight about 90% -325 mesh (particle size less
than about 45 Am) aluminum alloy powder and the balance ethanol.
The -325 mesh aluminum alloy powder was nominally comprised by
weight of about 7.5-9.5% Si, 3.04.0% Cu, .ltoreq.2.9% Zn, 0.2-0.3%
Mg, .ltoreq.1.5% Fe, .ltoreq.0.5% Mn, .ltoreq.0.35% Sn, and the
balance aluminum. The aluminum alloy powder-coated parent metal was
then placed into the graphite foil box such that the uncoated
surfaces of the parent metal contacted the inner surfaces of the
graphite foil box.
[0180] A fabric preform measuring about 4.0 inches (102 mm) long by
about 4.0 inches (102 mm) wide by about 0.06 inch (1.6 mm) thick
was made within the graphite foil box and on the aluminum alloy
powder coated surface of the parent metal by stacking four layers
of 12 harness satin weave (HSW) silicon carbide fabric (ceramic
grade NICALON.RTM. silicon carbide fibrous material obtained from
Dow Corning Corporation, Midland, Mich.) onto the parent metal.
About 0.5 inch (13 mm) of a 500 grit (average particle diameter of
about 17 .mu.m) alumina powder (E1 ALUNDUM.TM., Norton Company,
Worcester, Mass.) was poured over the 12 HSW fabric preform and
leveled. The sides of the graphite foil box that extended beyond
the level of the alumina powder covering the 12 HSW fabrics were
folded over onto the alumina powder to form a lid for the graphite
foil box.
[0181] A lay-up was formed in a graphite refractory container by
placing and leveling about 0.5 inch (13 mm) of a 500 grit (average
particle diameter of about 17 .mu.m) alumina powder into the bottom
of the graphite refractory container. The graphite foil box and its
contents comprising the aluminum alloy powder-coated parent metal
and the 12 HSW silicon carbide fabric preform were placed into the
graphite refractory container and onto a 500 grit (average particle
diameter of about 17 .mu.m) alumina. Additional 500 grit alumina
was placed into the graphite refractory container into the void
defined by the inner surface of the graphite refractory container
and the outer surface of the graphite foil box. The 500 grit
(average particle diameter of about 17 .mu.m) alumina powder also
covered the top lid of the graphite foil box and its contents.
[0182] The lay-up comprising the graphite refractory container and
its contents was placed into a retort lined resistance heat furnace
and the retort door was closed. The furnace and its contents were
heated to about 100.degree. C. at a rate of about 300.degree. C.
per hour. At about 100.degree. C., the retort was evacuated to
about 30.0 inches (762 mm) mercury (Hg) vacuum and maintained at
about 30.0 inches (762 mm) Hg vacuum to about 150.degree. C. At
about 150.degree. C., nitrogen was introduced into the retort at a
flow rate of about 4 liters per minute. The furnace and its
contents were then heated to about 900.degree. C. at about
300.degree. C. per hour. After about 200 hours at about 900.degree.
C., the furnace and its contents were cooled to about room
temperature at a rate of about 300.degree. C. per hour. At about
room temperature, the retort door was opened and the lay-up was
removed. The lay-up was disassembled, the preform was removed from
within the graphite foil box, and it was observed that an oxidation
reaction product comprising aluminum nitride had grown into and
substantially completely embedded the silicon carbide fabric
preform thereby forming a ceramic matrix composite body reinforced
with a plurality of fabric layers of 12 HSW ceramic grade
NICALON.RTM. silicon carbide as reinforcement. The ceramic matrix
also comprised a metallic constituent comprising residual unreacted
parent metal.
[0183] Table II contains a summary of the parameters used to
practice the metal removal step of the instant invention on Samples
A and B. Specifically, Table II contains the dimensions of the
sample, the filler material used for metal removal, the
infiltration enhancer precursor, the processing temperature, the
processing time at the processing temperature, and the processing
atmosphere.
[0184] FIG. 6 shows a cross-sectional schematic of the setup used
in this series of tests to remove the metallic constituent from
Samples A and B.
[0185] After the formation of the silicon carbide fiber reinforced
alumina composite body of Sample A had been achieved, the metal
removal process was effected. Specifically, a filler material
mixture was formed, comprising by weight about 90 percent filler,
which comprised 1000 grit (average particle diameter of about 5
.mu.m) Al.sub.2O.sub.3 (E67 tabular alumina, Norton Co., Worcester,
Mass.) and about 10 percent by weight -325 mesh (particle diameter
less than about 45 .mu.m) magnesium powder (AESAR.TM., Johnson
Matthey, Seabrook, N.H.). The filler material mixture was mixed in
a plastic jar on a rotating jar mill for about an hour.
[0186] A graphite foil box having an inner cavity measuring about 3
inches (76 mm) long by about 3 inches (76 mm) wide and about 2.5
inches (64 mm) deep was made from graphite foil (PERMA FOIL, TT
America, Portland, Oreg.). The graphite foil box was made from a
piece of graphite foil, measuring about 8 inches (203 mm) long by
about 8 inches (203 mm) wide by about 0.15 inches (4 mm) thick.
Four parallel cuts about 2.5 inches (64 mm) from the side and about
2.5 inches (64 mm) long, were made into the graphite foil. The
graphite foil was then folded into a graphite foil box
2TABLE II Infiltration Processing Time Filler Material Enhancer
Processing At Processing Sample ID Composite.sup.1 Geometry For
Metal Removal Precursor Temperature Temperature Atm. A
SiC.sub.f/Al.sub.2O.sub.3.sup.7 bar 1000 grit Al.sub.2O.sub.3.sup.2
10% - 325 mesh Mg.sup.3 850.degree. C. 10 h N.sub.2 B
SiC.sub.5/AlN.sup.7 irregular 1000 grit Al.sub.2O.sub.3 10% - 325
mesh Mg.sup.3 750.degree. C. 10 h N.sub.2 .sup.1SiC fiber
reinforced composite .sup.2E-67 alumina, Norton Co., Worcester, MA.
.sup.3AESAR .RTM., Johnson Matthey Corporation, Seabrook, New
Hampshire
[0187] and stapled together. Metal was removed from Sample A by
first pouring about 0.5 inch (13 mm) of the mixture of filler
material and magnesium powder into one of the graphite foil boxes.
The filler material mixture was levelled and hand tapped until
smooth. A bar of the silicon carbide fiber reinforced alumina
composite of Sample A, and measuring about 1.7 inches (43.8 mm)
long by about 0.25 inch (6.3 mm) wide by about 0.2 inch (4.5 mm)
thick was placed onto the filler material mixture within the
graphite foil box and covered with another about 0.5 inch (13 mm)
of the filler material mixture which was again levelled and hand
tapped until smooth.
[0188] The graphite foil box containing Sample A was then placed
into a graphite refractory container having inner dimensions of
about 9 inches (229 mm) long by about 9 inches (229 mm) wide by
about 5 inches (127 mm) deep and having a wall thickness of about
0.5 inch (13 mm). The graphite refractory container and its
contents were then placed into a controlled atmosphere resistance
heated furnace, the furnace door was closed and the furnace was
evacuated to about 30 inches (762 mm) Hg. After about 15 hours at
about 30 inches (762 mm) of mercury vacuum, the vacuum was shut off
and nitrogen gas was introduced into the furnace chamber at a flow
rate of about 1 liter/minute. The operating pressure of the chamber
was about 16.7 pounds per square inch (1.2 kg/cm.sup.2) with a
nitrogen flow rate of about 1 liter/minute. The furnace was heated
to about 850.degree. C. at about 200.degree. C. per hour. After
about 10 hours at about 850.degree. C., the power to the furnace
was interrupted and the graphite refractory container and its
contents were allowed to cool within the furnace to about room
temperature. Once at room temperature, the graphite refractory
container and its contents were removed and the lay-up for Sample A
was disassembled to reveal that the metallic constituent comprising
an aluminum alloy in the silicon carbide fiber reinforced alumina
composite had been drawn out from the composite body during the
metal removal process.
[0189] The setup for the removal of the metallic constituent from
Sample B was substantially the same as that described for Sample A
of this Example and is schematically illustrated in FIG. 6. The
nitrogen flow rate to effect removal of the metallic constituent
from Sample B was about two liters per minute. The controlled
atmosphere furnace was heated to about the processing temperature
of about 750.degree. C. at a rate of about 200.degree. C. per hour,
held at about the processing temperature for about 10 hours. After
about 10 hours at the processing temperature, at least a portion of
the metallic constituent was removed from within the ceramic matrix
composite body.
[0190] Specifically, the metallic constituent spontaneously
infiltrated the filler material mixture comprising substantially a
1000 grit (average particle diameter of about 5 .mu.m) alumina and
a -325 mesh magnesium infiltration enhancer precursor. The furnace
and its contents were cooled to about room temperature. At about
room temperature, the setup was removed from the furnace,
disassembled, and weight loss due to the removal of the metallic
constituent from Sample B was noted.
EXAMPLE 4
[0191] The following Example demonstrates that fiber reinforced
ceramic composite bodies formed by the method of the present
invention maintain substantially their room temperature fracture
toughness at elevated temperatures. A series of fiber preforms were
made substantially in accordance with the methods described in
Example 2, except that the first layer and eighth layer of the
fabric preform comprised 12 harness satin weave (12 HSW) fabric
instead of 8 harness satin weave (8 HSW) fabric and the temperature
of the methyltrichlorosilane (MTS) bath used during the formation
of silicon carbide coatings was maintained at about 18.degree. C.
instead of about 25.degree. C. The lay-up for the growth of the
fiber reinforced ceramic composite body included an alumina-bonded
wollastonite shell fabricated substantially in accordance with the
methods described in Example 2, and the composite growth process
was substantially the same as that described in Example 2. The
resultant ceramic matrix composite bodies were subjected to a metal
removal treatment substantially the same as that described in
Example 2. The samples were subsequently machined to form
mechanical test samples which were used to determine both the
flexural strength and the fracture toughness of the fiber
reinforced ceramic composite bodies both as a function of test
temperature.
[0192] Table III summarizes the results of these tests. The methods
for measurement of the flexural strength was substantially in
accordance with the methods described in Example 2. The method of
Munz, Shannon and Bubsey (International Journal of Fracture, Vol.
16 (1980) R137-R141) was used to determine the fracture toughness
of the silicon carbide fiber reinforced ceramic composite bodies.
The fracture toughness was calculated from the maximum load of
Chevron notch specimens in four point loading. Specifically, the
geometry of each Chevron notch specimen was about 1.8 to 2.2 inches
(45 to 55 mm) long, about 0.12 inch (3 mm) wide and about 0.15 inch
(3.75 mm) high. A Chevron notch was cut in each specimen with a
diamond saw to permit the propagation of a crack starting at the
notch and traveling through the sample. The Chevron notched
specimens, having the apex of the Chevron notch pointing downward,
were placed into a fixture within a Universal test machine. The
notch of the Chevron notch specimen, was placed between two pins
about 1.6 inches (40 mm) apart and about 0.79 inch (20 mm) from
each pin. The top side of the Chevron notch specimen was contacted
by two pins about 0.79 inch (20 mm) apart and about 0.39 inch (10
mm) from the notch. The maximum load measurements were made with a
Syntec Model CITS-2000/6 universal testing machine (System
Integration Technology Incorporated, Stoughton, Mass.). A crosshead
speed of 0.02 inches/minute (0.58 millimeters/minute) was used. The
load cell of the universal testing machine was interfaced to a
computer data acquisition system. The Chevron notch sample geometry
and maximum load were used to calculate the fracture toughness of
the material. Several samples were used to determine an average
fracture toughness for a given group of parameters (e.g.,
temperature, fiber reinforced ceramic composite body, etc.)
[0193] Table III summarizes the results of the measurements of the
average flexural strength, the maximum flexural strength and the
average fracture toughness all as a function of temperature, for
Samples D, E and F, which were subjected to the metal removal
process. Moreover, the fracture toughness of an "as-grown" Sample C
(e.g., without any residual metallic constituent removed) is
compared to a treated Sample D (i.e., metallic constituent
removed). The data in Table III shows that the fracture toughness
of a fiber reinforced ceramic composite body with its metallic
constituent substantially completely removed is not significantly
diminished at elevated temperatures. In addition, the fracture
toughness of a sample which is subjected to the metal removal
process does not appear to vary significantly from the fracture
toughness of an untreated composite body.
EXAMPLE 5
[0194] The following Example demonstrates that fiber reinforced
ceramic composite bodies exhibiting excellent fracture toughness
can be produced by (1) coating a fabric preform with coatings
comprising silicon carbide (SiC)/boron nitride (BN); (2) growing an
oxidation reaction product by a reaction of a parent metal with an
oxidant which embeds the SiC/BN-coated fabric preform and (3)
removing at least some of the metallic constituent from the grown
fiber reinforced ceramic composite body. A ceramic grade
NICALON.RTM. silicon carbide fiber reinforced alumina composite
body plate measuring substantially the same as that in Example 2
was formed substantially in accordance with the method of Example
2. Specifically, the fabric preform lay-up, the formation of both
the boron nitride and silicon carbide coatings, the growth of the
alumina oxidation reaction product embedding the SiC/BN-coated
fabric preform and the removal of the metallic constituent from the
fiber reinforced ceramic body were performed substantially in
accordance with the method of Example 2.
3TABLE III Average Maximum Average Flexural Flexural Fracture
Sample Sample Test Strength Strength Toughness ID Condition Temp.
(MPa) (MPa) (MPa-m.sup.1/2) C As Grown RT -- -- 19 .+-. 1 D
Metallic RT 450 (31)* 563 21 .+-. 1 constituent removed E Metallic
1000.degree. C. 400 (7)* 432 23 .+-. 1 constituent removed F
Metallic 1200.degree. C. 350 (14)* 406 18 .+-. 1 constituent
removed *The number in parenthesis indicates the number of sample
test.
[0195] The fracture toughness of the fiber reinforced ceramic
composite body was measured substantially in accordance with the
method of Example 4, except that specimen size used to determine
the toughness measured from about 1.0 to about 1.2 inches (25 to 30
mm) long, about 0.15 inch (3.75 mm) high and about 0.12 inch (3 mm)
wide. The apex of the Chevron notch pointed up within the universal
test machine. The notch of the specimen was placed between two pins
about 0.39 inch (10 mm) apart and about 0.2 inch (5 mm) from each
pin. The top side of the specimen was contacted by two pins about
0.79 inch (20 mm) apart and about 0.39 inch (10 mm) from the notch.
Three specimens were tested to determine an average fracture
toughness for a specific test temperature.
[0196] The fracture toughness of the fiber reinforced ceramic
composite body of this Example was measured at about room
temperature, at about 1200.degree. C. and at about 1300.degree. C.
These values were about 35.3.+-.1 MPa-m.sup.1/2, 19.6.+-.1
MPa-m.sup.1/2 and 18.7.+-.1 MPa-m.sup.1/2, respectively.
EXAMPLE 6
[0197] The following Example demonstrates the intrinsic strength of
the ceramic matrix of a fiber reinforced ceramic composite
body.
[0198] A ceramic grade NICALON.RTM. silicon carbide fiber
reinforced alumina composite was formed substantially in accordance
with the methods of Example 2. Specifically, the fabric preform
lay-up, the formation of both the boron nitride and silicon carbide
coatings, the growth of the alumina oxidation reaction product
embedding the SiC/BN-coated fiber and the removal of the metallic
constituent from the fiber reinforced ceramic body were performed
substantially in accordance with the method of Example 2.
[0199] The intrinsic strength of the matrix was measured at about
room temperature with the short beam method according to ASTM
method D 2344-84 entitled "Standard Test Method for Apparent
Interlaminar Shear Strength of Parallel Fiber Composite By
Short-Beam Method."
[0200] The mechanical test specimens were machined to overall
dimensions of about 1 inch (25 mm) in length by about 0.16 inch (4
mm) in width by about 0.16 inch (4 mm) in thickness. Furthermore,
the orientation of the mechanical test specimens were such that all
the fibers were perpendicular to the thickness dimension, i.e.,
none of the fibers traversed the thickness dimension.
[0201] This test was specifically designed to measure the strength,
and in particular, the shear strength, of the matrix material
between two adjacent layers of the eight total layers of HSW
fabric.
[0202] A three-point flexural test was used. The thickness and
width of the test bars was measured with a precision of about 390
microinch (0.01 mm). The test bars were subjected to a stress which
was applied at three points by two lower span bearing points and
one upper span bearing point. The lower span bearing points were
about 0.67 inch (17 mm) apart and the upper load point was centered
over the lower span so that the load was applied substantially
symmetrically on the test bar. The flexural strength measurements
were made with a Syntec Model No. CITS-2000/6 universal testing
machine (System Integration Technology, Inc., Stoughton, Mass.)
having a 500 pound (2225 N) full-scale deflection load cell. A
computer data acquisition system was connected to the measuring
unit and strain gauges in the load cell recorded the test
responses. The cross-head speed during testing was about 0.05 inch
per minute (1.3 mm per minute).
[0203] The interlaminar shear strength was found to be about 62
MPa.
EXAMPLE 7
[0204] This Example characterizes the tensile strength of a fiber
reinforced ceramic composite body and shows the gradual and
progressive failure of such a body as opposed to the sudden and
catastrophic failure typical of most ceramic or ceramic composite
bodies.
[0205] A ceramic grade NICALON.RTM. silicon carbide fiber
reinforced alumina matrix composite was formed substantially in
accordance with the method of Example 2. Specifically, the fabric
preform lay-up, the formation of both the boron nitride and silicon
carbide coatings, the growth of the alumina oxidation reaction
product embedding the SiC/BN-coated fiber and the removal of the
metallic constituent from the fiber reinforced ceramic body were
performed substantially in accordance with the method of Example
2.
[0206] The tensile strength of the fiber reinforced ceramic
composite body was measured using the procedures described in ASTM
designations A 370 and E 8M-88.
[0207] FIG. 7 shows the approximate shape of the test specimen
which was machined using diamond grinding with the longitudinal
axis of the test specimen parallel to either the length or width
dimension of the fiber preform. The tensile test specimen measured
overall about 6 inches (152 mm) long by about 0.5 inch (13 mm) wide
by about 0.12 inch (3 mm) thick. The gage section measured about
0.75 inch (19 mm) long by about 0.35 inch (9 mm) wide. The test was
performed using an MTS Model 810 universal testing machine (MTS
Systems Corp., Eden Prarie, Minn.) operated at a crosshead speed of
about 0.25 mm per minute. The sample strain was monitored with an
MTS Model 632-11B-20 clip-on extensometer (MTS Systems Corp.).
[0208] At room temperature, the average tensile strength for 14
samples was about 331 MPa with a standard deviation of about 22
MPa. The Young's Modulus, as measured by the ratio of stress to
strain in the linear portion of the stress-strain curve, averaged
about 162 GPa and the average strain-to-failure was about 0.645
percent.
[0209] FIG. 8 shows a typical stress-strain curve for a fiber
reinforced ceramic composite body made substantially by the method
of Example 2. The stress-strain curve begins to deviate from
linearity at a stress of about 50-60 MPa, which deviation indicates
the onset of matrix microcracking and pull-out of the reinforcing
fibers from the surrounding matrix material.
[0210] FIG. 9 is a scanning electron micrograph taken at about
50.times.magnification of a fracture surface which has been exposed
as a result of a room temperature tensile test. Segments of the
reinforcing fibers which have been partially pulled out of the
surrounding matrix material are clearly visible.
EXAMPLE 8
[0211] This Example demonstrates that fiber reinforced ceramic
matrix composites produced according to the method of the present
invention retain almost all of their ambient temperature strength
at elevated temperatures, even after repeated thermal cycling.
[0212] A fabric preform 103 was made by stacking a plurality of
layers of 8 harness satin weave (8 HSW) fabric and 12 harness satin
weave (12 HSW) fabric made from NICALON.RTM. silicon carbide fiber
(ceramic grade, obtained from Dow Corning Corp., Midland, Mich.) on
top of each other. The nomenclature describing the orientations of
the fabrics is substantially the same as that used in Example 2 and
depicted in FIGS. 5a, 5b and 5c.
[0213] The fabric preform of the present Example was made by
stacking the layers of HSW fabric in the following sequence:
[0214] A first fabric layer comprising an 8 HSW fabric was rotated
about 90.degree. in the counterclockwise direction from the as-is
position about an axis 93 perpendicular to the plane of the fabric
and was placed on a supporting surface to start the fabric
preform;
[0215] A second fabric layer comprising an 8 HSW fabric was placed
on the first fabric layer in the as-is position so that the edges
of the second fabric layer were substantially aligned with the
edges of the first fabric layer;
[0216] A third fabric layer comprising a 12 HSW fabric was rotated
about 90.degree. in the counterclockwise direction from the as-is
position about an axis 93 perpendicular to the plane of the fabric
and was placed on the second fabric layer so that the edges of the
third fabric layer were substantially aligned with the edges of the
second fabric layer;
[0217] A fourth fabric layer comprising a 12 HSW fabric was placed
on the third fabric layer in the as-is position so that the edges
of the fourth fabric layer were substantially aligned with the
edges of the third fabric layer;
[0218] A fifth fabric layer comprising a 12 HSW fabric was rotated
about 90.degree. in the counterclockwise direction from the as-is
position about an axis 93 perpendicular to the plane of the fabric
and was placed on the fourth fabric layer so that the edges of the
fifth fabric layer were substantially aligned with the edges of the
fourth fabric layer;
[0219] A sixth fabric layer comprising an 8 HSW fabric was placed
on the fifth fabric layer in the as-is position so that the edges
of the sixth fabric layer were substantially aligned with the edges
of the fifth fabric layer;
[0220] A seventh fabric layer comprising an 8 HSW fabric was
rotated about 90.degree. in the counterclockwise direction from the
as-is position about an axis 93 perpendicular to the plane of the
fabric and was placed on the sixth fabric layer so that the edges
of the seventh fabric layer were substantially aligned were
substantially aligned with the edges of the sixth fabric layer,
thus completing the rectangular fabric preform which measured about
7 inches (178 mm) in length by about 5 inches (127 mm) in
width.
[0221] The fabric preform was clamped in substantially the same
fixture as was described in Example 2 and depicted in FIG. 5e. The
preform containment fixture 108 containing the fabric preform was
placed into a reactor chamber of a refractory alloy steel chemical
vapor infiltration apparatus having a graphite tube liner and
having overall dimensions of about 8 feet (2.4 meters) in length by
about 15.5 inches (394 mm) in inside diameter. The warp yams of the
first and seventh layers of the fabric preform were perpendicular
to the gas flow direction within the chamber as well as being
perpendicular to the longitudinal axis of the reactor chamber. The
reactor chamber was closed and evacuated to less than about 0.04
inch (1 mm) of mercury (Hg). The reactor chamber was then heated to
a temperature of about 820.degree. C. Argon gas was flowed into the
annulus region between the graphite liner and the steel reactor
wall at a rate of about 850 standard cubic centimeters per minute
(sccm). When the temperature within the reactor chamber reached
about 820.degree. C., a gas mixture comprising borontrichloride
(BCl.sub.3) flowing at about 700 sccm at a temperature of about
60.degree. C. and ammonia (NH.sub.3) flowing at about 1800 sccm was
introduced into the reactor chamber while maintaining a total
operating pressure of about 0.5 torr. After about 7 hours at a
temperature of about 820.degree. C., the gas mixture flowing into
the reactor chamber was interrupted, the power to the furnace
heating the reactor chamber was interrupted and the furnace and its
contents were naturally cooled. At a temperature below about
200.degree. C., the reactor chamber door was opened and the
graphite containment fixture was removed, cooled and disassembled
to reveal that the fibers of the fabric layers of the fabric
preform were coated and that the fabric layers comprising the
fabric preform were bonded together by a boron nitride coating. The
boron nitride coating had a thickness of about 0.48 micron.
[0222] The boron nitride coated fabric preform was then stored in a
vacuum desiccator until it was ready to be put back into the
chemical vapor infiltration apparatus for additional coating.
[0223] For the application of this subsequent coating, the boron
nitride coated and bonded fabric preform was placed back into the
reactor chamber of the chemical vapor infiltration apparatus. In
this instance, however, the warp yarns of the first and seventh
layers of the fabric preform were parallel to the gas flow
direction within the chamber, as well as being parallel to the
longitudinal axis of the reactor chamber. The reactor chamber was
closed and evacuated to about less than about 1 torr. Hydrogen gas
was introduced into the reactor chamber at a flow rate of about
5000 standard cubic centimeters per minute (sccm). The reactor
chamber was then heated to a temperature of about 935.degree. C.
Nitrogen gas was flowed through the annulus region at a rate of
about 850 sccm. Once the temperature of the contents of the reactor
chamber had substantially completely stabilized at about
935.degree. C., about 1500 sccm of hydrogen were diverted away from
direct entry into the reactor chamber and were first bubbled
through a bath of methyltrichlorosilane (MTS) maintained at a
temperature of about 45.degree. C. before entering the reactor
chamber. After about 20 hours at a temperature of about 935.degree.
C., the power to the furnace heating the reactor chamber was
interrupted and the about 1500 sccm of hydrogen that was being
directed through the MTS bath was again permitted to flow directly
into the reactor chamber to re-establish a direct hydrogen gas flow
rate of about 5000 sccm into the reactor chamber. After the reactor
chamber had cooled substantially, the hydrogen flow rate was
interrupted and the furnace and its contents were evacuated to less
than 1 torr. The pressure within the reactor chamber was then
brought back up to about atmospheric pressure with argon gas. After
the reactor chamber had cooled to a temperature below about
200.degree. C., the argon gas flow rate was interrupted and the
reactor chamber door was opened. The graphite containment fixture
was removed, cooled and disassembled to reveal that the boron
nitride bonded fabric preform had been coated with a second layer
of silicon carbide thereby forming a silicon carbide (SiC)/boron
nitride (BN)-coated fabric preform. The silicon carbide had a
thickness of about 1.9 microns.
[0224] Growth of an alumina oxidation reaction product through the
silicon carbide/boron nitride-coated fabric preform was then
carried out in substantially the same manner as was described in
Example 2 to form a fiber reinforced ceramic composite body
comprising a ceramic matrix comprising an aluminum oxide oxidation
reaction product and a metallic component in comprising some
residual unreacted parent metal, with said ceramic matrix
reinforced by the silicon carbide/boron nitride coated NICALON.RTM.
silicon carbide fibers (ceramic grade). Substantially complete
growth of the ceramic matrix only required about 72 hours,
however.
[0225] Once the ceramic composite body had been manufactured, at
least a portion of the metallic constituent comprising the ceramic
matrix was removed. This metal removal process was performed in
substantially the same manner as was described in Example 2.
[0226] Tensile test specimens were machined from the fiber
reinforced ceramic composite body and tested in substantially the
same manner as described in Example 7. Heating was provided by
positioning a resistance heated air atmosphere furnace in the
testing zone of the test machine. The samples were tested in air at
ambient as well as at elevated temperatures of about 1100.degree.
C., 1200.degree. C., and about 1370.degree. C. As shown in FIG. 10,
the tensile strength at these temperatures was about 260, 250, 260,
and about 230 MPa, respectively. Thus, these data show that the
fiber reinforced ceramic composite material retains substantially
all of its ambient temperature strength up to a temperature of
about 1200.degree. C., and almost all of its ambient temperature
strength at a temperature of about 1370.degree. C.
[0227] Next, the effect of repeated thermal cycling on the
material's tensile strength was assessed.
[0228] First, a ceramic grade NICALON.RTM. silicon carbide
reinforced alumina matrix composite was produced substantially in
accordance with the procedure described in Example 2 and likewise
subjected to the metal removal process of Example 2. Unlike the
procedure of Example 2, however, during the growth of oxidation
reaction product into the preform, the approximately 950.degree. C.
process temperature was maintained for about 100 hours instead of
about 90 hours. Moreover, the thickness of the silicon carbide
coating deposited onto the boron nitride coated NICALON.RTM.
silicon carbide fibers during chemical vapor infiltration was about
2.0 microns.
[0229] Substantially rectangular tensile test specimens were
diamond machined from the composite tile such that the length
dimension of the test specimen was oriented parallel to the length
or width dimension of the composite tile.
[0230] About half of the specimens were given a rapid thermal
cycling treatment before tensile testing; the others were tested
"as is". Specifically, the thermal cycling comprised subjecting
each composite test specimen to about 150 thermal cycles, each
thermal cycle comprising heating a test specimen from a starting
temperature to a temperature of about 1200.degree. C. in an argon
atmosphere at a rate of about 40.degree. C. per minute, holding at
a temperature of about 1200.degree. C. for about 2 minutes, and
cooling back to the starting temperature at a rate of about
10.degree. C. per minute. The starting temperature corresponded to
the final testing temperature. The two sets of tensile test
specimens were then tested in substantially the same manner as was
described in the preceding Example at about room temperature and
temperatures of about 1000.degree. F. (538.degree.) about
1500.degree. F. (816.degree. C.) and at about 2000.degree. F.
(1093.degree. C.).
[0231] FIG. 11 shows the tensile strength as a function of test
temperature for the two sets of composite test specimens. The data
show that the thermally cycled composite test specimen experienced
little loss in tensile strength compared to their counterparts
which were not thermally cycled. The significance of this result is
that the thermal cycling provided an opportunity for chemical
reaction between the fiber, the fiber coatings and the surrounding
matrix constituents. The thermal cycling operation also provided an
opportunity for cracking due to thermal expansion mismatch. The
lack of significant strength reduction indicates that any
microcracking induced by the thermal cycling was confined to the
matrix material and, furthermore, that the ability of the fibers to
pull out of the matrix under the applied tensile load was not
substantially affected by the thermal cycling. The different
tensile strength levels observed in comparing the data of FIG. 10
to that of FIG. 11 may be attributable to variations in preform
fabrication, specifically, such as the differences in the number of
each type of HSW fabric (e.g., 12 HSW vs 8 HSW).
EXAMPLE 9
[0232] This Example demonstrates the high temperature mechanical
performance of a fiber reinforced ceramic composite body under an
applied load over a prolonged period of time in an oxidizing
atmosphere.
[0233] The fiber reinforced ceramic composite body described herein
was fabricated substantially in accordance with the methods
outlined in Example 2. Specifically, the fabric preform lay-up, the
formation of both the boron nitride and silicon carbide coatings,
the growth of the alumina oxidation reaction product embedding the
SiC/BN-coated fiber and the removal of the metallic constituent
from the fiber reinforced ceramic body were performed substantially
in accordance with the method of Example 2.
[0234] In Example 7, it was demonstrated that at room temperature
(e.g., about 20.degree. C.) in a pure tensile test, a fiber
reinforced ceramic matrix composite sample begins to deviate from
linear stress/strain behavior at an applied stress of about 50-60
MPa, indicating that the matrix begins to microcrack at
approximately this stress level. These microcracks may allow for
oxygen in the surrounding atmosphere to find a path to the
underlying ceramic grade NICALON.RTM. silicon carbide fiber and/or
its SiC and BN coatings. Accordingly, stress rupture tests were
conducted at various elevated temperatures in air at applied
stresses above this 50-60 MPa microcracking threshold in order to
evaluate the impact of matrix microcracking and subsequent oxygen
ingress on the performance of the fiber reinforced ceramic
composite body.
[0235] The stress rupture test specimen had substantially the same
shape as that depicted in FIG. 7, with the exception that shoulders
were machined into each end of the test specimen so that the sample
could be gripped by a collar in the test fixture rather than
clamped. Mica powder was used in the collar to cushion the contact
zone between the collar and the shoulder portions of the stress
rupture test specimen. The test specimen measured about 5.5 inches
(140 mm) long overall by about 0.5 inch (13 mm) wide by about 0.12
inch (3 mm) thick. The gage portion of the test specimen measured
about 2 inches (51 mm) in length by about 0.2 inches (5 mm)
wide.
[0236] The tests comprised heating the samples to the desired test
temperature and loading each specimen in tension to a desired
stress and maintaining said stress at said temperature. The applied
stress was increased in a step-wise manner. The unit length change
of the specimen within the gage portion of the overall test
specimen was monitored with a Model 1102 ZYGO.TM. helium-neon laser
extensometer (Zygo Corp., Middlefield, Conn.).
[0237] The results of the stress rupture testing are presented for
FIG. 12. The particulars of the applied stress and the exposure
times are presented below.
Sample G
[0238] The test fixture, comprising the Sample G test specimen with
collars attached to each end, was loaded into a Model P-5 creep
testing machine (SATEC Inc., Grove City, Pa.). A tensile stress of
about 12.5 megapascals was applied to the test specimen using dead
loading. A resistance heated air atmosphere furnace was positioned
completely around the stress rupture test specimen and the furnace
and the stress rupture sample contained within were heated from
about room temperature to a temperature of about 1000.degree. C.
over a period of about 2 hours.
[0239] After the furnace chamber and its contents had reached a
temperature of about 1000.degree. C., the stress applied to the
sample was increased to about 75 MPa. After maintaining an applied
stress of about 75 MPa for about 70 hours, the applied stress to
the sample was increased to about 100 Mpa. After about 15 hours at
a stress of about 100 MPa, the sample broke. The furnace chamber
and its contents were allowed to cool naturally back down to about
room temperature.
Sample H
[0240] The Sample H test fixture was placed into the creep testing
machine at about room temperature and the Sample H stress rupture
test specimen was heated in the surrounding resistance heated air
atmosphere furnace to a temperature of about 1000.degree. C. over a
period of about 3 hours under an applied stress of about 5 MPa. At
a temperature of about 1000.degree. C., the applied tensile stress
on the sample was increased to about 70 MPa and the temperature
inside the furnace chamber was increased to about 1100.degree. C.
over a period of about 1 hour. After maintaining the sample in
tension at a stress of about 70 MPa at a temperature of about
1100.degree. for about 210 hours, the applied stress was increased
to about 83 MPa. After about an additional 6 hours, the stress was
increased to about 85 MPa. After maintaining an applied stress of
about 85 MPa on the sample for about 115 hours, the applied stress
was increased to about 88 MPa. After maintaining an applied stress
of 88 MPa for about 1.5 hours, the stress applied was increased to
about 90 MPa. After maintaining an applied tensile stress of about
90 MPa for about 3 hours, the applied stress was increased to about
91 MPa.
[0241] After maintaining an applied stress of about 91 MPa for
about 1.5 hours, the stress was further increased to about 92 MPa.
After maintaining an applied stress of about 92 MPa for about 1.3
hours, the applied stress was increased to about 95 MPa. After
maintaining an applied stress of about 95 MPa on the sample for
about 115 hours, the applied stress was increased to about 96 MPa.
After maintaining an applied stress of about 96 MPa for about 3
hours, the applied stress was increased to about 97 MPa. After
maintaining an applied stress of about 97 MPa for about 2 hours,
the applied stress was increased to about 99 MPa. After maintaining
an applied stress of about 99 MPa for about 1.5 hours, the applied
stress was increased to about 100 MPa. After maintaining an applied
stress of about 100 MPa for about 60 hours, the sample broke. The
furnace chamber and its contents were thereafter furnace cooled
from a temperature of about 1100.degree. C. down to about room
temperature.
[0242] The fractured sample was recovered from the test chamber and
the fracture surface was examined in the scanning electron
microscope. FIG. 13 is an approximately 50.times.magnification
scanning electron micrograph of a portion of the fracture surface.
Direct comparison of FIG. 13 with the previous scanning electron
micrograph of FIG. 9 shows much less fiber pull-out associated with
this Sample H specimen than with the fracture surface of the
Example 7 tensile test specimen. This decrease in the degree of
fiber pull-out of the present stress rupture may suggest
degradation of the fiber and/or one or more of its coatings over
the 500+hour duration of the stress rupture test. Conversely, the
ability of this fiber reinforced ceramic matrix composite body to
survive sustained exposure of this duration at a temperature of
about 1100.degree. C. at a stress level sufficient to expose the
reinforcing fibers and/or their coatings to atmospheric oxygen may
suggest the operation of a mechanism working to protect the
NICALON.RTM. fibers from chemical reactions such as atmospheric
oxidation.
[0243] FIGS. 14a, 14b and 14c are scanning electron micrographs
taken at about 2500.times., 500.times. and
10,000.times.magnification of a diamond polished cross-section of
the Sample H stress rupture test specimen at a region very close to
the fracture surface. Specifically, FIG. 14a shows a crack
breaching at least the SiC coating, thus potentially exposing the
NICALON.RTM. fiber and/or the BN debond coating to chemical
reaction with reactant supplied from outside the fiber and its
coatings. The higher magnification of this crack region shown in
FIG. 14b reveals the presence of a substance at least partially
filling the crack. Such a substance may comprise a reaction product
of one or both of the SiC and BN coatings and/or the NICALON.RTM.
fiber itself. The presence of such a reaction product may explain
the apparent degradation of the fiber pull-out mechanism as well as
the relative longevity of the material while under load at elevated
temperature. Specifically, the at least partial re-filling of a
matrix microcrack after such a crack forms may serve to reduce the
access of, for example, atmospheric oxygen to the reinforcing
fibers and their coatings. FIG. 14c shows a different matrix
microcrack in Sample H breaching an SiC coating.
[0244] This particular micrograph appears to show that the
substance substantially filling the crack in the SiC coating also
substantially comprises the space between the SiC coating and the
NICALON.RTM. fiber and the space between the SiC coating and the
alumina oxidation reaction product.
Sample I
[0245] Sample E was stress rupture tested at a temperature of about
1200.degree. C. The sample was loaded into the test rig in
substantially the same manner as was described for Sample G. A
tensile stress of about 12.5 MPa was applied to the test specimen
at about room temperature. The furnace chamber and its contents
were then heated from about room temperature to a temperature of
about 1200.degree. C. over a period of about 3 hours. At a
temperature of about 1200.degree. C., the applied stress was
increased to about 66 MPa. After maintaining a temperature of about
1200.degree. C. at an applied stress of about 66 MPa for about 256
hours, the applied stress was increased to about 70 MPa. After
maintaining an applied stress of about 70 MPa at a temperature of
about 1200.degree. C. for about 216 hours, the applied stress was
increased to about 75 MPa. After maintaining an applied stress of
about 75 megapascals at a temperature of about 1200.degree. C. for
about 288 hours, the applied stress was increased to about 80 MPa.
After maintaining an applied stress of about 80 MPa at a
temperature of about 1200.degree. C. for about 242 hours, the
applied stress was increased to about 87 MPa. After about 1 hour at
an applied stress of about 87 MPa at a temperature of about
1200.degree. C., the sample broke.
[0246] Concurrent with the stress rupture test, the strain of the
stress rupture test specimens was monitored in the gage portion of
the test specimen using the previously identified laser
extensometer to help assess the creep behavior of the fiber
reinforced ceramic matrix composite test specimen. Specifically,
the first portion of the stress rupture for Sample H was repeated.
Instead of testing the sample to failure, however, the temperature
was decreased from about 1100.degree. C. back down to about room
temperature after about 210 hours at about 1100.degree. C. under
the approximately 70 MPa applied tensile stress. FIG. 15 shows the
cumulative percent strain in the gage portion of the Sample H test
specimen resulting from this creep test. The significance of FIG.
15 is that during the course of this approximately 210 hour creep
test, Sample H shows essentially no change in elongation,
indicating substantially no plastic deformation of the sample.
Accordingly, no creep deformation of Sample H occurred under the
described test conditions.
[0247] Similarly, no creep deformation was observed in the Sample I
material which was stress rupture tested at a temperature of about
1200.degree. C. under an applied load of about 70 MPa for about 216
hours. In contrast, it has been demonstrated in the art that creep
deformation occurs in ceramic grade NICALON.RTM. silicon carbide
fibers at about 1200.degree. C. Accordingly, the present results
suggest that the present particular disposition of the reinforcing
fibers in the applied coatings and the surrounding matrix material
may provide enhanced creep resistance to the present fiber
reinforced ceramic matrix composite system.
[0248] Furthermore, the present results may suggest that the
particular disposition of the reinforcing fibers in the present
composite body provides protection to said fibers from degradation
(e.g., chemical attack) such as from atmospheric gases (e.g.,
oxygen and nitrogen) at elevated temperatures. Specifically, one
additional stress rupture test was conducted on the fiber
reinforced ceramic composite material of the present Example. The
sample was tested in substantially the same manner as was Sample I
except that after heating to a temperature of about 1200.degree.
C., the applied tensile stress was increased from about 12.5 MPa to
about 80 MPa. After about 1000 hours in air at a temperature of
about 1200.degree. C. and at a stress of about 80 MPa (as shown in
FIG. 12), the sample broke.
EXAMPLE 10
[0249] This Example demonstrates the fabrication of a ceramic grade
NICALON.RTM. silicon carbide fiber reinforced alumina matrix
composite, wherein the NICALON.RTM. fibers are first CVD coated
with dual boron nitride/silicon carbide coatings applied in
alternating layers starting with boron nitride.
[0250] A fabric preform was made by stacking 8 layers of 12 harness
satin weave (12 HSW) fabric made from NICALON.RTM. silicon carbide
fiber (ceramic grade, obtained from Dow Corning Corp., Midland,
Mich.) on top of each other substantially in accordance with the
procedure described for Sample A of Example 3.
[0251] The fabric preform comprising the 8 layers of 12 HSW
NICALON.RTM. silicon carbide fabric were then placed into the
graphite preform containment fixture 108 described in Example 2 and
depicted in FIG. 5e in substantially the same manner as was
described in Example 2. The preform containment fixture containing
the fabric preform was then placed into the reactor chamber of a
chemical vapor infiltration apparatus having an inside diameter of
about 4.5 inches (114 mm) and a length of about 18 inches (457 mm).
The warp yarns of the eighth layer of the fabric preform were
parallel to the gas flow direction within the chamber as well as
being parallel to the longitudinal axis of the reactor chamber. The
reactor chamber was closed and evacuated to less than about 0.6
torr. The reactor chamber was then heated to a temperature of about
800.degree. C. by means of inductive heating. When the temperature
within the reactor chamber reached about 800.degree. C., as
indicated by a thermocouple contained therein, a gas mixture
comprising ammonia (NH.sub.3) flowing at about 400 standard cubic
centimeters per minute (sccm) and boron trichloride (BCl.sub.3)
flowing at about 200 sccm was introduced into the reactor chamber
while maintaining a total operating pressure of about 0.6 torr.
After about 2 hours at a temperature of about 800.degree. C., the
gas mixture flowing into the reactor chamber was interrupted, the
power to the furnace heating the reactor chamber was interrupted
and the furnace and its contents were naturally cooled. After
sufficient cooling (e.g., less than about 200.degree. C.), the
reactor chamber door was opened and the preform containment fixture
was removed, cooled and disassembled to reveal that the fibers of
the fabric layers of the fabric preform were coated with boron
nitride, and furthermore, that the fabric layers comprising the
fabric preform were bonded together by the boron nitride coating.
The boron nitride coating thickness on the fibers was about 0.33
microns.
[0252] The boron nitride coated and bonded fabric preform was
stored in a vacuum desiccator pending subsequent coating.
[0253] Next, a silicon carbide coating was applied to the fibers of
the fabric preform.
[0254] The boron nitride coated and bonded fabric preform was
placed back into the reactor chamber of the above-described
chemical vapor infiltration apparatus. Because the fiber preform
was self-bonding at this stage, the graphite containment fixture
was unnecessary. The orientation of the fabric preform, however,
was substantially the same as that employed for depositing the
boron nitride coating onto the fibers in the previous deposition
reaction.
[0255] The reactor chamber door was closed and the reactor chamber
and its contents were evacuated to less than about 0.3 torr. The
reactor chamber and its contents were then heated from about room
temperature to a temperature of about 925.degree. C. at a rate of
about 50.degree. C. per minute. Hydrogen gas was then introduced
into the reactor chamber at a flow rate of about 750 standard cubic
centimeters per minute (sccm). When the reactor chamber and its
contents had equilibrated at a temperature of about 925.degree. C.,
as indicated by a thermocouple contained therein, additional
hydrogen flowing at a rate of about 750 sccm was bubbled through a
liquid bath of methyltrichlorosilane (MTS) maintained at a
temperature of about 21.degree. C., after which this gas was
introduced into the reactor chamber. The pressure in the reactor
chamber was stabilized at about 11 torr. After maintaining these
conditions of temperature, pressure and gas flow rate for about 3
hours, power to the resistance heated furnace which heated the
reactor chamber was interrupted and the about 750 sccm of hydrogen
that was being directed through the liquid MTS bath was diverted
around the MTS bath and permitted to flow directly into the reactor
chamber, thus establishing a direct hydrogen gas flow rate of about
1500 sccm into the reactor chamber. After the temperature of the
reactor chamber and its contents had dropped to about 800.degree.
C., the resistance heated furnace was re-energized and the
temperature of the reactor chamber and its contents was stabilized
at about 800.degree. C.
[0256] Another boron nitride coating was then deposited on the
coated fiber. Specifically, the flow of hydrogen gas into the
reactor was interrupted and the reactor chamber and its contents
were then evacuated to less than about 0.3 torr. Ammonia (NH.sub.3)
and borontrichloride (BCl.sub.3) gases were then introduced into
the reactor chamber in substantially the same manner as was
described previously at an operating pressure of about 0.6 torr so
as to deposit a coating of boron nitride onto the coated fibers
comprising the fabric preform. After depositing boron nitride for
about 1.5 hours at a temperature of about 800.degree. C. and at a
pressure of about 0.6 torr, the gas mixture flowing into the
reactor chamber was interrupted. The temperature of the reactor
chamber and its contents was raised from about 800.degree. C. back
up to about 925.degree. C. Hydrogen gas was then reintroduced into
the furnace chamber at a flow rate of about 750 sccm.
[0257] When the temperature of the reactor chamber and its contents
had stabilized at about 925.degree. C., a final coating of silicon
carbide was deposited onto the coated NICALON.RTM. silicon carbide
fibers comprising the fabric preform.
[0258] Specifically, substantially the same procedure was employed
in depositing this second silicon carbide coating as was employed
in depositing the first silicon carbide coating described earlier,
with the exception that the reactor chamber and its contents were
maintained at a temperature of about 925.degree. C. at an operating
pressure of about 11 torr for about 20 hours.
[0259] After depositing this second silicon carbide coating for
about 20 hours, the power to the furnace heating the reactor
chamber was interrupted and the about 750 sccm of hydrogen which
was bubbled through the liquid MTS bath was instead sent directly
into the reactor chamber without first being routed through the MTS
bath. After the furnace chamber and its contents had cooled down to
about less than about 200.degree. C., the flow of hydrogen gas into
the reactor chamber was interrupted and the reactor chamber was
evacuated to less than about 0.3 torr. The pressure in the furnace
chamber was then returned to atmospheric pressure using argon gas.
When the furnace chamber had reached substantially atmospheric
pressure, the chamber was opened and the coated fabric preform was
removed from the reactor chamber.
[0260] An alumina oxidation reaction product was grown into the
coated fiber preform in substantially the same manner as was
described for Sample A of Example 3 to form a ceramic composite
body comprising NICALON.RTM. silicon carbide fibers coated with, in
order from interior to exterior, about 0.2 micron boron nitride,
about 1.83 microns silicon carbide, about 0.2 micron boron nitride
and about 1.93 microns silicon carbide as measured along the radius
of the fiber cross-section, said coated NICALON.RTM. fibers
reinforcing a ceramic matrix, said ceramic matrix comprising an
alumina oxidation reaction product and a metallic constituent
comprising some residual unreacted parent metal.
[0261] At least a portion of the metallic component of the formed
composite body was then removed in substantially the same manner as
described in Example 2.
[0262] Flexural strength test specimens were machined and strength
tested at about room temperature in substantially the same manner
as was described in Example 2. FIGS. 16a and 16b are a scanning
electron micrographs at about 3500.times.magnification of a
polished cross-section of the fracture surface of the fiber
reinforced ceramic composite test specimen. In particular, FIG. 16a
shows a crack entering the outer silicon carbide layer and exiting
without going through the inner silicon carbide layer. Not all of
the cracks displayed this behavior, however, as evidenced by FIG.
16b which shows a crack entering through both outer and inner
silicon carbide coating layers and subsequently exiting through
both silicon carbide layers.
[0263] Demonstration that a NICALON.RTM. fiber reinforced ceramic
composite whose NICALON.RTM. fibers have coated thereon double
layers of boron nitride and silicon carbide can fracture or debond
between the inner and outer silicon carbide layers may suggest that
those fibers where this behavior occurs will be more resistant to
chemical degradation from external reactants at elevated
temperatures because such fibers are still protected by one group
of boron nitride and silicon carbide coatings.
EXAMPLE 11
[0264] This Example demonstrates that a coating of boron nitride
followed by a coating of silicon carbide on a NICALON.RTM. fiber
provide some protection from oxidation at elevated temperatures.
This Example also shows that the application of an additional set
of boron nitride and silicon carbide coatings supplied over the
first set provide significantly greater oxidation protection.
[0265] Thermogravimetric analyses were performed on Samples J, K, L
and M described below. Each test comprised placing a sample having
a mass of several tens to several hundreds of milligrams into an
alumina crucible which in turn was placed into the test chamber of
a Model SPA 409 Netzsch microbalance (Netzsch Inc., Exton, Pa.).
The chamber was sealed and substantially pure oxygen gas was
introduced into the test chamber at a flow rate of about 200
standard cubic centimeters per minute (sccm). The temperature of
the sample was then increased from substantially room temperature
to a temperature of about 1200.degree. C. at a rate of about
200.degree. C. per hour. After maintaining a temperature of about
1200.degree. C. for about 24 hours, the temperature was decreased
to about room temperature at a rate of about 200.degree. C. per
hour. The flow of the substantially pure oxygen gas was
interrupted. The microbalance continuously monitored and recorded
the mass of the test sample throughout the duration of the
test.
Sample J
[0266] Sample J comprised ceramic grade NICALON.RTM. fibers in the
"as-received" condition.
Sample K
[0267] Sample K comprised ceramic grade NICALON.RTM. fibers which
were coated with boron nitride substantially in accordance with the
method described in Example 10.
Sample L
[0268] Sample L comprised ceramic grade NICALON.RTM. fibers which
were coated with a layer of boron nitride and a layer of silicon
carbide substantially as described in Example 10.
Sample M
[0269] Sample M comprised a sample of the laminate which was
deposited on the reactor wall in accordance with the procedure of
Example 10. The laminate material specifically comprised, in
succession, layers of boron nitride, silicon carbide, additional
boron nitride and additional silicon carbide substantially as
described in Example 10.
[0270] Table IV shows the percentage weight gain for each of the
four samples as a function of the initial sample weight (e.g.,
fiber weight plus the weight of any coatings), the weight only of
initial NICALON.RTM. fiber, the weight only of the boron nitride
coating, and the weights of the silicon carbide and boron nitride
coatings. For Sample M only the percentage weight increase in terms
of the initial sample weight was measured.
[0271] The data show that coating a NICALON.RTM. fiber with both
boron nitride and silicon carbide substantially reduces the
elevated temperature oxidation of the fiber in oxygenated
environments, as evidenced by the weight increases of 0.47 and 0.65
percent, respectively, compared to the weight increase of 1.4
percent for an uncoated NICALON.RTM. fiber. Moreover, the Table
appears to indicate that the best oxidation resistance (e.g., the
least amount of weight increase) may occur when a dual duplex
coating of boron nitride and silicon carbide (e.g., four layers in
all) is applied. This result may suggest that this dual duplex
coating could not only protect a NICALON.RTM. fiber but could also
protect the underlying boron nitride/silicon carbide coatings and
in particular, the inner boron nitride debond coating.
[0272] Although only a few exemplary embodiments of the invention
have been described in detail above, those skilled in the art would
readily appreciate that the present invention embraces many
combinations and variations other than those exemplified.
EXAMPLE 12
[0273] This Example demonstrates, among other things, that the
addition of particulates of silicon carbide between the fibric
plies of fiber in a fiber reinforced ceramic composite body can
greatly reduce the extent of microcracking in the ceramic
4TABLE IV Weight Gains of coated and Uncoated NICALON .TM. Fibers
and CVI SiC/BN Coatings on Exposure to air at 1200.degree. C. for
24 hours % Weight Gain Based on Based on Based on Based on Initial
Initial CVI BN SiC/BN Sample Sample Fiber Coating Coating ID
description Wt. Wt. Wt. Wt. J NICALON .TM. 1.4 1.4 -- -- Fiber K BN
Coated 13.63 15.39 119.53 -- NICALON .TM. Fiber L SiC/BN Coated
0.47 2.0 14.81 0.65 NICALON .TM. Fiber M SiC/BN/SiC/ 0.08 -- -- --
BN Coated NICALON .TM. Fiber
[0274] matrix material which may occur between the plies during
composite fabrication.
[0275] A graphite containment fixture containing a fabric preform
was assembled in substantially the same manner as was described in
Example 2, with the following notable exceptions. First, the
NICALON.RTM. silicon carbide fabric (certamic grade obtained from
Dow Corning Corp., Midland, Mich.) entirely comprised 8-harness
satin weave (8 HSW). Moreover, about 5.75 grams of 220 grit 39
CRYSTOLON.backslash.dry silicon carbide particulate (average
particle size of about 66 microns, Norton Company, Worcester,
Mass.) was evenly applied to the top 6.75 inch (171 mm) face of
each of fabric plies 2-6. Still further, the dry silicon carbide
particulate was worked at least part way into the tows of fiber
making up each fiber ply with a brush.
[0276] The fabric preform comprising the eight stacked layers of
8-harness satin weave (8 HSW) fabric containing the five layers of
silicon carbide particulate were then placed into a chemical-vapor
infiltration (CVI) reactor and the fibers were coated with a first
layer of boron nitride (BN) followed by layer of silicon carbide
(SiC) substantially in accordance with the method described in
Example 2. As a result of chemical-vapor infiltration, about 0.51
micron of boron nitride and about 1.94 microns of silicon carbide,
as calculated based upon preform weight gain, were deposited onto
the reinforcement fibers and particulates in the preform.
[0277] A ceramic matrix comprising aluminum oxide and some aluminum
alloy metal was then grown into the coated preform by means of the
directed metal oxidation process described in Example 2 to form a
fiber reinforced ceramic matrix composite body.
[0278] The formed ceramic composite body was then subjected to
substantially the same metal removal process as described in
Example 2 to remove at least some of the metallic component of the
ceramic composite body.
[0279] The ceramic composite body was then sectioned using a
diamond saw, mounted in thermoplastic resin and polished using
progressively finer grades of diamond paste to produce a
sufficiently smooth surface for optical examination. FIG. 17b is an
approximately 50.times.magnification optical photomicrograph of
this polished cross-section of this fiber reinforced ceramic matrix
composite. Specifically, this figure shows particulates of silicon
carbide 306 embedded within an alumina ceramic matrix material 302
located between and embedding adjacent plies of fabric comprising
woven tows of the reinforcement NICALON.RTM. fibers 304. FIG. 17a
is also an approximately 50.times.optical photomicrograph of a
fiber reinforced ceramic matrix composite which was produced in
substantially the same manner as the composite material of the
present Example with the exception that no silicon carbide
particulates were placed between the fabric plies of NICALON.RTM.
fiber. The absence in FIG. 17b of the cracks 300 shown in FIG. 17a
are particularly noticeable and significant.
[0280] Thus, the present Example, among other things, demonstrates
that addition of a particulate material such as silicon carbide
between the plies of NICALON.RTM. silicon carbide fabric can
substantially reduce, if not completely eliminate, the phenomenon
of microcracking in the matrix material occupying the space between
the plies of NICALON.RTM. fabric.
EXAMPLE 13
[0281] This Example demonstrates, among other things, that there is
a preferred thickness for each of the boron nitride and silicon
carbide coatings which are applied to a preform comprising silicon
carbide reinforcement fibers if the optimal flexural strength is to
be achieved. More particularly, this Example demonstrates that for
ceramic composite bodies having about 35 to about 36 volume percent
reinforcement fibers in a matrix comprising predominantly aluminum
oxide, the optimum thickness of boron nitride is somewhere between
0.20 micron and 0.41 micron and the optimum thickness of silicon
carbide is somewhere above about 1.9 microns.
Sample N
[0282] A fabric preform was made in substantially the same manner
as was described for Sample A of Example 3, except that all eight
layers of fabric comprising ceramic grade NICALON.RTM. silicon
carbide fiber (obtained from Dow Corning Corp., Midland, Mich.)
comprised 12-harness satin weave (12 HSW) fabric. The fabric
preform was then placed into a graphite containment fixture whose
shape (but not necessarily size) was substantially as shown in FIG.
5e.
[0283] The graphite containment fixture containing the fabric
preform was then placed into a reactor chamber of a chemical-vapor
infiltration (CVI) apparatus having a outer diameter of about 4.5
inches (110 mm) and a length of about 18 inches (441 mm). The
reactor was inductively heated and oriented vertically such that
the reactive gases were introduced at the top of the reactor and
exhausted at the base of the reactor.
KANTHAL.backslash.iron-chromium-aluminum alloy wires were used to
suspend the graphite containment fixture about 11.5 inches (282 mm)
from the top of the reactor. The warp yarns of the eighth layer of
the fabric preform were parallel to the gas flow direction within
the chamber. The subsequent processing was then substantially the
same as that utilized to deposit the first boron nitride coating of
Example 10 with the exception that the reactor was maintained at
the coating temperature of about 800.degree. C. for about 135
minutes. Disassembly of the graphite containment fixture revealed
that the fibers of the fabric layers of the fabric preform were
coated and bonded together by a boron nitride coating. This boron
nitride coating had a thickness of about 0.3 micron as determined
by the weight gain of the fabric preform due to boron nitride
deposition.
[0284] Next, a silicon carbide coating was then applied on top of
the boron nitride coated and bonded fabric preform as follows. The
boron nitride coated and bonded fabric preform was suspended about
11.5 inches (282 mm) from the top of the reactor using
KANTHAL.backslash.iron-chromiu- m-aluminum alloy wires. The
orientation of the boron nitride coated and bonded fabric preform
was such that the warp yarns of the eighth layer of the 12-harness
satin weave fabric were parallel to the gas flow direction within
the chamber. The rest of the silicon carbide coating process was
substantially the same as that described in Example 10 for the
final silicon carbide coating. The boron nitride bonded fabric
preform was found to have been coated with a layer of silicon
carbide, thereby forming a silicon carbide (SiC)/boron nitride
(BN)-coated fabric preform. The silicon carbide coating had a
thickness of about 2.3 microns as determined by the weight gain of
the preform during silicon carbide deposition.
[0285] A ceramic matrix comprising aluminum oxide and some aluminum
alloy was then grown into the coated fabric preform in
substantially the same manner as was described for Sample A of
Example 3.
[0286] Unlike the composite material described in Example 2, the
composite material of the present example was neither subjected to
the metal removal process nor the elevated temperature heat
treatment.
Sample O
[0287] The Sample O composite material was prepared in
substantially the same way as described for the Sample N composite
material with the exception that the coating temperature of about
800.degree. C. for boron nitride deposition was maintained for
about 180 minutes. The resulting coating thicknesses for boron
nitride and silicon carbide were about 0.41 and about 2.30 microns,
respectively.
Sample P
[0288] The composite material of Sample P was produced in
substantially the same manner as was the material for Sample N with
the exception that the coating temperature of about 800.degree. C.
for deposition of boron nitride was maintained for about 70
minutes. The coating thicknesses of boron nitride and silicon
carbide which resulted were about 0.20 and about 2.30 microns,
respectively.
Sample Q
[0289] The composite material of Sample Q was produced in
substantially the same manner as was described for Sample N with
the exception that the temperature of about 925.degree. C. for
deposition of the silicon carbide coating was maintained for about
16 hours instead of about 20 hours. The coating thicknesses of
boron nitride and silicon carbide which resulted were about 0.29
and about 1.78 microns, respectively.
Sample R
[0290] The composite material of Sample R was produced in
substantially the same manner as was described for Sample Q with
the exception that the temperature of about 800.degree. C. for
deposition of the boron nitride coating was maintained for about 70
minutes. The coating thicknesses of boron nitride and silicon
carbide which resulted were about 0.21 and about 1.90 microns,
respectively.
[0291] Test specimens for determining the mechanical strength of
the above-described composite materials were diamond machined to
the specimen dimensions given in Example 2.
[0292] The mechanical strength of these test specimens was then
determined by stressing each specimen in four-point flexure
substantially as described in Example 2 until the sample failed.
These flexural strength tests were conducted at ambient temperature
and are reported in Table V. An examination of Table V reveals that
for a relatively constant silicon carbide coating thickness of
about 2.2 to 2.3 microns, the greatest flexural strength (taken
from an average of 7 specimens for each composite material sample)
is realized for a boron nitride coating thickness of about 0.3
micron. Reductions in strength were observed for boron nitride
coating thicknesses of 0.20 and 0.41 microns, respectively. In
addition, specifically by comparing the average strength values for
Sample Q to that of Sample N and that of Sample R to that of Sample
P, it is clear that silicon carbide coating thicknesses in the
range of 2.2-2.3 microns produce greater strengths than silicon
carbide coating thicknesses in the range of 1.8-1.9 microns.
[0293] Thus, this Example demonstrates that there exists a
desirable range of coating thicknesses of each of boron nitride and
silicon carbide to be applied to the fibers of the present
reinforced ceramic composite materials to yield optimum ambient
temperature strength. Specifically, the applied thickness of boron
nitride should be greater than about 0.20 micron but less than
about 0.41 micron. Furthermore, the applied thickness of silicon
carbide should be greater than about 1.78 microns and preferably
greater than about 1.90 microns.
5TABLE V Sample Vol. % BN* SiC* Avg. Str.** No. Fiber (.mu.m)
(.mu.m) (MPa) N 35.3 0.30 2.20 541 O 36.0 0.41 2.30 487 P 36.0 0.20
2.30 457 Q 36.0 0.29 1.78 388 R 36.0 0.21 1.90 358 *nominal coating
thickness calculated on the basis of weight gain during CVI
**average of 7 samples
EXAMPLE 14
[0294] This Example demonstrates the effect of the
chemical-vapor-infiltra- tion (CVI) coating thicknesses of both
boron nitride and silicon carbide as well as the volume fraction of
reinforcement fibers on an elevated temperature flexural strength
of a fiber reinforced ceramic composite material.
Samples S-X
[0295] Fabric preforms comprising ceramic grade NICALON.RTM.
silicon carbide fiber (obtained from Dow Corning Corp., Midland,
Mich.) was assembled in substantially the same manner as was shown
in Example 2. Unlike those of Example 2, however, the fabric
preform of the present Example only measured about 3.2 inches (77
mm) square.
[0296] The fabric preforms were then chemical-vapor-infiltrated
(CVI) with boron nitride (BN) and silicon carbide (SiC) in
substantially the same manner as was described in Example 2 with
the following notable exceptions.
[0297] Twelve fabric preforms were simultaneously coated with boron
nitride, and for the silicon carbide deposition, six boron nitride
coated preforms were simultaneously coated. Furthermore, during the
boron nitride deposition, a coating temperature of about
742.degree. C. as indicated by a thermocouple contained within the
reactor chamber was maintained for about 5.5 hours and the total
pressure within the reactor chamber was maintained between about
1.2 and about 1.6 torr. The gaseous reactants consisted of ammonia
(NH.sub.3) flowing at a rate of about 1200 standard cubic
centimeters per minute (sccm) and boron trichloride (BCl.sub.3)
flowing at a rate of about 300 sccm. For the silicon carbide
deposition, a coating temperature of about 928.degree. C. was
maintained for about 24 hours at a total reactor chamber operating
pressure of about 11 torr.
[0298] The differences in the boron nitride and silicon carbide
coating thicknesses which were obtained, as shown in Table VI, are
accounted for by the relative location of a particular preform
within the reactor. Generally speaking, for both boron nitride and
silicon carbide depositions, the closer the preform was to the gas
reactant source, the thicker was the coating. During boron nitride
deposition, the preforms were arranged six deep in groups of two.
Likewise, during silicon carbide deposition, the boron nitride
coated preforms were arranged three deep in groups of two. Thus,
during boron nitride deposition, the Sample W preform was closest
to the gas reactant source, while the Sample V preform was farthest
away. Likewise, during silicon carbide deposition, the Sample X
preform was closest to the gas reactant source, while the Sample T
preform was farthest away.
6TABLE VI Sample Vol. % BN* SiC* Avg. Str.** No. Fiber (.mu.m)
(.mu.m) (MPa) S 41 0.33 2.03 427 T 41 0.33 1.75 358 U 37 0.48 2.10
406 V 37 0.24 207 377 W 33 0.51 2.20 367 X 32 0.49 2.43 312
*nominal coating thickness calculated on the basis of weight gain
during CVI **average of 4 samples
[0299] Each fabric preform coated with boron nitride and silicon
carbide was then infiltrated with a ceramic matrix comprising
aluminum oxide and some aluminum alloy using a directed metal
oxidation process substantially as described in Example 2.
[0300] At least some of the metallic component of the formed fiber
reinforced ceramic composite bodies was then removed. This metal
removal process was substantially as described in Example 2 with
the following exceptions. A nitrogen gas flow rate of about 10,000
sccm instead of about 4000 sccm was maintained throughout the
heating and cooling. Also, three fiber reinforced composite bodies
were simultaneously processed.
[0301] Unlike the method of Example 2, however, no subsequent heat
treatment of the formed ceramic composite bodies of the present
Example was performed.
[0302] The mechanical strength of each of the fiber reinforced
ceramic composite bodies of Samples S-X was measured at a
temperature of about 1200.degree. C. in a four-point flexure mode.
Specifically, test specimens were diamond machined to the
dimensions specified in Example 2 and tested in four-point bending
until failure using the procedure as outlined in Example 2. The
resulting flexural strength based upon an average of four data
points per material is reported for each of Samples S-X in Table
VI.
[0303] An examination of the data reported in Table VI reveals a
number of trends.
[0304] Specifically, a comparison of the data for Sample S with
those for Samples U, V and W shows that boron nitride coating
thicknesses of about 0.24 micron and 0.48 micron or thicker
resulted in flexural strengths which were suboptimal. Thus, for the
fiber reinforced composite bodies of the present invention, a boron
nitride coating thickness on the reinforcement fibers of about 0.3
micron ought to be at least close to optimal.
[0305] Moreover, inspection of Table VI reveals a correlation
between the thickness of the boron nitride coating and the volume
fraction of the NICALON.RTM. reinforcement fiber in the coated
preform (and ultimately the composite body). This correlation
results from the "spring-back" phenomenon which occurs after
coating the fabric preforms with boron nitride when the graphite
containment fixture is disassembled and the boron nitride coated
fabric preform is removed. It seems that the thicker the boron
nitride coating which is applied, the more the coated fabric
preform expands in its thickness dimension upon its removal from
the graphite containment fixture, thus the lower the overall volume
fraction of NICALON.RTM. fiber.
[0306] In addition, a comparison of the strength data for Samples
S, U and W shows that the flexural strength of the composite
material increases as the volume fraction of reinforcement fibers
increases.
[0307] Finally, a comparison of the data for Sample T with the data
for Sample S and, similarly, a comparison of the data for Sample X
with the data for Sample W reveals that for a given volume fraction
of reinforcement fibers, a silicon carbide thickness of about 1.75
microns, while a thickness of about 2.43 microns of silicon carbide
is excessive in that it also results in suboptimal strength. The
importance of the proper silicon carbide thickness is highlighted
by a comparison of the data for Sample T with for Sample W.
Specifically, this comparison shows that the optimization of the
silicon carbide coating thickness at about 2.2 microns compensates
for a difference in volume fraction of reinforcement fibers of
about eight points or about 20 percent.
[0308] Thus, this Example demonstrates that for the fiber
reinforced ceramic composite materials described herein, there
exists desirable boron nitride and silicon carbide coating
thicknesses to apply to the reinforcement fibers to optimize the
high temperature strength of the formed composite materials.
Specifically, the thickness of boron nitride applied should be
greater than about 0.24 micron but less than about 0.51 micron. The
thickness of silicon carbide applied should be greater than about
1.75 microns but less than about 2.43 microns. These desired
thickness ranges agree well with the optional ranges of Example 13
derived for optional ambient temperatures composite strength.
Furthermore, within the range of fiber reinforcement of about 32 to
41 volume percent of the composite body, the strength of the
composite body increases with the volume fraction of
reinforcement.
[0309] Table VII summarizes many of the processing parameters
followed in fabricating the fiber reinforced ceramic composite
bodies described in the foregoing Examples.
EXAMPLE 15
[0310] This Example demonstrates, among other things, an improved
method of coating a fabric preform. Specifically, this Example
demonstrates a set of coating conditions which result in coatings
of more uniform thickness throughout the fabric preform.
[0311] A fabric preform 103 was made by stacking a plurality of
layers of 8 harness satin weave (8 HSW) fabric and 12 harness satin
weave (12 HSW) fabric made from ceramic grade NICALON.RTM. silicon
carbide fiber (obtained from Dow Corning Corp., Midland, Mich.) on
top of each other in substantially the same manner as was described
in Example 8. The fabric preform had dimensions of about 9 inches
(229 mm) long by about 6 inches (152 mm) wide by about 0.125 inch
(3.2 mm) thick.
[0312] The fabric preform was clamped in substantially the same
kind of fixture as was described in Example 2 and depicted in FIG.
5e. The preform containment fixture 108 containing the fabric
preform was placed into a reactor chamber of a refractory alloy
steel chemical vapor infiltration (CVI) apparatus having a graphite
tube liner and having overall dimensions of about 8 feet (2.4
meters) in length by about 15.5 inches (394 mm) in inside diameter.
The warp yarns of the first and seventh layers of the fabric
preform were perpendicular to the gas flow direction within the
chamber as well as being perpendicular to the longitudinal axis of
the reactor chamber. The reactor chamber was closed and evacuated
to less than about 0.5 torr. The reactor chamber was then heated to
a temperature of about 800.degree. C. When the temperature within
the reactor TABLE VII
7 TABLE VII Matrix Fabrication Parameters Residual Metal Example
Sample ID Fiber Plies Atmosphere Temperature Time (h) Removed
Comments 1 -- 3 bundles of 500 fibers O.sub.2 1000.degree. C. 54 no
2 -- 12 HSW-(6) 8 HSW-12 HSW air 950.degree. C. 90 yes Some heat
treated (see Table I); parent metal containing 8.5-11 wt % Si 3 A
(8) 12 HSW air 1000.degree. C. 60 yes 3 B (4) 12 HSW N.sub.2
900.degree. C. 200 yes 4 C 12 HSW-(6) 8 HSW-12 HSW air 950.degree.
C. 90 no 4 D 12 HSW-(6) 8 HSW-12 HSW air 950.degree. C. 90 yes 4 E
12 HSW-(6) 8 HSW-12 HSW air 950.degree. C. 90 yes 4 F 12 HSW-(6) 8
HSW-12 HSW air 950.degree. C. 90 yes 5 -- 12 HSW-(6) 8 HSW-12 HSW
air 950.degree. C. 90 yes 6 -- 12 HSW-(6) 8 HSW-12 HSW air
950.degree. C. 90 yes 7 -- 12 HSW-(6) 8 HSW-12 HSW air 950.degree.
C. 90 yes 8 -- (2) 8 HSW-(3) 12 HSW-(2) 8 HSW air 950.degree. C. 72
yes 8 -- (2) 8 HSW-(3) 12 HSW-(2) 8 HSW air 950.degree. C. 100 yes
9 G,H,I (8) 12 HSW air 950.degree. C. 90 yes 10 -- (8) 12 HSW air
1000.degree. C. 60 yes 12 -- (8) 8 HSW air 950.degree. C. 90 yes 5
layers of SiC particulate 13 N-R (8) 12 HSW air 1000.degree. C. 60
no 14 S-X 12 HSW-(6) 8 HSW-12 HSW air 950.degree. C. 90 yes
[0313] chamber reached about 800.degree. C., a gas mixture
comprising borontrichloride (BCl.sub.3) flowing at about 1200 sccm
at a temperature of about 60.degree. C. and ammonia (NH.sub.3)
flowing at about 2100 sccm was introduced into the reactor chamber
while maintaining a total operating pressure of about 0.5 torr.
After about 4 hours at a temperature of about 800.degree. C., the
gas mixture flowing into the reactor chamber was interrupted, the
power to the furnace beating the reactor chamber was interrupted
and the furnace and its contents were naturally cooled. At a
temperature below about 200.degree. C., the reactor chamber door
was opened and the graphite containment fixture was removed, cooled
and disassembled to reveal that the fibers of the fabric layers of
the fabric preform were coated and that the fabric layers
comprising the fabric preform were bonded together by a boron
nitride coating. The boron nitride coating had a thickness of about
0.48 micron.
[0314] The boron nitride coated fabric preform was then stored in a
vacuum desiccator until it was ready to be put back into the
chemical vapor infiltration apparatus for additional coating.
[0315] For the application of this subsequent coating, the boron
nitride coated and bonded fabric preform was placed back into the
reactor chamber of the chemical vapor infiltration apparatus. In
this instance, however, the warp yarns of the first and seventh
layers of the fabric preform were parallel to the gas flow
direction within the chamber, as well as being parallel to the
longitudinal axis of the reactor chamber. More specifically, the
boron nitride coated fabric preforms were supported by a graphite
fixture as shown in FIG. 18A. The graphite fixture alone is shown
in FIG. 18B. A total of 8 boron nitride coated fabric preforms can
be further coated simultaneously in a single reactor run by placing
2 such loaded fixtures front-to-back in the reactor chamber.
[0316] The CVI reactor chamber was closed and evacuated to about
less than about 1 torr. Hydrogen gas was introduced into the
reactor chamber at a flow rate of about 11,000 standard cubic
centimeters per minute (sccm). The reactor chamber was then heated
to a temperature of about 950.degree. C. The reactor pressure was
equilibrated at about 250 torr. Once the temperature of the
contents of the reactor chamber had substantially completely
stabilized at about 950.degree. C., about 1800 sccm of hydrogen
were diverted away from direct entry into the reactor chamber and
were first bubbled through a bath of methyltrichlorosilane (MTS)
maintained at a temperature of about 45.degree. C. before entering
the reactor chamber. After about 48 hours at a temperature of about
950.degree. C., the power to the furnace heating the reactor
chamber was interrupted and the about 1800 sccm of hydrogen that
was being directed through the MTS bath was again permitted to flow
directly into the reactor chamber to reestablish a direct hydrogen
gas flow rate of about 11000 sccm into the reactor chamber. After
the reactor chamber had cooled substantially, the hydrogen flow
rate was interrupted and the furnace and its contents were
evacuated to less than 1 torr. The pressure within the reactor
chamber was then brought back up to about atmospheric pressure with
argon gas. After the reactor chamber had cooled to a temperature
below about 200.degree. C., the argon gas flow rate was interrupted
and the reactor chamber door was opened. The graphite support
fixtures were removed, cooled and disassembled to reveal that the
boron nitride bonded fabric preforms had been coated with a second
layer of silicon carbide thereby forming a silicon carbide
(SiC)/boron nitride (BN)-coated fabric preform. The silicon carbide
had a thickness of about 2-3 microns. Significantly, the silicon
carbide coating was of more uniform thickness from the interior of
the preform to an exterior surface in the present Example than in
the previously described Examples. In other words, the thickness of
silicon carbide deposited at the exterior of the preform was not as
great as in earlier Examples; thus, the coated preforms of the
present Example were more permeable than some of the coated
preforms of the previous Examples. Thus, the results of this
Example suggest that it may be possible to apply silicon carbide
coatings to the present fabric preforms having nominal thickness
greater than about 2 to 3 microns without creating isolated pores
in the preform.
[0317] Growth of an alumina oxidation reaction product through the
silicon carbide/boron nitride-coated fabric preform was then
carried out in substantially the same manner as was described in
Example 2 to form a fiber reinforced ceramic composite body
comprising a ceramic matrix comprising an aluminum oxide oxidation
reaction product and a metallic component comprising some residual
unreacted parent metal, with the ceramic matrix embedding the
silicon carbide/boron nitride coated NICALON.RTM. silicon carbide
fibers. Substantially complete growth of the ceramic matrix only
required about 72 hours, however. Because of the more permeable
nature of the coated fabric preforms of the present Example, it is
believed that the time required for complete growth is even less
than this value.
[0318] Thus, this Example demonstrates an efficient technique for
coating a plurality of preforms simultaneously as well as
conditions which result in silicon carbide coatings of more uniform
thickness.
EXAMPLE 16
[0319] This Example demonstrates the fatigue characteristics of the
present fiber reinforced ceramic composite materials. Specifically,
this Example demonstrates the lifetimes for samples of ceramic
grade NICALON.RTM. silicon carbide fiber reinforced alumina matrix
composite bodies as a function of the maximum applied stress for
bodies tested in air at various temperatures and subjected to
low-frequency cycling in tension.
[0320] Samples which were tested at about 20.degree. C. and at
about 1000.degree. C. were fabricated in substantially the same
manner as was described in Example 2, including the residual metal
removal process. The fiber reinforced ceramic composite bodies
which were tested at temperatures of about 1100.degree. C. and
about 1370.degree. C. were fabricated substantially in accordance
with Example 8, which fabrication also included the residual metal
removal process.
[0321] The geometry of the test specimen was that of a "double
dogbone", that is, similar to that geometry in FIG. 7 except
further comprising another reduced section. Specifically, test
specimens were diamond machined to an overall length of about 5
inches (127 mm), about 0.55 inch (14 mm) maximum width and having a
gage section measuring about 1.3 inches (33 mm) in length by about
0.25 inch (6 mm) in width.
[0322] Each sample was placed into the test chamber of a universal
testing machine at about 25.degree. C. The test chamber was then
heated to the desired elevated temperature in air. When the
temperature of the specimen had stabilized, a sinusoidal tensile
stress was applied to the specimen. The minimum applied tensile
stress was about 10 percent of the maximum applied tensile stress.
The testing apparatus was configured so as to record the number of
tensile stress cycles required to cause failure. These fatigue data
are illustrated in Table VIII. Those test specimens which were
still intact following 10,000 cycles of applied tensile stress at
temperatures of 1100.degree. C. or 1370.degree. C. were then
tensile tested in air using a uniformly increasing load at the same
temperature at which they were cycled in applied stress until
failure was observed. The table shows that such test specimens
retained over 50 percent of their original strength following the
tensile cycling at elevated temperature. The table also
demonstrates that the fiber reinforced ceramic composite material
was capable of surviving over 2 million cycles of a tensile stress
applied between about 8 and about 83 MPa at a frequency of about 5
Hz at a temperature of about 1000.degree. C. in air. The data
generated at the elevated temperatures are presented graphically in
FIG. 19 which shows the number of cycles to produce failure as a
function of the maximum applied tensile stress. The arrows
connected to several of the data points and pointing to the right
indicate that the particular data point represents a lower bound of
the material's life (e.g., failure of the specimen had not been
achieved at the indicated maximum stress and number of tensile
cycles). Further examination of the data presented in FIG. 19
suggest an endurance limit for the fiber reinforced ceramic
composite material of about 80 MPa at a temperature of about
1000.degree. C. in air.
8TABLE VIII 2-D Nicalon .TM. /Al.sub.2O.sub.3 Low Cycle Fatigue
Data Test Max. Stress Residual Str. Temper- (R = 0.1) Frequency
Cycles at Temp. ature MPa (ksi) Hz to failure MPa (ksi) Room 72
(11) 1 172,912* -- Temper- 72 (11) 1 235,858* -- ature 138 (20) 1
150,956* -- 138 (20) 1 166,459* -- 166 (24) 1/10 351,600/118,360 --
1000.degree. C. 83 (12) 5 2,195,000* -- (1800.degree. F.) 103 (15)
0.3 142,577* -- 103 (15) 5 117,843 -- 103 (15) 5 175,620 -- 124
(18) 1 27,712 -- 124 (18) 5 66,676 -- 1100.degree. C. 60 (9) 1
10,000* 185 (26) (2000.degree. F.) 120 (17) 1 10,000* 149 (21) 180
(26) 1 603 -- 1370.degree. C. 58 (8) 1 10,000* 153 (22)
(2500.degree. F.) 115 (16) 1 2,480 -- 173 (25) 1 128 -- *Test
stopped prior to failure RT & 1000.degree. C. data generated by
GE 1100 & 1370.degree. C. data generated by Williams
International
EXAMPLE 17
[0323] This Example illustrates a modified stress rupture test
whose purpose or objective was to further simulate at least some of
the conditions which might be present in a turbine engine.
Specifically, the test illustrated by the present example is
similar in many respects to the stress rupture test described in
Example 9 with the exception that a temperature cycle was added or
superimposed to the test system during the elevated temperature
exposure under the applied dead load.
[0324] A fiber reinforced ceramic composite body was fabricated in
substantially the same manner as was described in Example 2,
including removing a substantial fraction of the residual metallic
component.
[0325] Tensile test specimens were diamond machined in
substantially the same manner as was described in Examples 7 and 9
and loaded into the test fixture described in Example 9, which in
turn was loaded into the Model P-5 creep testing machine (SATEC
Inc., Grove City, Pa.). A tensile stress of about 12.5 MPa was then
applied to the test specimen using dead loading. A
resistance-heated air atmosphere furnace was positioned completely
around the test fixture portion of the creep testing machine and
the furnace and the sample contained within were heated from about
20.degree. C. to a temperature of about 1100.degree. C. Each
thermal cycle then consisted of maintaining a temperature of about
1100.degree. C. for about 1 hour in air, then uniformly decreasing
the temperature of the test specimen to a temperature of about
600.degree. C. over a period of about 45 minutes, maintaining a
temperature of about 600.degree. C. for about 1 hour and finally
uniformly increasing the temperature of the specimen back up to a
temperature of about 1100.degree. C. over a period of about 45
minutes. As in Example 9, sample strain was monitored with a Model
1102 ZYGO.TM. helium-neon laser extensometer (Zygo Corporation,
Middlefield, Conn.). The sample test data were then recorded in the
form of sample strain as a function of test duration, which data
are illustrated graphically in FIG. 20. Referring to FIG. 20, after
the application of about 12 thermal cycles at an applied tensile
stress of about 25 MPa (the first few cycles being used to check
out the functioning of all of the test equipment), the stress was
then increased to about 50 MPa. After the application of about 115
thermal cycles at a stress of about 50 MPa, the stress was then
further increased to about 70 MPa. After about 22 cycles at a
stress of about 70 MPa, the sample failed. FIG. 20 also shows an
enlargement of the 70 MPa region which specifically illustrates the
change in strain in the test specimen in response to the change in
specimen temperature. All together, the test specimen survived a
total of 142 thermal cycles or about 500 hours of test
duration.
[0326] A second thermal cycling test was then conducted on a
substantially similar fiber reinforced composite test specimen in
substantially the same manner as described above with the exception
that the applied tensile stress throughout the duration of the
testing was about 50 MPa. This second test specimen survived
thermal cycling under this applied tensile stress in air for a
total of about 282 thermal cycles (or about 987 hours) before
failure occurred.
[0327] Thus, this Example demonstrates that the present fiber
reinforced ceramic composite materials are capable of surviving
hundreds of hours under tensile loads in air under conditions of
varying elevated temperature.
EXAMPLE 18
[0328] This Example demonstrates, among other things, the
deposition of a chemically modified coating layer on a
reinforcement filler material and subsequent encapsulation by a
matrix material to form a composite body. More specifically, the
present Example demonstrates the incorporation of a source of
silicon into a boron nitride based coating material.
[0329] An artisan of ordinary skill will readily appreciate that
numerous modifications may be made to the above-identified Examples
without departing from the spirit of the present invention.
Accordingly, the Examples should be considered as illustrative of
the invention and should in no way be construed as limiting to the
scope of the invention as defined in the claims appended
hereto.
Sample AA (no silicon doping)
[0330] A fabric preform comprising ceramic grade Nicalon.RTM.
silicon carbide fiber was fabricated in substantially the same
manner as described in Example 2. The fabric preform of the present
Example measured about 3 inches (76 mm) square by about 0.125 inch
(3 mm) thick.
[0331] The fabric preform was then coated with a material
comprising boron nitride in substantially the same manner as in
Example 2 except that the usable inside diameter of the coating
chamber was about 5 inches (127 mm) instead of about 9.45 inches
(240 mm), the temperature inside the reactor was maintained at
about 736-740.degree. C., the operating pressure was maintained at
about 1.1 to 1.2 Torr, the time at temperature was about 4 hours,
the flowrate of the ammonia reactant was about 342 standard cubic
centimeters per minute (sccm), and the flowrate of the boron
trichloride was about 85 sccm. An average of about 0.32 micron of
boron nitride was deposited on the Nicalon.RTM. silicon carbide
filaments, as calculated from the weight gain of the fabric
preform.
[0332] Next, a coating comprising silicon carbide was deposited on
top of the boron nitide coated fabric preform. The deposition
conditions were substantially the same as those described for the
silicon carbide deposition of Example 2 except that the coating was
deposited at a temperature of about 980.degree. C., at a pressure
of about 11 Torr and and for a duration of about 19 hours. Based
upon the preform weight gain, it was estimated that the average
silicon carbide coating thickness on the fibers was about 1.32
microns.
[0333] Following this CVD coating process, a matrix comprising
aluminum oxide was formed by directed metal oxidation of a parent
metal comprising aluminum. More specifically, the matrix was formed
in substantially the same manner as the matrix described in Example
2 with the exception that the nickel oxide particulate was about
minus 200 mesh (substantially all particles smaller than about 75
microns) and the dwell temperature of about 950.degree. C. was
maintained for about 96 hours.
[0334] Following matrix formation, most of the residual, unreacted
parent metal within the formed alumina matrix composite was
removed. The metal removal technique was substantially the same as
that described in Example 2 except that the metal removal lay-up
was loosely covered with a Grafoil.RTM. graphite foil lid (Union
Carbide Co., Cleveland, Ohio), and that the filler material mixture
for infiltration comprised by weight about 10 percent ground
magnesium particulate (-100+200 mesh, Hart Metals, Tamaqua, Pa.)
and the balance grade C75 unground alumina particulate (Alcan
Chemicals Div. of Alcan Aluminum Corp., Cleveland, Ohio). A weight
loss in the composite body of about 8.4 percent was recorded.
[0335] Unlike some of the flexural test specimens in Example 2, the
composite body of the present Example was not heat treated.
Sample AB
[0336] Sample AB features an attempt to deliberately add a source
of silicon to the CVD gas stream for the purposes of doping the
resulting boron nitride based coated with silicon.
[0337] Sample AB was fabricated by substantially the same
procedures as was Sample AA with the following exceptions: The
boron nitride coating run was conducted at a pressure of about 1.6
to 1.7 Torr. Also, to the NH.sub.3 and BCl.sub.3 gas streams was
added about 200 sccm of hydrogen (H.sub.2) and silicon
tetrachloride (SiCl.sub.4). The H.sub.2 was bubbled through liquid
SiCl.sub.4 at ambient (e.g., about 20.degree. C.) temperature,
thereby acting as a carrier gas for SiCl.sub.4. About 0.32 micron
of the BN based coating was applied to the fibers of NICALON.RTM.
silicon carbide, based upon the weight gain of the fabric
preform.
[0338] During the silicon carbide deposition by CVD, the dwell at
about 980.degree. C. was maintained for about 11.5 hours. About
1.79 microns of silicon carbide were deposited.
Sample AC
[0339] Sample AC was prepared in substantially the same manner as
was Sample AB with the following exceptions: The boron nitride
based coating was deposited for about 5 hours at a temperature of
about 735.degree. C. to 740.degree. C. and at a pressure of 1.3-1.4
Torr. The BCl.sub.3 and NH.sub.3 gas flowrates were about 65 sccm
and 262 sccm, respectively. The SiCl.sub.4 with its H.sub.2 carrier
was admitted to the coater at a flowrate of about 496 sccm. A
coating about 0.42 micron thick was deposited. The silicon carbide
coating was deposited identically to that described for Sample
AA.
Sample AD
[0340] Sample AD is also a silicon carbide fiber reinforced alumina
matrix composite material wherein the fibers are coated with a
duplex coating featuring a boron nitride containing layer followed
by a silicon carbide coating layer. As with Samples AB and AC, the
BN layer of the present sample was also modified, but by a
different route. Instead of the SiCl.sub.4 being carried into the
reactor by H.sub.2, however, pure SiCl.sub.4 gas with no carrier
was employed. In particular, it was realized that the vapor
pressure of SiCl.sub.4 at ambient temperature was sufficient to
produce a flowrate of about 21 sccm under the process conditions.
The BCl.sub.3 and NH.sub.3 gas flowrates were 85 sccm and 342 sccm,
respectively. The modified BN layer was deposited for about 5 hours
at a pressure of about 1.3 to 1.4 Torr. An approximately 0.37
micron thick coating was deposited.
[0341] The SiC deposition was identical to that in Sample AA.
Sample AE
[0342] Sample AE was prepared almost identically to Sample AD. The
only significant 6difference between these two Samples is that for
the BN based coating, the present Sample featured gas flowrates of
60 sccm, 69 sccm and 276 sccm for the SiCl.sub.4, BCl.sub.3 and
NH.sub.3, respectively. An approximately 0.35 micron thick coating
was deposited.
Sample AF
[0343] Sample AF was prepared in substantially the same manner as
was Sample AE except that the gas flowrates for the modified BN
deposition were 40 sccm, 82 sccm and 326 sccm of SiCl.sub.4,
BCl.sub.3 and NH.sub.3, respectively. The thickness of the
resulting coating was about 0.35 micron.
[0344] A chemical elemental analysis was performed on the modified
BN coatings on some of the Samples. This elemental analysis was
performed by an outside contractor using scanning Auger depth
profiling. In addition to the expected B, N and Si, considerable
quantities of C and O were also detected. The results of this
elemental analysis are reported in terms of atomic percent in Table
IX. The column labeled "Si:B ratio" provides the ratio of silicon
to boron atoms in the precursor SiCl.sub.4 and BCl.sub.3 gases.
[0345] To gauge the effect of the attempted silicon modification of
the boron nitride coating on the strength of the composite bodies
into which the coated were encapsulated, a number of flexural
strength sample test bars were diamond machined from each Sample.
The four point flexural strength testing was performed at ambient
temperature and at about 1200.degree. C. in air substantially as
described in Example 2. The mean flexural strength based on a
sample size of 3 is reported in Table X, as well as presented in
graphical form in FIG. 21. The flexural strength shows little
dependence upon the ratios of the boron and silicon precursor
gases.
[0346] To assess the resistance of the present composite bodies to
chemical attack by oxygen and moisture at elevated temperatures,
three of the present Samples were selected for such corrosion
testing. Specifically, Samples AD, AE and AF were subjected to an
atmosphere consisting of 90 percent by volume water vapor, balance
oxygen at a temperature of about 800.degree. C. and ambient
pressure. After an approximately 88 hour exposure, the recession
distance (e.g., corrosion length) was measured from a machined
surface of the composite body. This recession distance was about
600 microns to 800 microns for Sample AE and about 500 microns to
about 700 microns for Sample AF, but only about 2 microns to 40
microns for Sample AD. By comparison, a typical Si/SiC matrix
composite body featuring an unmodified boron nitride fiber coating
exhibits recession distances of about 1700 microns to about 2200
microns. Although the matrix material is not thought to
significantly affect the recession rate, since all test samples
feature exposed filament ends, if anything, the alumina matrix
composites of Sample AD, AE and AF might be expected to corrode
faster than the composite body formed by melt infiltration since
the alumina matrx has a higher population density of microcracks.
The fact that these Samples exhibited less corrosion (less
recession distance) suggests that the modified BN coating on the
fibers is more chemically protective than regular, unmodified
BN.
9 TABLE IX Si: B Atom Elemental Analysis, Atom % Sample ratio
(precursor gases) B N Si C O AA 0 AB 2.3 AC 7.7 39 38 0 18 5 AD
0.25 39 41 1-2 13 5 AE 0.87 37 42 3 10 10 AF 0.49 38 40 3 13 6
[0347]
10 TABLE X Four Point Flexural Si:B Atom Strength (MPa) Sample
ratio (precursor gases) at 20.degree. C. at 1200.degree. C. AA 0
483 +/- 51 374 +/- 15 AB 2.3 328 +/- 29 251 +/- 5 AC 7.7 455 +/- 16
311 +/- 31 AD 0.25 429 +/- 20 307 +/- 21 AE 0.87 360 +/- 45 320 +/-
36 AF 0.49 427 +/- 53 362 +/- 10
* * * * *