Note: Descriptions are shown in the official language in which they were submitted.
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Toughness Enhanced Silicon-containing Composite Bodies,
and Methods for Making Same
'This patent document is a Continuation-in-Part of U.S. Patent Application
Serial
No. 09/990,175, filed on November 20, 2001, which application claims the
benefit of
U.S. provisional patent application serial numbers 60/329,358, and 60/252,489,
filed on
October 15, 2001, and November 21, 2000, respectively.
TECHNICAL FIELD
This invention relates to toughened silicon-ceramic composite bodies,
especially
those produced by a reactive infiltration process, e.g., reaction-bonded
bodies. Reaction-
bonded silicon carbide having a boron carbide filler or reinforcement
particularly
exemplifies the invention.
BACKGROUND ART
Reaction bonded silicon carbide (SiC) ceramics combine the advantageous
properties of high performance traditional ceramics, with the cost
effectiveness of net
shape processing. These materials provide high surface hardness, very high
specific
stiffness, high thermal conductivity, and very low coefficient of thermal
expansion
(CTE). The processing consists of two steps. First, a carbon containing near
net shape
porous preform is fabricated; and second, the preform is reactively
infiltrated with molten
Si to form a primarily SiC body.
Reaction bonded silicon carbide ceramic offers extremely high levels of
mechanical and thermal stability. It possesses low density (similar to A1
alloys) and very
high stiffness (~70% greater than steel). These properties lead to components
that show
little deflection under load, allow small distances to be precisely controlled
with fast
machine motion, and do not possess unwanted low frequency resonant vibrations.
In
addition, due to the high stiffness and hardness of the material, components
can be ground
and lapped to meet stringent flatness requirements. Moreover, as a result of
very low
coefficient of thermal expansion (CTE) and high thermal conductivity, reaction
bonded
SiC components show little motion with temperature changes, and are resistant
to
distortion if localized heating occurs. Also, due to an excellent CTE match
with Si,
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reaction bonded SiC ceramics are well suited as substrates for Si handling
operations.
Furthermore, both Si and SiC possess refractory properties, which yields a
composite
with good performance in many high temperature and thermal shock applications.
Finally, dense, high purity SiC coatings can be applied when extremely high
purity and/or
superior resistance to corrosion are required.
Reaction bonded SiC ceramics have many outstanding properties, including high
specific stiffiiess, low coefficient of thermal expansion, and high thermal
conductivity.
However, they typically also exhibit low fracture toughness, and are therefore
not suited
for many applications where impact loading will occur.
Thus, materials investigators have experimented with various techniques for
enhancing the toughness or impact resistance of such inherently brittle
ceramic-rich
materials. Perhaps the most popular approach has been to incorporate fibrous
reinforcements and attempt to achieve crack deflection or fiber debonding and
pull-out
mechanisms during the crack propagation process.
More recently, some have alloyed the brittle silicon phase with different
metals
such as aluminum, to enhance toughness. For example, the assignee's aluminum-
toughened SiC provides a nominally 75.°!° increase in fracture
toughness relative to its
standard reaction bonded SiC. This toughness allows the composite to be used
in
applications where some impact will occur. In addition, the composite can be
used in thin
walled component designs that would be difficult to produce with a low
toughness
ceramic.
The presence of the aluminum results in an increase in thermal conductivity
relative to the standard SiC product, which is valuable in heat sink
applications or in
components where localized heating can occur. In addition, the thermal
conductivity is in
excess of that of most metal matrix composites because no additional metallic
alloying
elements are used, e.g., magnesium.
Silicon is usually thought of as being a brittle material, but this statement
pertains
to its ambient temperature characteristics. A review of the literature finds
that Si
undergoes a brittle to ductile transition in the 500°C temperature
range. (J. Samueles,
S.G. Roberts, and P.B. Hirsch, "The Brittle-to-Ductile Transition in Silicon,"
Materials
Science and En~ineerin~, A105/106 (1988), pp. 39-36.) Warren observed that
this
temperature can be influenced according to the density of dislocations in the
silicon.
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Specifically, he observed that the transition temperature decreased when
dislocations
were introduced to a silicon surface by a grinding operation. (P.D. Warren,
"The Brittle-
Ductile Transition itz Silicon: The Influence of Pre-Existing Dislocation
Arrangettzents,"
Scripta Met., 23 (1989), pp. 637-42.)
Moreover, Gogotsi et al. of Drexel have achieved metal-like ductile machining
of
silicon at room temperature by controlling the stress state at the cutting
tool-to-workpiece
interface. (Y. Gogotsi, C. Baek, and F. I~irscht, "Ranzan nzicrospectroscopy
study of
processing-induced phase trartsforrtzations and residual stress state itt
silicon",
Semicond. Sci. Tech. 14, (1999), pp. 936-44). Analysis of the Silicon crystal
structure
directly under the cutting tool by Raman microspectroscopy shows that the room-
temperature ductile machining is obtained through a pressure-induced
transformation of
Silicon from the cubic diamond phase into a metallic beta-tin structure. This
latter phase
has mechanical properties of a typical metal and deforms plastically under the
tool. Thus,
under the correct loading conditions, ductile behavior of Silicon at room
temperature can
be obtained. Phase transformations in Silicon, as well as in boron carbide,
can lead to
additional energy dissipation, which is important in many applications,
possibly including
armor applications.
U. S. Patent No. 3,857,744 to Moss discloses a method for manufacturing
composite articles comprising boron carbide. Specifically, a compact
comprising a
uniform mixture of boron carbide particulate and a temporary binder is cold
pressed.
Moss states that the size of the boron carbide particulate is not critical;
that any size
ranging from 600 grit to 120 grit may be used. The compact is heated to a
temperature in
the range of about 1450°C to about 1550°C, where it is
infiltrated by molten silicon. The
binder is removed in the early stages of the heating operation. The silicon
impregnated
boron carbide body may then be bonded to an organic resin backing material to
produce
an armor plate.
U.S. Patent No. 3,859,399 to Bailey discloses infiltrating a compact
comprising
titanium diboride and boron carbide with molten silicon at a temperature of
about
1475°C. The compact further comprises a temporary binder that,
optionally, is
carbonizable. Although the titanium diboride remains substantially unaffected,
the
molten silicon reacts with at least some of the boron carbide to produce some
silicon
carbide in situ. The flexural strength of the resulting composite body was
relatively
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modest at about 140 MPa. A variety of applications is disclosed, including
personnel,
vehicular and aircraft armor.
U.S. Patent No. 3,796,564 to Taylor et al., filed in 1967, discloses a hard,
dense
carbide composite ceramic material particularly intended as ceramic armor.
Granular
boron carbide is mixed with a binder, shaped as a preform, and rigidized. Then
the
preform is thermally processed in an inert atmosphere with a controlled amount
of molten
silicon in a temperature range of about 1500°C to about 2200°C,
whereupon the molten
silicon infiltrates the preform and reacts with some of the boron carbide. The
formed
body comprises boron carbide, silicon carbide and silicon. Taylor et al. state
that such
composite bodies may be quite suitable as armor for protection against low
caliber, low
velocity projectiles, even if they lack the optimum properties required for
protection
against high caliber, high velocity projectiles. Although they desire a
certain amount of
reaction of the boron carbide phase, they also recognize that excessive
reaction often
causes cracking of the body, and they accordingly recognize that excessive
processing
temperatures and excessively fine-grained boron carbide is harmful in this
regard. At the
same time, they also realize that excessively large-sized grains reduce
strength and
degrade ballistic performance.
More recently, the assignee has likewise developed a reaction bonded boron
carbide composite material because this system was thought to show potential
as a
candidate armor material, even if previous efforts suffered from various
shortcomings.
As examples of this "potential", hardness is believed to be important in
making an armor
material having high mass efficiency. Moreover, many armor applications such
as
aircraft and body armor, require low mass. Boron carbide possesses both of
these
characteristics. None of the individual components of this reaction bonded
boron carbide
system possesses inherent toughness.
U.S. Patent No. 3,725,015 to Weaver discloses a process for making low
porosity,
essentially defect free, composite refractory shapes via a reactive
infiltration process. A
carbon-containing preform is infiltrated with a molten metal containing at
least two
constituents. One of the constituents is capable of reacting with the carbon
to form a
metal carbide in situ in the preform. The other constituent is added such that
the infiltrant
alloy has a thermal expansion close to that of the refractory material making
up the matrix
of the preformed shape, thereby regulating, and preferably eliminating
residual stress and
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microcracking upon cooling to ambient temperature form the processing
temperatures.
Weaver furthermore discloses providing to the infiltrant alloy a metal
corresponding to
the metal of the refractory material for the purpose of preventing the
infiltrant from
leaching out of the metal constituent of the refractory material. For example,
the
incorporation of about 6 percent by volume of boron in silicon saturates the
alloy
sufficiently to prevent its dissolving boron out of a boron carbide refractory
material
matrix.
Commonly Owned International Patent Application No. PCT/LJS99/16449, filed
on July 23, 1999, and which published as Publication No. WO 01/07377 on
February 1,
2001, teaches that reaction-bonded or reaction-formed silicon carbide bodies
may be
formed using an infiltrant comprising silicon plus at least one metal, e.g.,
aluminum.
Modifying the silicon phase in this way permits tailoring of the physical
properties of the
resulting composite, and other important processing phenomena result: Such
silicon
carbide composite materials are of interest in the precision equipment,
robotics, tooling,
armor, electronic packaging and thermal management, and semiconductor
fabrication
industries, among others. Specific articles of manufacture contemplated
include
semiconductor wafer handling devices, vacuum chucks, electrostatic chucks, air
bearing
housings or support frames, electronic packages and substrates, machine tool
bridges and
bases, mirror substrates, mirror stages and flat panel display setters.
DISCLOSURE OF THE INVENTION
It is an object of the instant invention to produce, a composite material that
is
lightweight, stiff, strong and substantially pore-free.
It is an object of the instant invention to produce a composite material that
has
utility in precision equipment applications requiring a degree of impact
resistance.
It is an object of the instant invention to produce a silicon-containing
composite
material of enhanced toughness.
It is an object of the instant invention to produce a toughened silicon-
containing
composite material without requiring complex reinforcement debonding/pull-out
mechanisms.
It is an object of the instant invention to produce a toughened silicon-
containing
composite material without resort to adding significant amounts of tough
metals.
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These objects and other desirable attributes can be achieved through the
engineering of silicon-containing composite bodies. According to the instant
invention, a
silicon-ceramic composite body is produced, preferably by a reactive
infiltration
technique. The ceramic component, such as a reinforcement or filler, is
selected such that
it has a higher coefficient of thermal expansion (CTE) than does the silicon
phase. At
least at some point during processing, the silicon phase is at a temperature
above its
normal ductile/brittle transition temperature of about 500°C, and
preferably above its
melting point. The formed composite body containing the silicon phase is then
cooled
below its ductile/brittle transition. During cooling, the ceramic phase
shrinks more than
does the silicon phase, thereby placing the latter in a state of compressive
stress. By the
time the composite body has cooled to substantially ambient temperature, i.e.,
about 20°C,
the induced compressive stress in the silicon phase is sufficient as to impart
a measurable
degree of semi-ductile character to the silicon phase, even at ambient
temperature.
This toughening effect has been observed in certain reaction-bonded boron
carbide composites (RBBC). In this system, a molten infiltrant containing
silicon and one
or more sources of boron is contacted to a porous mass that contains at least
some boron
carbide, and also containing at least some free carbon. The molten infiltrant
infiltrates the
porous mass without a pressure or vacuum assist to form a composite body of
near
theoretical density. The silicon component of the infiltrant reacts with the
free carbon in
the porous mass to form in-situ silicon carbide as a matrix phase. Further,
the tendency of
the molten silicon to react with the boron carbide component can be suppressed
or at least
greatly attenuated by the alloying or doping of the silicon with the boron
source. The
resulting composite body thus comprises boron carbide dispersed or distributed
throughout the silicon carbide matrix. Typically, some residual, unreacted
infiltrant phase
containing silicon and boron is also present and similarly distributed or
interspersed
throughout the matrix.
Regardless of the exact mechanism which is responsible for the observed
result,
the display of pseudo-ductility in silicon-containing composites such as RBBC
is a
surprising result; namely, because none of the constituents of this composite
system taken
by themselves exhibit toughness or ductility under ambient conditions, e.g.,
standard
temperature and pressure. Such composites showing enhanced toughness should
find
numerous commercial applications because ceramic engineers are often concerned
about
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the prospect of brittle failure, and are frequently engaged in searching for
ways to
increase toughness or impact resistance of ceramic-rich materials.
DEFINITIONS
"Areal Density", as used herein, means the mass of an armor system per unit
area.
"CTE", as used herein, means Coefficient of Thermal Expansion, and unless
otherwise noted, refers to the linear (not volumetric) CTE.
"Fine-grained", as used herein, means that the morphological features making
up
the microstructure of the reaction-bonded silicon carbide bodies of the
instant invention
are smaller than the microstructural features of most of the prior,
commercially available
reaction-bonded silicon carbide. Preferably, the microstructure of the instant
reaction-
bonded silicon carbide bodies is engineered such that the vast majority of
morphological
features do not exceed about 100 microns in size.
"Free Carbon", as used herein, means carbon that is intended to react with
molten
silicon to form silicon carbide. This term usually refers to carbon in
elemental form, but
is not necessarily limited thereto.
"Inert Atmosphere", as used herein, means an atmosphere that is substantially
non-reactive with the infiltrant or the porous mass or preform to be
infiltrated.
Accordingly, this definition includes gaseous constituents that might
otherwise be thought
of as mildly reducing or mildly oxidizing. For example, forming gas,
comprising about 4
percent hydrogen, balance nitrogen, might be considered to be an inert
atmosphere for
purposes of the present disclosure, as long as the hydrogen does not reduce
the filler
material and as long as the nitrogen does not appreciably oxidize the
infiltrant or filler
material.
"Mass Efficiency", as used herein, means the areal density of rolled
homogeneous
steel armor required to give the same ballistic performance as that of the
targets being
tested, expressed as a patio.
"Reaction Bonded Silicon Carbide", or "RBSC", refers to a ceramic composite
body produced by reaction-bonding, reaction-forming, reactive infiltration, or
self
bonding.
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"Reaction-Bonded Boron Carbide", or "RBBC", as used herein, means a class or
subset of reaction-bonded silicon carbide composites in which the filler or
reinforcement
of the composite, i.e., the phase being bonded, includes boron carbide.
"Reaction-Bonding", "Reaction-Forming", "Reactive Infiltration" or "Self
Bonding", as used herein, means the infiltration of a porous mass comprising
carbon in a
form that is available to react with an infiltrant comprising silicon to
produce a ceramic
composite body comprising at least some silicon carbide produced in-situ.
"Substantially inert filler material", as used herein, means a filler material
that is
substantially non-reactive chemically with the molten infiltrant material
under the specific
processing conditions.
"Total areal density", as used herein, means the areal density of ceramic
armor
material plus the areal density of any other material that should properly be
considered a
part of the assembly of components making up an armor system. Examples of
other
materials would be fiber reinforced polymeric materials frequently used to
back up or
encase a ceramic armor plate.
BRIEF DESCRIPTION OF THE FIGURES
Figure 1 is a cross-sectional view of a feeder rail as described in Example 1.
Figures 2A and 2B are front and side views, respectively, of a set-up used to
prepare the boron carbide reinforced silicon carbide composite tiles of
Example 1.
Figure 3 is an optical photomicrograph of a polished cross-section of the RBBC
material produced in accordance with Example 2.
Figure 4 is an optical photomicrograph of a polished cross-section of the SiC-
filled RBSC material produced in accordance with Example 4.
Figures SA andSB are approximately 350X magnification SEM fractographs of
RBSC and RBBC composite materials of Examples 4 and 2, respectively.
Figure 6 is an approximately 600X magnification SEM photomicrograph of a
crack in a polished section of the RBBC material of Example 2.
Figures 7A-7C illustrate several applications of the armor material embodiment
of
the instant invention.
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MODES FOR CARRYING OLTT THE INVENTION
Toughness enhancement and other desirable attributes can be achieved in
otherwise brittle materials through the application and engineering of silicon-
containing
composite bodies. The toughening phenomenon was first observed in a reaction-
bonded
boron carbide composite, in which a porous mass of boron carbide particulate
plus some
free carbon in the form of a carbonaceous resin was infiltrated with a mass of
molten
silicon containing some alloyed boron. The boron addition to the silicon
infiltrant,
however, does not appear to contribute to the semi-ductile behavior that was
seen in the
silicon phase of a fractured sample of the formed RBBC composite. Without
wishing to
be bound to any particular theory or explanation, it appears that the
toughening effect
results from constraining of the residual silicon phase in the composite body
by the higher
CTE boron carbide particles during cooling. Specifically, the boron carbide
phase places
the silicon phase in a state of compression during cooling of the composite to
ambient
temperature following thermal processing above the melting point of the
silicon infiltrant.
Thus, the silicon-toughening phenomenon should be generic to composite systems
in
which another component within the composite material places the silicon
component in a
state of compressive stress during or following thermal processing.
A convenient way to accomplish this toughening of silicon is to incorporate
the
silicon as a component of a composite material in which another component of
the
composite has a larger CTE than does the silicon. Particularly preferred is a
reactive
infiltration approach, such as that used to make reaction-bonded silicon
carbide
composites. There, molten silicon metal in bulk form is contacted to a porous
mass
containing substantially inert ceramic material such as silicon carbide
particulate, plus
some free carbon. The molten silicon wets the porous mass, and wicks into it.
Some of
the molten silicon chemically reacts with the free carbon to produce
additional silicon
carbide (in-situ). When the entire mass has been infiltrated by molten
silicon, the
infiltrated mass is solidified to yield a substantially pore-free composite
body featuring
silicon carbide, plus any other inert ceramic that was originally provided,
plus some
residual silicon metal distributed throughout the composite body. However, the
toughening effect has not been observed in these composite bodies that contain
substantially only SiC and Si, perhaps because there is insufficient
difference in CTE
between SiC and Si to create a sufficiently large compressive stress state in
the Si upon
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cooling to ambient temperature. By way of review, B4C has a CTE of about 5.4
ppm/K,
whereas the CTE of Si and SiC is each nominally about 4 ppm/K.
The fact that some of the silicon infiltrant reacts as it infiltrates the
porous mass in
a reaction-bonded composite system should not make a large difference on the
toughening of the residual silicon phase. Accordingly, the effect should also
be present in
those silicon infiltration composite systems that do not involve reaction of
the silicon
infiltrant, which are sometimes referred to as "siliconization." U.S. Patent
No. 3,951,587
to Alliegro et al. is an example of siliconization. Moreover, it may be
possible that other,
non-infiltration-based techniques may also exhibit the effect, such as powder
metallurgical techniques, whereby the silicon component of the composite may
be
admixed in powder form with the inert ceramic filler, consolidated to a
preform such as
by dry pressing, then sintered. Still further, it may be possible, through
utilization of
presses such as hot presses or hot isostatic presses, to produce bulk silicon
featuring this
semi-ductile behavior. Specifically, bulk silicon in a molten condition, or at
least above.
its ductile-brittle transition temperature, is cooled under externally applied
bulk pressure.
As mentioned earlier, Gogotsi et al. have observed metal-like ductility of
bulk silicon,
specifically, from the localized pressure of a cutting tool tip during a
machining
operation. Ibid.
As long as the volumetric loading of the higher CTE component is sufficient,
the
silicon component should be in a compressive stress state, regardless of
whether the
silicon is interconnected or isolated. As the volumetric loading of the higher
CTE phase
is reduced, however, at some point, the higher CTE phase may become isolated
instead of
being the matrix, or part of the matrix. Then, the stress state on the silicon
phase becomes
more complex, and it is unclear whether the toughening phenomenon can be
maintained
at such reduced loadings. For example, where the silicon phase makes up the
matrix and
the higher CTE material is dispersed as a noncontiguous phase within this
matrix, upon
cooling, both components would be expected to be in a tensile stress state in
radial
directions.
In any event, the phenomenon is definitely observed in reaction-bonded boron
carbide composites that are highly loaded in boron carbide, e.g., about 75
volume percent,
even though the silicon is interconnected, and thus could be thought of as
forming a co-
matrix. Strictly speaking, the matrix here consists of the in-situ silicon
carbide phase, and
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the boron carbide forms the reinforcement phase. The boron carbide, however,
being
present in substantially greater amount than the silicon carbide, strongly
influences the
CTE of the matrix, i.e., increases it. In this sense, then, the boron carbide
is sometimes
thought of as being part of the matrix.
In the RBBC system, calculation has shown that the compressive stresses
induced
in the silicon phase are extremely large. Here, the constraining particles,
namely the
boron carbide phase, also possesses a large stiffness (elastic modulus). If it
is necessary
to apply very large compressive stress to the silicon to cause it to behave in
a
pseudoductile fashion, then it may be necessary for the constraining phase to
also possess
a high elastic modulus. Otherwise, the applied stress simply can be
accommodated by
straining the constraining phase. In other words, if large compressive
stresses are
required, then aluminum probably could not be the constraining phase. Most of
the
candidate constraining phases are expected to be ceramic-based, but certain
high modulus
metals such as tungsten might qualify.
Without wishing to be bound to any particular theory or explanation, it may be
that by solidifying the silicon component of a composite material under
conditions
whereby the silicon is placed into a state of compressive stress, 'the
population of atomic
scale defects such as dislocations is increased. As mentioned earlier, there
is some
experimental evidence in the prior art that introducing surface dislocations
in silicon by a
grinding process decreased the ductile-brittle transition temperature. Another
possibility
is that solidifying in the compressive state permits the appearance of one or
more high
pressure phases of silicon, such phases) exhibiting the semi-ductile
character, and which
phases) would not be favored thermodynamically under normal, ambient pressure
conditions. Such a high pressure phase may itself exhibit the observed semi-
ductile
behavior, or perhaps may function in a way similar to a transformation-
toughening
mechanism. There, the high pressure form may be denser than the ambient
pressure
crystallographic form, and when the compressive stress state is diminished or
removed,
for example, when a tensile crack moves into close proximity, the high
pressure phase
transforms to the ambient pressure phase, with a concurrent volume increase
associated
with this transformation. The volume increase tends to dissipate at least some
of the
tensile stress on the neighboring material, thereby inhibiting further crack
growth by
requiring the application of additional strain energy to the material to raise
the tensile
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stress and start the crack propagating again. Still another possibility is
that the
compressive state of the silicon "locks in" the high temperature
crystallographic form at
ambient temperature. This high temperature form of silicon has the "beta tin"
crystal
structure, and is ductile. Thus, this isomorph is sometimes referred to as
"metallic
silicon." In contrast, the diamond form of silicon is sometimes referred to as
a
"metalloid" or a "semi-metal".
Yin et al. state that their calculations show that silicon will transform at
ambient
temperature from it usual diamond crystalline form to the beta-tin form under
high
hydrostatic pressure. (M.T. Yin and Marvin L. Cohen, "Microscopic Theory of
the Phase
Transformation and Lattice Dynamics of Si", Phys. Rev. Letters, 45, 12, (1980)
pp.1004-7.) Thus, this high pressure phase is the beta-tin phase, according to
these
authors.
Whatever the operative mechanism, a fracture toughening phenomenon has now
been observed in boron carbide -containing composite materials made by a
silicon
infiltration technique. This result was quite unexpected because none of the
starting
materials is particularly tough, and none of the components of the fornled
composite body
was expected to be tough, or to display any degree of ductility during
fracture.
In accordance with a preferred embodiment of the present invention, a
substantially pore-free, mechanically strong composite material of increased
fracture
toughness is produced that includes at least some boron carbide, preferably in
a large
volume fraction or combined with one or more exceptionally hard, stiff
materials such as
silicon carbide to yield a large fraction of very hard, very stiff phase.
Furthermore,
through careful control of the processing conditions, e.g., to suppress
reaction of the
boron carbide phase, a superior material can be produced that exhibits
exceptional
specific stiffness and resistance to impact from ballistic projectiles,
particularly small
arms fire. In addition, the composite bodies produced according to the present
invention
maintain dimensional tolerances upon thermal processing better than do hot
pressed and
sintered bodies.
Thus, for economy and manufacturing flexibility, among other reasons, the
preferred embodiment boron carbide -based composite bodies of the instant
invention are
produced by a reactive infiltration technique, usually termed "reaction
forming" or
"reaction bonding", whereby a molten infiltrant comprising silicon is
contacted to a
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porous mass comprising carbon and at least one hard ceramic filler material
that includes
boron carbide. The molten silicon-based material infiltrates the
interconnected porosity
in the porous mass or preform. The molten silicon contains one or more sources
of boron
in a quantity sufficient to attenuate the tendency of the boron carbide
component to
chemically react with the molten silicon. Concurrent with the infiltration,
the silicon
reacts with at least some of the carbon in the porous mass or preform to form
silicon
carbide. The amount of infiltrant is generally provided in such a quantity
that the carbon
in the porous mass or preform is completely reacted to silicon carbide, with
sufficient
additional infiltrant supplied to fill any remaining void space between the
filler material
and the in-situ silicon carbide. The resulting composite materials feature a
matrix of the
in-situ silicon carbide. Dispersed throughout the matrix is the ftller and
residual,
unreacted infiltrant material. As the residual infiltrant is often
interconnected, it is
sometimes considered as part of the matrix of the composite.
In terms of the preferred processing conditions, atmospheres that are
compatible
with this type of infiltration include vacuum or inert atmospheres such as
argon, although
vacuum is preferred. The vacuum does not have,to be "hard" or high vacuum;
that
provided by a mechanical "roughing" pump is entirely adequate. Although the
infiltration tends to be more robust at the higher temperatures, it is also
more aggressive,
which could give rise to unwanted side reactions, particularly of the boron
carbide
component. Further, it is more difficult to confine the infiltrant spatially
at higher
temperatures. Moreover, higher processing temperatures are more likely to give
rise to
exaggerated grain growth. For all of these reasons, the preferred processing
temperatures
are those that are generally low yet consistent with reliable infiltration.
For infiltrating
silicon-based metals into a boron carbide -containing particulate mass in a
rough vacuum
environment, temperatures in the range of about 1450°C to 1600°C
should be
satisfactory.
Boron carbide is an especially attractive filler material candidate where the
mass
of the formed composite article is of concern because of its low theoretical
density of
about 2.45 to 2.55 grams per cubic centimeter. (The range in reported
theoretical density
may be because boron carbide is not a line compound per se, but instead can
exhibit a
limited range of stoichiometry.) Because the Young's Modulus of boron carbide
is
comparable to that of silicon carbide (about 450 GPa), boron carbide has a
higher specific
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stiffness than does silicon carbide. High specific stiffness is a valuable
property in
applications such as those requiring precise motion and control of motion,
especially
where large loads or high accelerations are involved. Moreover, boron carbide
is even
harder than silicon carbide. Thus, a RBSC composite body featuring boron
carbide as a
reinforcement or filler material (i.e., "RBBC") may offer higher hardness yet
lower
specific gravity as compared to a RBSC~ composite having silicon carbide as
the filler
material.
Under most of the prior silicon infiltration conditions, however, boron
carbide is
at least somewhat reactive with the molten silicon. Although one reaction
product of
such reaction is more in situ silicon carbide, where one is attempting to
maximize the
boron carbide loading, it would be desirable if the boron carbide could remain
substantially unaffected by the infiltrant, e.g., if the silicon did not react
with the boron
carbide. 'This problem is solved by dissolving some boron into the molten
silicon, thereby
reducing the activity of the silicon for reaction with boron carbide. Although
pure silicon
will eventually become saturated in boron and carbon as it reacts with the
boron carbide
phase in the porous mass or preform, this approach is not preferred, unless
this porous
mass or preform is "sacrificial", and not the ultimate article of commerce
being produced.
In many instances, reaction of the boron carbide reinforcement with the
silicon infiltrant
has led to cracking of the resulting silicon carbide composite body. Instead,
what is
preferred is to provide a source of boron to the silicon-based infiltrant
prior to the
infiltrant making contact with the boron carbide in the porous mass or
preform. Any
boron-containing substance that can be dissolved in silicon may be useful in
the context
of the instant invention; however, elemental boron or boron carbide are
particularly
preferred.
One can envision any number of techniques for adding a boron source material
to
the silicon infiltrant. The approach preferred according to the instant
invention is to
support the preform to be infiltrated on, and to feed the infiltrant into the
preform by way
of, kiln furniture consisting of a porous preform comprising boron carbide.
Specifically,
a silicon-containing infiltrant can infiltrate kiln furniture (later referred
to as a "feeder
rail" or "beam") containing at least some boron carbide. The kiln furniture
may be
provided in either the porous condition, e.g., as a preform; or in the
"already infiltrated"
condition, e.g., as a composite body. The preform that ultimately is intended
to become
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an article of commerce upon infiltration, sometime referred to as the "object"
preform, is
supported on the kiln furniture. The silicon-containing infiltrant dissolves
at least some
of the boron carbide of the kiln furniture, and may even become saturated with
carbon
and/or boron. When this molten silicon then continues to infiltrate into the
object
preform that is in contact with the kiln furniture, the infiltrating silicon
will react very
little if at all with the boron carbide in the object preform. Any cracking of
the kiln
furniture as a consequence of silicon reacting with the boron carbide in the
kiln furniture
should not unduly affect the continued infiltration of the silicon into the
object preform.
Of course, the supporting kiln furniture is not required to contain boron
carbide per se.
Many boron-containing substances in which the boron is able to dissolve in the
silicon
component of the infiltrant should be satisfactory; however, substances such
as boron
oxide may not be sufficiently refractory under the thermal processing
conditions. Further,
the boron source is not required to be located in the kiln furniture; it may
be alloyed or
otherwise introduced into the silicon component of the infiltrant at most any
point prior to
the silicon making contact with the boron carbide of the object preform. For
example, it
is useful when building the "lay-up" for infiltration to supply boron carbide
particulate to
the bottom of the vessel housing the molten silicon infiltrant, dispersed, for
example, as
loose powder between the feeder rails. Moreover, the presence of a boron
nitride coating
on the porous mass or preform to be infiltrated also helps to suppress the
boron carbide
reaction.
The hardness of the composite is proportional to the volume fraction of hard
phases such as silicon carbide or boron carbide making up the composite
material. One
technique for maximizing the amount of hard phase in the composite body is to
produce a
preform that is highly loaded volumetrically in the hard phases. Highly loaded
preforms
can be produced by utilizing a distribution of filler material particle sizes
sufficiently
wide so that small particles can nest or fit within the interstices of larger
particles.
Because these two parameters of maximizing the loading of hard fillers in the
preform
while capping or limiting the size of the largest particles inherently are at
odds with one
another, careful attention to processing parameters is required to achieve
both in the same
material. Fortunately, one can still produce a preform that is relatively
highly loaded in
hard filler while limiting the size of the filler bodies in such a way that,
for example, at
least 90 percent by volume are smaller than about 100 microns in diameter.
Even with
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this upper limit on the size of the largest particles, it is still possible,
for instance, to
produce preforms that are 65 volume percent or more loaded in hard ceramic
phases such
as boron carbide.
An important, but by no means the only important embodiment or application of
these boron carbide containing composites is that of armor for defeating
ballistic
projectiles. A particularly high-performing armor material can be realized
when the
composite system is designed such that it features a sufficiently high
volumetric loading
of the hard ceramic phases, particularly the boron carbide phase, as to meet
overall
hardness levels believed to be important, and then also so designed as to
limit the size of
the largest grains or crystals making up the composite body. The alloying of
the silicon
infiltrant with a boron source material facilitates the preservation of boron
carbide,
particularly of the finer grains of boron carbide, in the final composite
structure, by
suppressing the chemical reaction of boron carbide with silicon.
Although most any of the known techniques may be employed to produce a
porous preform that can be infiltrated by a molten infiltrant comprising
silicon, the
techniques that seem to be better able at producing preforms that are highly
loaded with
one or more fillers are those that utilize a liquid phase, for example,
sediment casting, slip
casting or thixotropic casting.
Again, it is thought that the toughening phenomenon of the instant invention
is not
limited to reaction-bonded composite systems containing boron carbide, but can
be
generalized to composite systems having one or more components whose CTE is
greater
than that of silicon. In this regard, other substantially inert filler
materials might include
titanium diboride or aluminum nitride. Here, the thermal expansion
coefficients are about
8.1 and 4.5 ppm/I~, respectively.
The filler material making up the porous mass to be infiltrated may be
provided in
a number of different morphologies, including particulates, platelets, flakes,
whiskers,
continuous fibers, microspheres, aggregate, etc. Particulates are often
preferred for
reasons of economy and availability.
While not possible through visual inspection, it is possible using diffraction
techniques to distinguish the silicon carbide matrix that is reaction-formed
from any
silicon carbide that may be present as a reinforcement or filler material.
Specifically, the
reaction-formed silicon carbide typically is of the beta polymorph, at least
under the
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instant processing conditions. In contrast, most commercially available
silicon carbide,
particularly the commodity grades, is the alpha form that is so commonly used
as a filler
material. Accordingly, one can provide at least approximate quantitative data
as to the
relative amounts of each that are present in the composite body.
A wide range of sizes of filler material bodies can be successfully
infiltrated using
the reaction-forming process, e.g., bodies ranging from several millimeters in
size down
to bodies on the order of a micron in size. Again, when the goal is to produce
structures
that possess high strength, high ballistic~resistance and/or low surface
roughness, the filler
bodies should not be permitted to get much larger than about 100 microns in
size.
In addition to keeping the starting size of the filler relatively fine, the
porous mass
of filler material containing boron carbide or any other constituent of the
composite
whose reaction with silicon is thermodynamically, but perhaps not kinetically
favored,
and whose reaction with silicon is not desired, should not be exposed to
excessive
temperatures, especially during infiltration. As will be shown in more detail
below, a
porous mass of boron carbide particulate has been successfully infiltrated
with a boron-
doped silicon infiltrant at a temperature of about 1550°C without
causing chemical attack
on the boron carbide phase.
Further on the subject of producing composite bodies containing large amounts
of
hard phase(s), these high loadings should not be accomplished through
production of
large amounts of the in-situ silicon carbide phase, but instead through the
engineering of
highly loaded masses of the filler material. For example, the porous mass to
be infiltrated
preferably contains free or elemental carbon as the carbon source to form the
in-situ
silicon carbide. The amount of this free carbon should be limited to
(generally) no more
than about 10 percent by volume of the porous mass, and preferably, no more
than about
5 or 6 percent. Thus, in general, the amount of silicon carbide produced in-
situ should be
limited to no more than about 24 volume percent of the final composite body.
Among the
problems that result from excessive reaction during the infiltration process
are
temperature spikes due to the exothermic nature of the chemical reaction of
silicon and
carbon. Such temperature spikes can cause cracking due to localized thermal
expansion.
Also, the conversion of elemental carbon to silicon carbide entails a
volumetric expansion
of about 2.35 times. Thus, large amounts of reaction are also detrimental from
the
standpoint that the large volumetric change can also lead to cracking.
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A source of reactable or "free" carbon in the porous mass is not required to
achieve infiltration, although.such a source assists in the infiltration
process. When such
a carbon source is utilized, however, it usually takes the form of elemental
carbon, such
as graphite. For many applications, particularly those requiring high
stiffness, it is
desirable that the silicon carbide of the resulting composite body be at least
partially
interconnected. This outcome is more readily achieved if the carbon in the
porous mass
or preform is interconnected. Further, interconnected carbon in the porous
mass or
preform assists the infiltration process in terms of speed and reliability. In
a preferred
embodiment, the carbon is introduced to the porous mass as a resin. This
mixture may
then be molded to the desired shape. Curing the resin renders the porous mass
self
supporting, e.g., as a preform. During subsequent thermal processing, or
during an
intervening firing step, typically in a non-oxidizing atmosphere; the resin
pyrolyzes to
carbon in interconnected form to yield a preform containing at least about 1
percent by
volume of carbon. The resin infiltration and pyrolysis cycle may be repeated
one or more
times if an increase in the carbon content is needed.
The following non-limiting examples further illustrate the instant invention.
EXAMPLE 1
This example demonstrates the production via reactive infiltration of a Si/SiC
composite body containing a boron carbide reinforcement, i.e., "RBBC". More
specifically, this Example demonstrates the infiltration of a silicon-
containing melt into a
preform containing an interconnected carbon phase derived from a resinous
precursor,
and silicon carbide and boron carbide particulates.
Preforms were prepared by a sedimentation casting process. Specifically, about
28 parts of water were added to 100 parts of ceramic particulate and 8 parts
of
KRYSTAR 300 crystalline fructose (A.E. Staley Manufacturing Co.) to make a
slurry.
The ceramic particulate content consisted of about equal weight fractions of
220 grit
TETRABOR° boron carbide (ESK GmbH, Kempten, Germany, distributed
by
MicroAbrasives Corp., Westfield, MA) having a median particle size of about 66
microns
and 500 grit CRYSTOLON green silicon carbide (St. Gobain/Norton Industrial
Ceramics)
having a median particle size of about 13 microns (Grade 500 RG). The solids
and
liquids were added to a plastic jar and roll mixed for about 40 hours. The
slurry was de-
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aired in about 760 mm of vacuum for about 5 minutes. About 15 minutes prior to
casting,
the slurry was re-roll mixed to suspend any settled particulates.
A graphite support plate was placed onto a vibration table. A rubber mold
having
a cavity of the desired shape to be cast was wetted with a surfactant (Sil-
Clean, Plastic
Tooling Supply Co., Exton, PA). The wetted rubber mold was then placed onto
the
graphite plate and allowed to dry. The slurry was poured into the cavity.
Vibration was
commenced.
The residual liquid on the top of the casting was blotted up with a sponge
periodically during sedimentation. After the particulates had fully settled
(about 3 hours),
vibration was halted. The graphite plate, the rubber mold and the castings
inside were
transferred from the vibration table to a freezer maintained at a temperature
of about
minus 20°C. The casting was thoroughly frozen in about 6 hours, thereby
forming a self
supporting preform.
From the freezer, the frozen preform was demolded and placed onto a graphite
setter tray. The graphite tray and preform were then immediately placed into a
nitrogen
atmosphere furnace at ambient temperature. The furnace was energized and
programmed
to heat to a temperature of about 50°C at a rate of about 10°C
per hour, to hold at about
50°C for about 8 hours, then to heat to a temperature of about
90°C at a rate of about
10°C per hour, to hold at about 90°C for about 8 hours, then to
heat to a temperature of
about 120°C at a rate of about 10°C per hour, to hold at about
120°C for about 4 hours,
then to heat to a temperature of about 600°C at a rate of about
50°C per hour, to hold at
about 600°C for about 2 hours, then to cool down to about ambient
temperature at a rate
of about 100°C per hour. This firing operation pyrolyzed the fructose,
yielding a well-
bonded preform containing about 2.7 percent by weight carbon.
The above-mentioned steps were employed to produce two "beam" or feeder rail
preforms and a number of tile preforms. Each tile preform had a mass of about
174
grams and had overall dimensions of about 100 mm square by about 9 mm thick.
Each
rail preform had a cross-section as depicted in Figure 1 and measured about
220 mm long
by about 15 mm wide by about 25 mm thick. During infiltration of the tile
preforms,
these rails would serve as a conduit for conducting molten infiltrant toward
and into the
tile preforms.
Next, a set-up to confine the infiltration process was prepared.
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Referring to Figures 2A and 2B, the interior surfaces of a Grade ATJ graphite
tray
31 (Union Carbide Corp., Carbon Products Div., Cleveland, OH) measuring about
790
mm by about 230 mm by about 51 mm deep were spray coated with a boron nitride
slurry
or paint 33 using a Model 95 Binks spray gun. The boron nitride paint was
prepared by
diluting about 1800 grams of LUBRICOAT boron nitride paste (ZYP Coatings, Oak
Ridge, TIC with deionized water to a volume of about 1 gallon (3.7 liters).
Two
relatively light coats of this boron nitride paint were applied, with brief
ambient
temperature drying in air between coats.
The boron nitride-coated tray was then placed into a larger graphite chamber
35
having interior dimensions of about 825 mm long by about 270 mm wide by about
320
mm in height. The chamber also featured means for supporting a parallel row of
graphite
dowel rods.
Referring now specifically to Figure 2B, two plies of PANEX°30 low
oxidation
carbon cloth 44 (Grade PW03, plain weave, 115 g/m2, Zoltek Corp., St. Louis,
MO)
weighing about 48 grams and measuring about 790 mm by about 230 mm was placed
on
the floor of the coated graphite tray 31, 33. Four boron carbide rail preforms
42, each
having a mass of about 190 grams and a length of about 200 mm, were placed on
top of
the cloth and arranged parallel to the length dimension of the tray. Silicon
in lump form
21 (Grade LP, Elkem Metals Co., Pittsburgh, PA) and comprising by weight about
0.5 .
percent Fe (max) and the balance Si, was then distributed more or less
uniformly over the
carbon cloth between the individual preform rails. Calculations showed that
about 1510
grams of silicon infiltrant would be required to completely react the
elemental carbon and
fill the interstices in the cloth, feeder rail preforms and tile preforms;
however, about 10%
additional silicon was provided to the set-up.
Graphite dowel rods 49 measuring about 0.25 inch (6 mm) in diameter and spray
coated with boron nitride 33 were placed into graphite holders or supports 47.
A total of
fifteen square tile preforms 41 (only four are shown in the Figure) similarly
spray coated
with boron nitride 33 were placed across the two rails edgewise in each half
of the tray.
As the boron nitride tended to act as a barrier material hindering over-
infiltration, the
surface of the tiles that were to contact the boron carbide preform rails were
left uncoated.
The top of the chamber was covered with a loose-fitting (non-hermetically
sealing) graphite lid 34 featuring a number of approximately 1 cm diameter
through-holes
CA 02463089 2004-04-05
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36 to permit atmosphere exchange. The holes were covered with a piece of
graphite felt
3~ which was held in place with a graphite block 40 which served as a dead
load, thereby
completing the set-up.
The completed set-up was then placed into a vacuum furnace at about ambient
temperature (e.g., about 20°C). The air was evacuated using a
mechanical roughing
pump, and a rough vacuum of less than about 100 millitorr residual pressure
was
thereafter maintained. The lay-up was then heated from ambient temperature to
a
temperature of about 1350°C at a rate of about 200°C per hour.
After maintaining a
temperature of about 1350°C for about 1 hour, the temperature was
further increased to a
temperature of about 1550°C at a rate of about 200°C per hour.
After maintaining a
temperature of about 1550°C for about 1 hour, the temperature was
decreased to a
temperature of about 1450°C at a rate of about 100°C per hour.
Without holding at this
temperature, the lay-up temperature was further decreased to a temperature of
about
1300°C at a rate of about 25°C per hour, which was immediately
followed by a cooling at
a rate of about 200°C per hour to approximately ambient temperature.
Following this heating schedule, the chamber and its contents was recovered
from
the vacuum furnace, disassembled and inspected. The silicon infiltrant had
melted and
infiltrated through the carbon cloth, thereby converting the carbon cloth to
silicon carbide
cloth. The molten silicon infiltrant had also infiltrated through the rail
preforms and into
the square tile preforms, and reacting with the elemental carbon therein, to
form dense,
silicon carbide matrix composite bodies having a boron carbide reinforcement.
Because
each tile preform was supported by the rails in line contact, only low-to-
moderate hand
force was sufficient to separate the RBBC composite tiles from the feeder rail
composite
material.
EXAMPLE 2
The technique of Example 1 was substantially repeated,' except that no silicon
carbide particulate was used in fabricating the preform, and the particle size
distribution
of the boron carbide was modified such that substantially all particles were
smaller than
about 45 microns. Following the pyrolysis step, the preforms contained about
75 percent
by volume of the boron carbide particulate and about 4 percent by volume of
carbon.
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After infiltration, the ceramic material contained nominally 75 vol. % boron
carbide, 9 vol. % reaction-formed SiC, and 16 vol. % remaining Si (i.e., an
RBBC
composite). A polished section was examined using a Nikon Microphot-FX optical
microscope. An optical photomicrograph of the material is shown in Figure 3.
It is
clearly evident that, by careful selection of processing conditions, including
addition of a
source of boron to the silicon infiltrant, little growth and interlocking of
the particles has
occurred, thus allowing a relatively fine microstructure to be maintained. For
instance,
the photomicrograph shows little visible reaction between the silicon and
boron carbide as
a result of the infiltration process.
EXAMPLE 3
The technique of Example 2 was substantially repeated, except that, before
supplying the silicon infiltrant to the lay-up, a monolayer of TETRABOR~ boron
carbide
particulate (220 grit, ESK) was sprinkled onto the carbon cloth between the
feeder rails.
The amount of silicon was concomitantly increased to account for the added
boron
carbide, and to maintain an excess supply of silicon of about 10 percent, as
in Example i .
EXAMPLE 4
'The technique of Example 2 was substantially repeated, except that silicon
carbide
particulate was substituted for the boron carbide particulate. As in Example
2, however,
the particle size distribution of the silicon carbide blend was such that
substantially all
particles were smaller than about 45 microns. Following the pyrolysis step,
the preforms
contained about 75 percent by volume of the silicon carbide particulate and
about 4
percent by volume of carbon.
After infiltration with molten Si, the resultant bodies consisted of 84 vol. %
SiC
(75 original and 9 reaction formed) and 16 vol. % Si (i.e., an RBSC
composite). A
typical microstructure (optical photomicrograph) of the material is shown in
Figure 4.
In the optical photomicrograph, it is not possible to differentiate between
the
original SiC and the reaction-formed SiC. As with the reaction bonded boron
carbide of
Example 2, the reaction bonded SiC ceramic shown in Figure 4 displays little
interlocking
and clustering of the SiC, thus allowing a relatively fine microstructure to
be maintained.
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EXAMPLE 5
The technique of Example 4 was substantially repeated, except that the feeder
rail
preforms were fabricated using the Example 2 procedure, i.e., made from boron
carbide
particulate rather than- silicon carbide particulate. Thus, the molten
infiltrant that
reactively infiltrated the silicon carbide tile preforms contained a boron
source; that is,
consisted of an alloy of silicon and boron.
Characterization of Mechanical and Physical Properties
After the fabrication step, various mechanical and physical properties of the
instant reaction-bonded ceramic composite materials were measured. Density was
determined by the water immersion technique in accordance with ASTM Standard B
311.
Elastic properties were measured by an ultrasonic pulse echo technique
following ASTM
Standard D 2845. Hardness was measured on the Vickers scale with a 2 kg load
per
ASTM Standard E 92. Flexural strength in four-point bending was determined
following
MIL-STD-1942A. Fracture toughness was measured using a four-point-bend-chevron-
notch technique and a screw-driven Sintech model CITS-2000 universal testing
machine
under displacement control at a crosshead speed of lmm/min. Specimens
measuring 6 x
4.8 x 50 mm were tested with the loading direction parallel to the 6 mm
dimension and
with inner and outer loading spans of 20 and 40 mm, respectively. The chevron
notch,
cut with a 0.3 mm wide diamond blade, has an included angle of 60° and
was located at
the midlength of each specimen. The dimensions of the specimen were chosen to
minimize analytical differences between two calculation methods according to
the
analyses of Munz et al. (D.G. Munz, J.L. Shannon, and R.T. Bubsey, "Fracture
Touglzraess Calculation from Maxirnuna Load in Four' Point Bead Tests of
Chevron Notch
Specimens," Int. J. Fracture, 16 8137-41 (1980)).
Results of density, Young's modulus, flexural strength and fracture toughness
of
some of these reaction-bonded ceramics are provided in Table I. When
appropriate, the
results are provided as a mean +/- one standard deviation.
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Table I
Example 4 Example 2
RBSC RBBC
Density (lcg/m') 3060 2570
Young's Modulus (GPa) 384 +/- 382 +/-
2 6
Flexural Strength (MPa)284 +/- 278 +/-
14 14
Fracture Tou hness 3.9 +/- 5.0 +/-
MPa-ml~z 0.5 0.4
The density of the silicon carbide-based reaction-bonded composite material is
about 6% lower than monolithic silicon carbide due to the presence of the
silicon phase,
which has relatively low density. This reduced density is important for
applications, such
as armor, that are weight specific. The boron carbide-based reaction-bonded
composite
material has very low density and is similar to that of monolithic boron
carbide.
The Young's moduli of the RBSC and RBBC ceramics are essentially the same,
and compare favorably with other high performance ceramic materials. The
specific
results are as predicted based on the Young's modulus values for dense silicon
carbide,
boron carbide and silicon: approximately 450, 450 and 120 GPa, respectively.
In
particular, on a weight specific basis, the RBBC composite has a very high
Young's
modulus.
The fracture toughness of the RBSC composite material of nominally 4 MPa-ml~z,
is consistent with most SiC-based ceramics. Surprisingly, the RBBC composite
material
shows a 28% increase in toughness relative to the reaction bonded SiC
material, despite
the fact that no ductile phase was added. A possible explanation for this
increased
toughness was found by examining fracture surfaces, as is explained in the
next section.
Analysis of Fracture Surfaces
The relatively high fracture toughness of the RBBC ceramic was unexpected. To
gain an understanding for this result, the fracture surfaces of the RBSC and
RBBC
ceramics were studied and compared. The SEM fractographs for the two materials
are
provided in Figures SA and SB, respectively.
A significant difference between the two fracture surfaces is seen. The RBSC
ceramic shows brittle, transgranular fracture of the SiC particles. Also,
brittle fracture of
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the Si matrix is seen. In addition, some indications of interfacial cracking
between the Si
and SiC are seen. The RBBC ceramic shows brittle, transgranular fracture of
the boron
carbide particles. However, the silicon component shows some highly unexpected
ductile
behavior with the characteristic chisel-like rupture pattern. In addition, no
evidence of
failure at the interfaces between the particles and matrix is seen. It is felt
that the
observed semi-ductile failure of the silicon phase is contributing to the
relatively high
toughness of the RBBC ceramic (Table I).
Figure 6 is an SEM photomicrograph of a crack propagating through a polished
surface of the RBBC sample of Example 2. In at least two locations, a phase
bridges the
crack, similarly suggesting the presence of a ductile phase.
Silicon normally undergoes a brittle to ductile transition at around
500°C;
however, the transition temperature decreases as the dislocation density in
the silicon
increases. In the Warren study (Ibid.), more surface dislocations were
introduced to the
surface of a sample by grinding, which reduced the brittle-to-ductile
transition
temperature.
In the RBSC system, once the silicon phase solidifies, little additional
stress will
be induced in the silicon phase on cooling from the processing temperature
because both
silicon and silicon carbide have nominally the same CTE. Thus, the dislocation
density in
the silicon should be low. However, the situation is very different in the
RBBC ceramic.
Upon cooling from the process temperature, the boron carbide and silicon will
shrink at
different rates, with boron carbide shrinking faster due to its higher CTE,
i.e., about 5.4
ppm/I~ versus about 4 ppm/K. Thus, the silicon will become highly stressed and
thus will
have a high dislocation density. This high dislocation density may reduce the
ductile-
brittle transition temperature in silicon to below ambient temperature.
EXAMPLE 6
This Example demonstrates that the presence of boron associated with the
silicon
infiltrant, either as a dissolved species or as a precipitated species, is not
responsible for
the enhanced toughness observed in the RBBC composite body of Example 2.
Specifically, the RBSC composite body of Comparative Example 2 was measured
for fracture toughness, as described above. The fracture toughness was no
higher than for
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RBSC composite material not containing boron, e.g., the Example 4 material; in
fact, it
was slightly lower.
Thus, something other than boron is responsible for the increase in fracture
toughness of RBBC.
Utilit~of RBBC as Armor for Defeating Ballistic Projectiles
1. Hardness Data
Hardness is a very important parameter for armor materials. Previous work has
demonstrated that high mass efficiencies are only obtained versus hard armor
piercing
projectiles when the projectiles are fractured, and that to effectively
fracture the
projectile, an armor must have high hardness. (See, for example, M.L. Wilkins,
R.L.
Landingham, and C.A. Honodel, "Fifth Progress Report of Light Armor Program,"
Report No. UCRL-50980, University of CA, Livermore, Jan. 1971; also C. Hsieh,
"Ceramic-Faced Aluminum Armor Panel Development Studies," Appendix 9 of Report
No. JPL-D-2092, Jet Propulsion Laboratory, Feb. 1985.)
However; it is difficult to compare the many hardness data in the open
literature
because results can be highly dependent on test method and technique.
Therefore, for the
instant invention many different commercial materials were obtained. Hardness
measurements were then made on both the commercial materials and the new
reaction
bonded ceramics of the instant invention in an identical manner so that true
comparisons
could be made. The results are provided in Table II.
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Table II
Material Vickers' Hardness with
2 k Load ~ lmna2
7.62 mm M2 AP Bullet (Tool Steel) ~ 926 +/- 26
14.5 mm BS-41 Bullet (WC/Co) 1644 +/- 30
Sintered A1N 1044 +/- 63
Pure Si 1243 +/- 21
90% Sintered A1203 1250 +/- 89
Hot Pressed A1N 1262 +/- 51
99.5% Sintered A1203 1499 +/- 74
Hot Pressed A1203 2057 +/- 82
Hot Pressed TiB2 2412 +/- 135
Hot Pressed TiC 2474 +/- 188
Hot Pressed SiC 2640 +/- 182
Hot Pressed B4C 3375 +/- 212
RBSC ~ 2228 +/- 274
RBBC 2807 +/- 54
The RBSC and RBBC ceramics have very high hardnesses that are well in excess
of both tool steel and WC/Co projectiles. In both cases, the RBSC and RBBC
composites
have hardnesses that more-or-less reflect the weighted average hardness of the
constituents. In particular, because of the very high hardness of the boron
carbide
component, the RBBC composite system has a very high hardness value.
2. Ballistic Testing
Candidate ceramic armor materials were provided in the form of square tiles
measuring about 100 mm on a side.
To produce an armor target for testing, the ceramic tile is attached to a
SpectraShield~ polymer composite backing layer (AlliedSignal Inc., Morristown,
NJ).
This material is supplied as a 54 inch (1370 mm) wide roll consisting of 2
plies of
unidirectional fibers embedded in a resin matrix, with the fibers of one ply
being
orthogonal to the fibers of the other ply. A number of 12-inch (305 rnm) wide
sheets
were cut from the roll. The appropriate number of these sheets were then
laminated and
consolidated in an autoclave at an applied pressure of about 150 psi (1.3 MPa)
at a
temperature of about 121 °C for about 60 minutes, thereby forming a
rigid polymer
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composite plate. Following consolidation, a backing plate measuring about 12
inches
(305 mm) square was cut from the 54 by 12 inches (1370 by 305 mm) plate using
a band
saw or water jet. An approximately 5 inch (120 mm) square region in the center
of the
backing plate was lightly abraded using 120 grit sandpaper. After cleaning the
surfaces to
be bonded with isopropyl alcohol, a candidate armor tile was bonded to the
center of the
backing plate using two plies of 76 microns thick urethane film adhesive. The
bond was
cured under full vacuum in an oven maintained at a temperature of about 121
°C for about
30 minutes, thereby forming a ballistic test coupon.
The weight of the backing plate was varied according to the number of
laminates
used; the weight of the ceramic tile was varied according to the thickness
dimension to
which the ceramic tile was ground. In each instance, however, the total areal
density
(ceramic tile plus backing material) was maintained at roughly the same
amount.
A target for ballistic testing was assembled as follows: The ballistic test
coupon
was placed in front of 28 plies of KM2 (600 denier) blanket with rip-stop
nylon and
camouflage cordura covers to simulate the outer tactical vest (OTV) of a body
armor.
'The OTV simulant and test coupon were located in front of a 100 mm thick
block of
Roma Plastiline modeling clay that had been conditioned at a temperature of
about 35°C
for about 6 hours. The test coupon and OTV simulant were secured to the clay
block with
duct tape, and the clay block was backed up by a steel support structure that
was secured
to the test table, thereby completing the target.
The targets were shot at zero degrees obliquity using 7.62 mm caliber
projectiles
at varying velocities. The basic unit of ballistic penetration resistance used
in this testing
is the Vso, the velocity of the projectile at which partial penetration and
complete
penetration of the target are equally likely.
3. Ballistic Data
The instant RBBC materials of Example 2 were evaluated as candidate armors,
and compared to the SiC-filled RBSC material of Example 4, as well as to
commercial
hot pressed B4C (the control). In one series of tests, the reaction bonded SiC
and
commercial hot pressed B4C were tested ballistically versus ball rounds as the
ballistic
projectile; and in a second set of tests, the reaction bonded boron carbide
and the hot
pressed boron carbide were tested versus armor piercing (AP) rounds.
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The results of ballistic testing are provided in Tables III and IV. In Table
III, test
results versus a 7.62 mm M80 ball round for the RBSC of Example 4 and
commercial hot
pressed B4C (control) are provided. In Table IV, test results versus a 7.62 mm
AP M2
round for the RBBC of Example 2 and commercial hot pressed B4C are provided.
In each
case, the tables provide the areal density of the system, the mass efficiency
of the target,
and the normalized mass efficiency relative to the hot pressed B4C control.
The mass
efficiencies in the tables were determined based on available data for rolled
homogeneous
steel armor (RHA) versus the same threats. Specifically, the mass efficiency
was
calculated as the areal density of RHA required to give the same performance
divided by
the areal density of the tested targets.
Table III
Armor System Mass E~ciency Normalized Mass
Areal Density (RHA Equivalefat) Efficiency
Hot Pressed B4C 23.5 (4.82) 4.56 1.00
(control)
RBSC ~ 23.9 (4.89) 5.11 1.12
(Example 4)
Table IV
Armor Systent Mass E~ciency Normalized Mass
Areal Density (RHA Equivalent) E~ciency
Hot Pressed B4C ~ 29.0 (5.95) 4.53 1.00
(control)
RBBC ~ 30.2 (6.18) 4.85 1.07
The ballistic results show that the armor designs employing lower cost
reaction
bonded ceramics had mass efficiencies equivalent to armors of the same design
using hot
pressed ceramics. This has enabled the production of cost effective armor
products for
various applications. In Figures 7A and 7C, for example, the aircraft armor
and personnel
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armor tiles were fabricated from SiC-filled RBSC. The vehicle armor plate of
Figure 7B
was fabricated from RBBC.
INDUSTRIAL APPLICABILITY
The toughness enhanced composite materials of the instant invention possess
exceptional hardness and stiffness, low specific gravity and relatively high
flexural
strength. Although the instant disclosure has focused primarily on the
potential
application of the instant materials as armor to defeat ballistic projectile
threats, they
should also find many applications where rigidity and low specific gravity are
important
materials properties, such as in semiconductor capital equipment components
(e.g., wafer
chucks, wafer handling arms, process chambers, etc.). Precision equipment
industries
such as metrology and precision optics will benefit from these composite
materials. The
instant composite materials might also be attractive as abrasives or wear-
resistant parts.
Further, these boron carbide composites may find applications in the nuclear
industry,
specifically, in applications where neutron absorption is important.
An artisan of ordinary skill will readily appreciate that numerous variations
and
modifications can be made to the invention as disclosed and exemplified above
without
departing from the scope of the invention as set forth in the appended claims.
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