U.S. patent application number 12/083080 was filed with the patent office on 2009-08-06 for bulk metallic glass/graphite composites.
Invention is credited to Jorg F. Loffler, Marco Siegrist.
Application Number | 20090194205 12/083080 |
Document ID | / |
Family ID | 37075793 |
Filed Date | 2009-08-06 |
United States Patent
Application |
20090194205 |
Kind Code |
A1 |
Loffler; Jorg F. ; et
al. |
August 6, 2009 |
Bulk Metallic Glass/Graphite Composites
Abstract
A composite material based on a bulk metallic glass is
disclosed. In an amorphous alloy phase forming a substantially
continuous matrix, a second phase comprising graphite particles is
embedded. The alloy is preferably zirconium based. The particles
may have a carbide surface layer, which may be formed phase
comprising carbide particles may also be present. The composite
material has high plasticity, high yield strength, good elasticity
and low coefficient of friction, which renders it a good candidate
for applications like joints, frictional bearings or Springs.
Inventors: |
Loffler; Jorg F.;
(Schneisingen, CH) ; Siegrist; Marco; (Bonstetten,
CH) |
Correspondence
Address: |
LERNER, DAVID, LITTENBERG,;KRUMHOLZ & MENTLIK
600 SOUTH AVENUE WEST
WESTFIELD
NJ
07090
US
|
Family ID: |
37075793 |
Appl. No.: |
12/083080 |
Filed: |
August 29, 2006 |
PCT Filed: |
August 29, 2006 |
PCT NO: |
PCT/CH2006/000466 |
371 Date: |
March 16, 2009 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
|
60722409 |
Oct 3, 2005 |
|
|
|
Current U.S.
Class: |
148/561 ;
148/403 |
Current CPC
Class: |
C22C 1/1068 20130101;
C22C 16/00 20130101; C22C 45/10 20130101; C22C 32/0084
20130101 |
Class at
Publication: |
148/561 ;
148/403 |
International
Class: |
C22C 45/10 20060101
C22C045/10 |
Claims
1. A composite material comprising: a substantially amorphous first
phase forming a substantially continuous matrix, said first phase
consisting essentially of an alloy; and a second phase embedded in
said matrix, said second phase comprising graphite particles.
2. The composite material according to claim 1, wherein said
graphite particles have a size in the range between 1 and 100
micrometers.
3. The composite material according to claim 1, wherein said
graphite particles have a size in the range between 25 and 75
micrometers.
4. The composite material according to claim 1, wherein said second
phase occupies between 3 volume percent and 20 volume percent of
said composite material.
5. The composite material according to claim 1, wherein said second
phase is selected such that, under compressive deformation of said
composite material up to yield, it induces a distribution of shear
bands spaced apart by less than about 5 micrometers around said
graphite particles.
6. The composite material according to claim 1, wherein said alloy
of the first phase, in the liquid state, is capable of wetting said
second phase.
7. The composite material according to claim 1, wherein said alloy
comprises at least about 40 atomic percent of a metal having a
negative enthalpy of formation for the reaction with graphite to
form a metal carbide.
8. The composite material according to claim 1, wherein said alloy
comprises at least about 40 atomic percent of zirconium.
9. The composite material according to claim 1, wherein said alloy
consists essentially of
Zr.sub.52.5Cu.sub.17.9Ni.sub.14.6Al.sub.10Ti.sub.5.
10. The composite material according to claim 1, wherein at least a
fraction of said graphite particles in said second phase have a
core consisting essentially of graphite and a surface layer
comprising at least one metal carbide.
11. The composite material according to claim 10, wherein said
surface layer has a thickness of at least 100 nanometers.
12. The composite material according to claim 10, wherein said
graphite particles have a size of at least about 25
micrometers.
13. The composite material according to claim 10, wherein said
surface layer comprises at least one metal carbide formed in situ
by a reaction of graphite with said alloy.
14. The composite material according to claim 10, wherein said
surface layer consists essentially of zirconium carbide.
15. The composite material according to claim 1, further comprising
a third phase embedded in said matrix, wherein said third phase
comprises crystalline particles.
16. The composite material of claim 15, wherein said third phase
comprises crystalline particles composed of the same elements as
said alloy of the first phase.
17. The composite material of claim 15, wherein said third phase
comprises carbide particles.
18. The composite material according to claim 17, wherein said
carbide particles comprise at least one metal carbide formed in
situ by a reaction of graphite with said alloy.
19. The composite material according to claim 17, wherein said
carbide particles consist essentially of zirconium carbide.
20. The composite material according to claim 17, wherein said
carbide particles have a size of less than or equal to about 10
micrometers.
21. Use of a composite material according to claim 1 for
manufacturing an object for use in a device selected from: a
frictional bearing, a joint, and a spring.
22. A method for manufacturing a composite material, said method
comprising: heating an alloy above its liquidus temperature to form
a liquid alloy; dispersing graphite powder in the liquid alloy to
form a finely dispersed mixture; cooling the mixture below its
glass transition temperature sufficiently rapidly for forming a
composite material comprising a substantially amorphous first phase
forming a substantially continuous alloy matrix and a second phase
embedded in said matrix, said second phase comprising graphite
particles.
23. The method of claim 22, wherein said alloy is heated above its
liquidus temperature by induction melting on top of said graphite
powder.
24. The method of claim 22, wherein said mixture is remelted at
least once for a time sufficiently long for a distinct carbide
layer to form on the surface of said graphite particles.
25. The method of claim 22, wherein said mixture is remelted at
least once for a time sufficiently long for a fraction of said
graphite particles reacting with at least one metal component of
said alloy to form metal carbide particles.
Description
CROSS-REFERENCE TO RELATED APPLICATION
[0001] The present application is a national phase entry under
U.S.C. .sctn.371 of International Application No.
PCT/CH2006/000466, filed Aug. 29, 2006, published in English, which
claims benefit of U.S. Provisional Patent Application No.
60/722,409, filed Oct. 3, 2005. The disclosures of all of said
applications are incorporated by reference herein.
FIELD OF THE INVENTION
[0002] The present invention relates to a composite material having
a first, amorphous alloy phase forming a substantially continuous
matrix and having a second, reinforcing phase embedded in the
matrix.
BACKGROUND OF THE INVENTION
[0003] Many amorphous metallic alloys with good glass-forming
ability have been developed over the last few years. These bulk
metallic glasses (BMGs) possess very interesting mechanical,
magnetic, thermophysical and structural properties. They display,
for example, up to double the fracture strength and four times the
elasticity of their crystalline counterparts and thus have a very
high potential for use as structural materials. Unfortunately,
these properties cannot be fully exploited due to the alloys'
brittle fracture behavior. With no crystalline structure,
deformation via dislocation movement is impossible, but takes place
in one or a few highly-localized shear bands. Even though BMGs may
show-some type of "ductile" fracture mechanism on a microscopic
scale, metallic glasses are generally brittle because the fracture
energy is concentrated in a very small volume of the sample.
Drastic enhancement of the plasticity of BMGs would lead to a
revolutionary new material for structural applications.
[0004] Achieving this increase in plasticity is a very timely
topic, and many researchers are working on it using different
approaches. Foreign-particle-reinforced BMGs, in-situ-formed BMG
composites, porous Pd-based BMGs, and monolithic Pt-based BMGs
displaying a high Poisson ratio have been investigated. All
approaches have one thing in common, namely the intention to
increase shear-band density so that fracture energy will be
distributed over a larger volume of the sample.
[0005] Foreign-particle reinforcement appears to be particularly
promising. The materials so far employed as foreign particles for
reinforcement include ductile metals like W, Ta, Nb, Mo or steel
and ceramics like WC, TiC, SiC or ZrC.
[0006] However, it is believed that the properties, in particular,
plasticity in relation to yield strength, of such previously
disclosed foreign-particle-reinforced BMGs are not yet at their
optimum.
[0007] Reinforcement by structures other than particles, such as
sheets, fibers or wires, has also been suggested. In particular,
reinforcement by carbon fibers or carbon nanotubes has been
disclosed in: [0008] Kim, C. P., Busch, R., Masuhr, A., Choi-Yim,
H. & Johnson, W. L. "Processing of Carbon-Fiber-Reinforced
Zr.sub.41.2Ti.sub.13.8Cu.sub.12.5Ni.sub.10.0Be.sub.22.5 Bulk
Metallic Glass Composites". Appl. Phys. Lett. 79, 1456-1458 (2001).
[0009] Bian, Z. et al., "Carbon-nanotube-reinforced
Zr.sub.52.5Cu.sub.17.9Ni.sub.14.6Al.sub.10T.sub.5 bulk metallic
glass composites", Appl. Phys. Lett. 81, 4739-4741 (2002); [0010]
Bian, Z. et al., "Carbon-nanotube-reinforced Zr-based bulk metallic
glass composites and their properties", Adv. Funct. Mater. 14,
55-63 (2004).
[0011] However, BMGs reinforced by carbon fibers will have
anisotropic properties, while alloys reinforced by carbon nanotubes
have a strong tendency to crystallize because the small nanotubes
act as heterogeneous nucleation sites. They have been shown to be
even more brittle than the corresponding monolithic BMG. In
addition, carbon nanotubes are relatively expensive to produce.
[0012] In-situ formed BMG composites have been shown to display
good combinations of yield strength and plasticity. In particular,
examples for zirconium carbide (ZrC) reinforcement are disclosed in
the following publications: [0013] Wang, W. H. & Bai, H. Y.
"Carbon-addition-induced bulk ZrTiCuNiBe amorphous matrix composite
containing ZrC particles". Mater. Lett. 44, 59-63 (2000). [0014]
Kato, H., Hirano, T., Matsuo, A., Kawamura, Y. & Inoue, A.
"High strength and good ductility of
Zr.sub.55Al.sub.10Ni.sub.5Cu.sub.30 bulk glass containing ZrC
particles". Scr. Mater. 43, 503-507 (2000). [0015] Hirano, T.,
Kato, H., Matsuo, A., Akihisa & Inoue, A. "Synthesis and
mechanical properties of Zr.sub.55Al.sub.10Ni.sub.5Cu.sub.30 bulk
glass composites containing ZrC particles formed by the in-situ
reaction". Mater. Trans. JIM 41, 1454-1459 (2000). [0016] Chen, F.
et al. "Crystallization of Zr.sub.55Al.sub.10Ni.sub.5Cu.sub.30 Bulk
Metallic Glass Composites Containing ZrC Particles". Mater. Trans.
JIM 43, 1-4 (2002).
SUMMARY OF THE INVENTION
[0017] It is an object of the present invention to provide a
composite material with an amorphous alloy phase that has a high
plasticity, in particular, a material having a high plasticity
while at the same time having a high yield strength.
[0018] It is a further object of the present invention to provide a
composite material with an amorphous alloy phase that has improved
thermal stability.
[0019] It is a further object of the present invention to provide a
composite material with an amorphous alloy phase that has improved
tribological properties, in particular, a low coefficient of
friction and a good resistance to abrasion.
[0020] These and other objects are achieved by a composite material
according to claim 1.
[0021] Thus, there is provided a composite material having at least
two phases. The composite material comprises [0022] a substantially
amorphous first phase forming a substantially continuous matrix,
the first phase consisting essentially of a (metallic) alloy; and
[0023] a second phase embedded in said matrix, said second phase
comprising graphite particles.
[0024] By embedding graphite particles in the amorphous alloy
matrix (BMG matrix), a remarkable increase in plasticity is
achieved, while yield strength and elasticity remain comparable to
that of the monolithic amorphous alloy. At the same time, favorable
properties under friction, such as a low coefficient of friction
(COF) and high resistance to abrasion are obtained, which makes the
composite material according to the present invention a good
candidate for dry sliding applications, such as frictional (plain)
bearings. Furthermore, the thermal stability is increased as
compared to the monolithic matrix material, i.e., the onset of
crystallization is shifted to higher temperatures.
[0025] In the context of the present invention, the term "graphite"
is to be understood to designate a form of elemental carbon in
which substantially all carbon atoms (or at least their vast
majority) exist in a sp2-hybridized state. In a perfect graphite
structure, the carbon atoms are arranged in layers having a
hexagonal structure. However, in more general terms, the term
"graphite" is to be understood to also include structurally less
well-defined materials such as pyrolytic carbon, in which the
layers are to some extent covalently bonded, or soot, in particular
carbon black, which may consist essentially of carbon in a
predominantly amorphous state.
[0026] The term "particle" is to be understood to designate a small
body having no well-defined axis of symmetry or approximate
symmetry and having roughly similar dimensions along all directions
in space. In particular, the term "particle" is to be understood as
excluding fibers, which extend along a preferred direction, or
(nano-)tubes, which likewise have a preferred direction
(approximate axis of symmetry). Fibers or nanotubes, when used as a
reinforcing material, impart very different properties to a
material than particles do.
[0027] The particles used in the present invention may have a broad
size range. There is no actual lower limit to the size; however,
for some matrix alloys, it has been found difficult to prevent
graphite particles smaller than about 10 micrometers from
completely reacting with the matrix alloy to form carbides.
Therefore, a minimum dimension of about 10 micrometers, preferably
about 25 micrometers is preferred. There is also no "hard" upper
limit for the size of the graphite particles. However, in practical
terms, often the size will be less than about 200 micrometers.
Preferably, the graphite particles have a size in the range between
about 25 and about 75 micrometers. The term "size" is to be
understood as designating an average of the dimensions of the
particles over all directions in space.
[0028] A very broad range for the volume fraction of the second
(reinforcing) phase relative to the total volume of the composite
material is possible. There is no actual lower limit. However, more
significant effects are expected if the second phase locally
occupies at least about 1 volume percent of the total volume of the
composite material. On the other hand, the theoretical maximum
volume fraction of the second phase is only limited by the
situation in which the particles are so densely distributed that
all the particles would touch each other. This volume fraction is
estimated to be well above 50 volume percent for standard graphite
powders. A preferred range for the volume fraction of the second
phase is between about 1 volume percent and about 20 volume
percent. The most preferred range depends on the average particle
size. For larger particles, a lower volume fraction appears to be
advantageous. By the way of example, for particles with sizes in
the range between about 25 and about 45 micrometers, a volume
fraction of the second phase in the range between about 3 and about
10 volume percent is preferred, leading to a marked increase in
plasticity and a marked decrease in COF, while yield strength and
hardness are only moderately affected. For particles with sizes in
the range from about 45 to about 75 micrometers, the preferred
volume ratio is between about 1 and about 6 volume percent. Of
course, the exact amount will also depend on the desired
application. It will also depend on the processing conditions, as
will become more clear below, when the formation of carbide layers
is discussed.
[0029] The volume fraction or concentration of the second phase may
vary over the volume of an object made from the material, e.g., the
concentration of the second phase may be higher near the surface of
an object than in the bulk. This is especially advantageous if the
object is to be used in a dry bearing application, where mainly the
surface properties are of interest.
[0030] An important effect of the graphite particles is to induce
the formation of closely spaced shear bands throughout the material
if a sample of the material is deformed. After compressive
deformation of the composite material up to yield, the density of
shear bands can easily be determined from optical images of the
fracture surface. Preferably, the properties of the second phase
(shape and size distribution of the graphite particles, volume
fraction etc.) are chosen such that shear bands are spaced apart by
less than about 5 micrometers around the graphite particles.
Preferably, the shear bands are preferably substantially uniformly
(homogeneously) distributed over regions substantially larger than
an average graphite particle, preferably over the whole matrix.
[0031] In order to ensure good mixing and close contact (on the
atomic scale) of the phases and a uniform distribution of the
second phase in the matrix, the matrix alloy, in the liquid state,
is preferably capable of wetting the particles of the second phase.
In other words, the surface of the particles is preferably wettable
by the (liquid) matrix alloy. Wettability is usually quantified by
the so-called contact angle. A surface is normally considered to be
wettable by a liquid if the contact angle is below 90.degree..
Liquid metals which are known to be capable of wetting graphite
surfaces include, in particular, zirconium, titanium, copper and
iron. Alloys containing a major proportion (e.g., at least about
40%) of one or more of these metals are expected to have a good
wetting behavior.
[0032] Another quantity which is relevant in the formation of the
composite materials according to the present invention is the
reactivity between the alloy and graphite to form carbides. Such a
reaction on the particle surface is, to some extent, desired as it
ensures a very close atomic bonding between the two phases, and as
it can be used to tailor the properties of the composite material.
This will be discussed in more detail below. A reaction between the
alloy and the graphite particles will take place if the enthalpy of
formation of metal carbides is negative. Therefore, it is preferred
that the alloy comprises at least about 40 atomic percent of one or
more metals having a negative enthalpy of formation for the
reaction with graphite to form a metal carbide. Examples include
zirconium and titanium.
[0033] Preferably, the alloy is a Zr-based alloy, i.e., it
comprises at least about 40 atomic percent of zirconium. Zirconium
is known to have a negative enthalpy of formation with graphite to
form ZrC (-106 kJ/mol) and to have a good wetting behavior for
graphite. Due to the well-known reaction behavior between Zr and
graphite, it may reasonably be expected that all Zr-based BMGs are
suitable matrix alloys for the composite materials of the present
invention. Many Zr-based alloys with good glass-forming abilities
(Zr-based BMGs) have been developed to date with widely varying
compositions. A non-exhaustive list of examples includes: [0034]
Zr.sub.58Cu.sub.22Fe.sub.8Al.sub.12 (Bio 1); [0035]
Zr.sub.57Nb.sub.5Al.sub.10Cu.sub.15.4Ni.sub.12.6 (Vit 106); [0036]
Zr.sub.41.2Ti.sub.13.8Cu.sub.12.5Ni.sub.10Be.sub.22.5 (Vit 1);
[0037] Zr.sub.46.75Ti.sub.8.8Ni.sub.10Cu.sub.7.5Be.sub.27.5 (Vit
4); [0038] Zr.sub.60Al.sub.10Cu.sub.30,
Zr.sub.55Al.sub.15Ni.sub.25, Zr.sub.55Cu.sub.30Al.sub.10Ni.sub.5,
Zr.sub.55Ti.sub.5Al.sub.10Cu.sub.20Ni.sub.10, and
Zr.sub.52.5Ti.sub.5Al.sub.12.5Cu.sub.20Ni.sub.10; [0039] Alloy of
composition (Zr.sub.xCu.sub.100-x).sub.80(Fe.sub.40A.sub.60).sub.20
with. Such alloys have been investigated extensively in WO
2006/026882, whose contents are incorporated herein by reference
for teaching bulk metallic glasses following specific principles in
their atomic composition. In particular, alloys with x=62, 64, 66,
68, 72.5, 77, 79, 81 or 83 have been investigated. [0040] Further
examples from WO 2006/026882 include:
(Zr.sub.95Ti.sub.5).sub.72Cu.sub.13Fe.sub.13Al.sub.2,
Zr.sub.70Cu.sub.13Fe.sub.13Al.sub.3Sn.sub.1,
Zr.sub.70Cu.sub.13Fe.sub.13Al.sub.2Cr.sub.2,
Zr.sub.70Cu.sub.13Fe.sub.13Al.sub.2Nb.sub.2,
Zr.sub.70Cu.sub.13Fe.sub.13Al.sub.2Zn.sub.2,
(Zr.sub.72Cu.sub.13Fe.sub.13Al.sub.2).sub.98Mo.sub.2,
(Zr.sub.72Cu.sub.13Fe.sub.13Al.sub.2).sub.98P.sub.2,
(Z.sub.95Hf.sub.5).sub.72Cu.sub.13Fe.sub.13Al.sub.2,
Zr.sub.70Cu.sub.11Fe.sub.11Al.sub.8,
Zr.sub.71Cu.sub.11Fe.sub.10Al.sub.8,
(Zr.sub.74Cu.sub.13Fe.sub.13).sub.90Al.sub.10,
Zr.sub.72Cu.sub.13Fe.sub.13Al.sub.2,
(Zr.sub.74Cu.sub.13Fe.sub.13).sub.98Al.sub.2,
Zr.sub.73Cu.sub.13Fe.sub.13Al.sub.1,
Zr.sub.72Cu.sub.13Fe.sub.13Al.sub.2,
Zr.sub.71Cu.sub.13Fe.sub.13Al.sub.3,
Zr.sub.72Cu.sub.12Fe.sub.12Al.sub.4,
Zr.sub.70Cu.sub.13Fe.sub.13Al.sub.4,
Zr.sub.72Cu.sub.11Fe.sub.11Al.sub.6
Zr.sub.72Cu.sub.11.5Fe.sub.11Al.sub.5.5,
Zr.sub.73Cu.sub.11Fe.sub.11Al.sub.5,
Zr.sub.71Cu.sub.11Fe.sub.11Al.sub.7,
Zr.sub.69Cu.sub.11Fe.sub.11Al.sub.9,
Zr.sub.70Cu.sub.10.5Fe.sub.10.5Al.sub.9,
Zr.sub.70Cu.sub.10Fe.sub.11Al.sub.9,
Zr.sub.70Cu.sub.11Fe.sub.10Al.sub.9,
Zr.sub.69Cu.sub.10Fe.sub.10Al.sub.11,
Zr.sub.69Cu.sub.10Fe.sub.11Al.sub.10,
Zr.sub.70Cu.sub.13Fe.sub.13Al.sub.2Sn.sub.2,
Zr.sub.72Cu.sub.13Fe.sub.13Sn.sub.2,
(Zr.sub.74Cu.sub.13Fe.sub.13).sub.98Sn.sub.2,
(Zr.sub.79Cu.sub.21).sub.80(Fe.sub.40Al.sub.60).sub.20,
(Zr.sub.81Cu.sub.19).sub.80(Fe.sub.40Al.sub.60).sub.20,
(Zr.sub.83Cu.sub.17).sub.80(Fe.sub.40Al.sub.60).sub.20,
(Zr.sub.66Cu.sub.34).sub.80(Fe.sub.40Al.sub.60).sub.20,
(Zr.sub.64Cu.sub.36).sub.80(Fe.sub.40Al.sub.60).sub.20, and
(Zr.sub.62Cu.sub.38).sub.80(Fe.sub.40Al.sub.60).sub.20.
[0041] Extensive experiments, to be described in more detail below,
have in particular been performed for composites whose matrix alloy
is represented essentially by the chemical formula
Zr.sub.52.5Cu.sub.17.9Ni.sub.14.6Al.sub.10Ti.sub.5. This alloy has
excellent glass-forming ability and has become known as "Vit 105".
For a Vit 105 matrix reinforced by graphite particles in the size
range of about 25 micrometers to about 45 micrometers, a plasticity
of up to 15% under compression and a yield strength of up to 1.5
GPa has been achieved. For reinforcement with particles in the size
range of about 45 micrometers to about 75 micrometers, a plasticity
of up to 18.5% under compression and a yield strength of up to 1.85
GPa has been achieved.
[0042] The mechanical properties of the composite material can be
tailored by providing graphite particles having a core of graphite
covered at least partially by a interfacial carbide layer. In other
words, in a preferred embodiment, at least a fraction of the
graphite particles have a core consisting essentially of graphite
and an interfacial layer comprising at least one metal carbide, in
particular, an interfacial layer consisting essentially of
zirconium carbide. The layer is preferably formed in situ by a
reaction of graphite with at least one metal in the surrounding
matrix (interfacial carbide formation in situ).
[0043] The interfacial layer may be very thin and might amount to
only a few atomic layers. For many alloys, such a layer might even
be unavoidable due to unavoidable in-situ reactions during
processing. A thin interfacial layer, formed in situ, will ensure
an intimate contact between particle and matrix, without much
influence on other properties. A thicker layer will also alter the
mechanical properties of the composite material, such as
plasticity, yield strength and, in particular, hardness, which
increases with increasing thickness of the layer. For applications
where an increased hardness is desired, it is advantageous if the
interfacial layer has a thickness of at least 100 nanometers. On
the other hand, it is advantageous if the thickness does not exceed
about 1.5 to 2 micrometers and is preferably below about 1
micrometer, in order not to induce a too brittle fracture
behavior.
[0044] In order to avoid that the graphite particles fully
transform into carbide particles, it is advantageous if they are
not too small. In particular, it is advantageous if the graphite
particles, including the interfacial carbide layer, have a size of
at least about 25 micrometers. This is especially true for Zr-based
matrix alloys.
[0045] The present invention further provides a three-phase
composite material that, in addition to the alloy matrix phase and
the graphite particle phase, comprises a third phase embedded in
the matrix, wherein the third phase comprises particles.
[0046] The third phase may comprise crystalline particles that are
composed of the same elements as the matrix alloy. Such particles
are usually formed during the cooling of the matrix alloy from the
melt. In particular, these particles may be nanocrystals with a
mean size below about 1 micrometer. To some extent, their presence
might be unavoidable; however, the processing conditions may also
be deliberately chosen such that a considerable fraction of such
particles is formed, e.g., up to 30 or 50%. These particles will
normally be composed of the same elements as the matrix alloy,
however, with different atomic fractions of the individual
elements.
[0047] Alternatively or additionally, the third phase may comprise
carbide particles. While the carbide particles may be preformed and
added to the matrix, preferably the carbide particles are formed in
situ by a reaction of graphite with the alloy. Such particles may
be formed by different mechanisms. In one example, they may result
from a substantially complete transformation of relatively small
graphite particles with at least one metal component of the alloy
into the corresponding carbide, with at most traces of graphite
remaining. In another example, they are the result of a mechanism
in which metal carbide that has formed on the surface of (larger)
graphite particles has been separated from this surface, e.g., by
strong stirring, and dispersed in the matrix. In a preferred
embodiment, the carbide particles consist essentially of zirconium
carbide.
[0048] The carbide particles preferably have a size of less than or
equal to about 10 micrometers. In particular, such particles can
readily be formed in situ by a complete transformation of similarly
small graphite particles by a reaction with the matrix alloy.
[0049] The various composite materials according the present
invention may be used in a variety of different applications in
which one or more of the following properties are required: high
plasticity, high yield strength, high elasticity, high elastic
constants, low coefficient of friction, high resistance to
abrasion. One example are articles employed in a dry frictional
(plain) bearing. In particular, the two-phase composite materials
described above, with graphite particles having no or only a
minimal interfacial carbide layer, are promising for such
applications. If higher hardness is additionally required, the
three-phase composite materials described above, with both graphite
particles and carbide particles, are advantageous. Another example
where, in addition to low COF, high plasticity and high yield
strength are important, are joints, in particular small joints
which experience comparably high loads such as joints between
different parts of a mobile telephone. Also for such applications
the materials of the present invention are particularly
advantageous. A further field of application are springs. Metallic
glasses are known to display an elastic limit which is 2-4 times
larger than for their crystalline counterparts. However, for
technological applications the large elastic limit cannot be fully
exploited due to the brittle fracture behavior of the monolithic
material. The plasticity achieved with the composites in this
invention allow a spring design fully exploiting the potential of
the matrix material.
[0050] The composite material of the present invention may be
prepared by various methods. In an advantageous process an alloy
with good glass-forming capabilities is provided. A good glass
former is capable of retaining an amorphous state when cooled from
its melt at or above a critical cooling rate, where the critical
cooling rate is no more than about 1000 K per second, preferably no
more than about 100 K per second. The process then comprises the
following steps: [0051] heating the alloy above its liquidus
temperature to form a liquid alloy; [0052] dispersing graphite
powder in the liquid alloy to form a finely dispersed mixture;
[0053] cooling the mixture below its glass transition temperature
sufficiently rapidly for forming a composite material comprising an
amorphous, substantially continuous alloy matrix with a second
phase embedded in said matrix, said second phase comprising
graphite particles.
[0054] In such a process, the alloy may be heated above its
liquidus temperature by induction melting on top of the graphite
powder. Before the final cooling step, the alloy may optionally be
processed once or repeatedly at a temperature above the melting
(liquidus) temperature for a time sufficiently long for a carbide
layer to form on the surface of said graphite particles. If a
three-phase alloy is desired, the mixture may be processed once or
repeatedly at a temperature above said melting temperature for a
time sufficiently long for a fraction of said graphite particles
reacting with at least one metal component of said alloy to form
metal carbide particles. In this case, it is advantageous if the
graphite powder initially dispersed into the alloy has a bimodal
size distribution of the graphite particles, with a fraction of
particles smaller than about 25 micrometers, preferably smaller
than about 10 micrometers, and another fraction of particles larger
than about 25 micrometers.
BRIEF DESCRIPTION OF THE DRAWINGS
[0055] The invention will be described in more detail in connection
with exemplary embodiments illustrated in the drawings, in
which
[0056] FIG. 1 shows an optical microscopy image of a
graphite-particle-reinforced BMG composite with 5 vol. %
graphite;
[0057] FIG. 2 shows differential scanning calorimetry (DSC) scans
of monolithic Vit 105 and BMG-graphite composites with varying
graphite particle volume content;
[0058] FIG. 3 shows x-ray diffraction (XRD) scans of BMG-graphite
composites with 3.5 vol. % graphite reinforcement produced at
different casting temperatures;
[0059] FIG. 4 shows compressive stress-strain curves and hardness
(inset) for as-cast 3 mm graphite-reinforced Vit 105 composites of
various reinforcement contents;
[0060] FIG. 5 shows compressive stress-strain curves for as-cast 3
mm graphite-reinforced Vit 105 composites with optimized size and
process parameters;
[0061] FIG. 6 shows a comparison of strength and plasticity derived
from FIGS. 4 and 5 to literature values;
[0062] FIG. 7A-7E shows scanning electron microscopy (SEM) images
of fracture surfaces and particle-shear band interactions;
[0063] FIG. 8 shows XRD scans for three samples designated as
sample 1 to 3, having 5 vol. %-graphite reinforcement and
increasing interfacial carbide content;
[0064] FIG. 9 shows stress-strain curves for samples 1 to 3;
[0065] FIGS. 10A and 10B show optical microscopy images of samples
1 and 3, respectively, illustrating different extents of carbide
formation along the matrix-particle interface;
[0066] FIG. 11 shows a diagram illustrating the hardness of
composites processed at 0.35 kW and 2.1 kW power input compared to
the monolithic matrix material;
[0067] FIG. 12 shows an optical micrograph of a graphite/ZrC
reinforced BMG;
[0068] FIG. 13 shows XRD scans of three kinds of composites
described above;
[0069] FIG. 14 show a schematic setup for measuring tribological
properties;
[0070] FIG. 15 shows an SEM image showing an overview of wear
tracks made at different parameters on a composite sample
containing 8 vol. % graphite;
[0071] FIG. 16 shows XRD scans of a composite with 4 vol. %
graphite and 3 vol. % graphite with in situ formed ZrC before and
after tribology testing;
[0072] FIG. 17 shows a diagram illustrating the COF in dependence
of graphite content for composites with and without in situ ZrC
formation. For the composites with in situ ZrC formation the COF in
the upper and lower regime is shown.
DETAILED DESCRIPTION OF THE INVENTION
[0073] For achieving the goal of obtaining improved BMG-based
materials, it is believed that foreign-particle reinforcement has
the brightest future, because it allows easy reproducibility and
direct tailoring of material properties.
Foreign-particle-reinforced BMGs display, for example, better
reproducibility of microstructure than in-situ-formed composites
because the reinforcement microstructure and volume content are
independent of processing parameters, in particular cooling rate.
Similarly, while porous BMGs display a combination of high
plasticity and yield strength, achieving a homogeneous pore
distribution is very difficult. Monolithic BMGs with high poison
ratios also appear promising, but the effect of enhanced plasticity
has only been observed so far in a very costly Pt-based alloy.
Foreign-particle reinforcement also has the great advantage that
the microstructure and thus the material properties can be
tailored. The latter can be adjusted by type, shape, size and
volume fraction of the reinforcement particles, as it is
state-of-the-art in crystalline metal-matrix composites (MMCs)
Foreign-particle reinforced BMGs also display high reproducibility
because they can be processed by standard MMC processing techniques
followed by a rapid quenching step. By using reinforcement
particles in the micrometer range, the heterogeneous nucleation
surface can be minimized, so that a high critical casting thickness
is still achieved with today's good glass-formers.
[0074] In the following, a new class of foreign-particle-reinforced
BMGs is discussed, where a fully amorphous alloy matrix was
reinforced with graphite particles. In the following examples,
Zr.sub.52.5Cu.sub.17.9Al.sub.10Ni.sub.14.6Ti.sub.5 (Vit 105) was
employed. The reinforcing particles had sizes of 25-44 .mu.m and,
in some cases, 44-75 .mu.m. Of course, the invention is in no way
limited to this base BMG or these particle sizes. A plasticity of
up to 18.5% was achieved without sacrificing the high yield
strength (1.85 GPa) of the metallic glass. Its microstructure can
be reproduced easily and is independent of cooling rate. This novel
composite displays the highest combination of yield strength and
ductility reported so far for foreign-particle-reinforced BMGs, and
the mechanical properties remain favorable, even if compared to
amorphous alloys or composites produced by the other methods
discussed above.
Characterization of Graphite-Reinforced BMGs
[0075] FIG. 1 shows, as an example, the graphite particle
distribution in the Vit 105 matrix for a composite containing 5
vol. % graphite at 25-44 .mu.m particle size, as obtained by
induction mixing. The particles are homogeneously distributed in
the glassy Vit 105 matrix and have shapes ranging from rectangular
to circular.
[0076] FIG. 2 shows DSC scans of monolithic Vit 105 and of the
composites with various reinforcement volume fractions ranging from
5 to 20 vol. % at 25-44 .mu.m particle size. A comparison of the
crystallization enthalpy of the composites with that of monolithic
Vit 105 shows that the matrix material is fully amorphous. The
addition of graphite, however, shifts the onset of crystallization
to a higher temperature, i.e. the composite has a higher thermal
stability than the monolithic metallic glass, and the
crystallization behavior changes. The first crystallization peak
increases with increasing graphite content at the expense of the
second crystallization event.
[0077] FIG. 3 shows XRD scans for a Vit 105 composite with 3.5 vol.
% graphite at 25-44 .mu.m particle size at different casting
temperatures. The two clearly seen amorphous humps result from the
glassy Vit 105 matrix, while can be attributed to crystalline ZrC.
While only traces of ZrC are observed in the lower scan, the ZrC
content increases significantly with increasing casting
temperature. The graphite peaks are not visible in the XRD scans
because carbon is too light to be detected compared to the other
elements present. It was found by energy-dispersive x-ray
diffraction (EDX), however, that no graphite particles had fully
transformed into carbides and that the content of Zr and Ti in the
matrix were within 0.5% of the nominal composition. Thus, the
carbides observed in XRD must be due to interfacial reaction
between the matrix material and the reinforcement particles.
Further evidence for this important observation will be provided
further below.
[0078] Compression tests conducted on the composites with the
lowest possible interfacial carbide formation show a large
improvement in plasticity, with only a slight decrease in yield
strength compared to monolithic Vit 105. FIG. 4 shows that the
plastic region has strongly increased, from 3% for monolithic Vit
105 to about 7% for 3.5 vol. % graphite, 13% for 5 vol. % graphite,
and 15% for 10 vol. % graphite, whereas the yield strength has
decreased only slightly from 1.85 GPa for the monolithic alloy to
1.7 GPa for 3.5 vol. % graphite, 1.6 GPa for 5 vol. % graphite, and
1.5 GPa for 10 vol. % graphite reinforcement, each time at 25-44
.mu.m particle size.
[0079] Further experiments have shown that an even higher
plasticity with only a slight reduction in yield strength can be
achieved by a variation in particle size. FIG. 5 shows the
stress-strain behavior of a composite containing 3.5 vol. %
graphite particles with a particle size of 44 to 75 .mu.m, where
particular care was taken to minimize the thickness of the
interfacial carbide layer. A plasticity of 18.5% in combination
with a yield strength of 1.85 GPa was achieved. Further
optimization appears possible.
[0080] For the samples with minimal carbide formation, the hardness
(measured as Vickers Hardness HV30, see standard DIN EN ISO 6507)
decreases with increasing graphite volume content, as is shown in
the inset to FIG. 4. Even small reinforcement volume fractions of
5% lead to significant softening of the material, and the hardness
decreases by about 25% for graphite contents of .gtoreq.10 vol. %.
On the other hand, composites displaying more carbides in XRD
displayed a higher hardness of up to 550 HV30.
[0081] FIG. 6 shows the yield strength and plasticity of the 5 vol.
% (25-44 .mu.m particle size), 10 vol. % (25-44 .mu.m), and
optimized (3 vol. %, 44-75 .mu.m) graphite-reinforced BMG
composites in comparison to other particle-reinforced BMG
composites found in literature (accuracy of literature values:
.+-.10%). Apparently, the graphite-reinforced BMG composites
represent a step improvement in their combination of fracture
strength and plasticity.
[0082] FIGS. 7A-7E show representative SEM images of fracture
surfaces and particle-shear band interactions for these
graphite-BMG composites (FIGS. 7A to 7D: 25-44 .mu.m particle size;
FIG. 7E: 44-75 .mu.m particle size). FIG. 7A shows a fracture
surface where a high density of vein patterns in the Vit 105 matrix
is observed around a graphite particle: a further proof that the
matrix is fully amorphous. The image in FIG. 7B shows how the
graphite particle obstructs the flow of the matrix material from
the top left to the bottom right during deformation (final
deformation event). FIG. 7C displays shear bands and steps on the
outer surface of the compression samples after failure (the
fracture surface is on the left side of the image), while FIG. 7D
shows particle-shear band interactions on the surface of a
compression sample after failure. The primary shear-band spacing
around the particles is in the range of 1-5 .mu.m, as can be
concluded from FIGS. 7C and 7D. FIG. 7E illustrates the high shear
band density in the matrix achieved with larger graphite particles
in the range of 44 to 75 .mu.m.
Interfacial Carbide Formation
[0083] A further object of investigation was the effect of
interfacial carbide formation on the mechanical properties. Zr
being the element in the matrix with the most negative enthalpy of
formation with graphite (H.sub.for=-106 kJ/mol followed by -77
kJ/mol for Ti), it is expected that zirconium carbide (ZrC) forms
on the surface of the graphite particles. It has previously been
reported that smaller graphite particles of size below about 10
.mu.m completely transform to ZrC in Zr based BMGs, forming an
in-situ composite. In contrast, in the present invention, complete
transformation of all graphite particles is expressly avoided.
[0084] Three samples designated as sample 1, 2 and 3 were prepared
with 25-44 .mu.m graphite particles and various amounts of
interfacial carbides, induced by casting the composites at
different temperatures. Samples 1, 2 and 3 were heated at a setting
of 1, 2.5 and 4, respectively, on the Buhler MAM1 system
(corresponding to 0.35 kW, 0.9 kW and 2.1 kW of power input).
[0085] The changes in carbide content with increasing casting
temperature can be seen in the XRD scans shown in FIG. 8. The two
clearly seen amorphous humps result from the glassy Vit 105 matrix.
DSC also confirmed the glassy structure of the matrix. The Bragg
peaks can be attributed to crystalline ZrC. While only traces of
ZrC are observed in the lower scan, the ZrC content increases
significantly with increasing casting temperature. The graphite
peaks are not visible in the XRD scans because carbon is too light
to be detected compared to the other elements present. EDX,
however, proved that no graphite particles had fully transformed
into carbides and that the content of Zr and Ti in the matrix were
within 1% of the nominal composition. Thus, the carbides observed
in XRD must be due to interfacial reaction between the matrix
material and the reinforcement particles. It may be concluded that,
by adjusting the power input in the final arc-melting step, it is
possible to vary the carbide content of the BMG composites. Thus,
in addition to the common methods of tailoring the mechanical
properties of metal-matrix composites such as varying particle
size, shape, hardness and volume fraction, the mechanical
properties of the presently proposed Bulk Metallic Glass composites
can be varied by merely varying the processing parameters.
[0086] FIG. 9 shows the results of compression tests conducted on
the three samples. Sample 1, with the lowest carbide content,
displays the highest plasticity, whereas sample 3, where most of
the carbide has formed, shows brittle fracture behavior.
[0087] FIGS. 10A and 10B show optical microscopy images of samples
1 and 3, respectively. In FIG. 10A, for sample 1 which was
processed at the very low power setting of 0.35 kW, only a very
thin interfacial reaction layer is visible, which has been broken
up by polishing the sample. In FIG. 10B, for sample 3 which was
processed at a high power setting of 2.1 kW, a distinctive reaction
layer with a thickness of about 1.5-2 .mu.m can be seen at the
graphite-matrix interface. This reaction layer was still mostly
intact after polishing and was thick enough to be identified as ZrC
by EDX. Clearly, even in this sample, the graphite particles have
not completely transformed into carbides, and a carbide layer
surrounding the graphite particles is found. In the other samples,
this layer is in the submicron range. The interfacial carbide phase
seen in FIG. 10B appears to be responsible for the brittle fracture
behavior of sample 3. On the other hand, the samples with higher
carbide content display a higher hardness than the samples where
only a little carbide has formed. Sample 3 in FIG. 10B, for
example, displayed a hardness of 476 HV30 (comparable to the
monolithic alloy), while sample 1 showed a hardness of 432
HV30.
[0088] The effect of the carbide layer on the hardness of the
composites is shown in FIG. 11. Graphite-BMG composites processed
at low power setting (0.35 kW) resulting in minimal carbide
formation display strong softening with increasing graphite volume
fraction. Samples processed at higher power setting (2.1 kW) with a
thicker carbide layer display significantly higher hardness than
composites with minimal carbide formation. At volume contents up to
5% the composites processed at 2.1 kW display even higher hardness
than the pure matrix material.
Three-Phase Composites: Tribological Properties
[0089] In this section, novel three-phase graphite/ZrC reinforced
BMGs are discussed. The tribological properties of these BMG
composites are compared to those of graphite-reinforced BMGs, as
discussed above, monolithic BMG and commercial bearing steel. As a
typical example, Vit 105
(Zr.sub.52.5Cu.sub.17.9Ni.sub.14.6Al.sub.10Ti.sub.5) was again used
as a base alloy for the BMGs; however, the invention is not limited
to this base alloy.
[0090] One way of changing the tribological properties of a system
is by changing the contact surface on a microscopic scale. The
contact surface can be significantly influenced by adding second
phase particles with different hardness than the amorphous matrix.
Graphite as a reinforcement phase is promising for optimizing
tribological properties because of its superlubricity and the
above-discussed possibility of in-situ formation of very hard ZrC
particles in Zr-based BMGs.
[0091] As will be seen in the following, the COF of the amorphous
alloy can be decreased significantly by adding the reinforcement
phases. The low COF combined with very high compressive yield
strength (.about.1.8 GPa) of these novel composites make them
potential candidates for self-lubricating friction bearing
materials.
[0092] As already discussed above, it is possible to tailor the
microstructure of the graphite-reinforced BMG composites by
adjusting the processing parameters. Three types of BMG composites
were thus produced. "Standard" graphite-reinforced BMGs as
described above, which show a clean graphite-matrix interface with
very little ZrC formation, was already discussed and illustrated in
FIG. 10A. Increasing the processing temperature in the final
casting step leads to a significant ZrC layer in the
particle-matrix interface with a thickness of about 2 .mu.m, as
also discussed above and shown in FIG. 10B. If such a sample is
remelted several times or smaller graphite particles are added, ZrC
crystals also appear in the matrix, leading to a three-phase
composite. This novel composite is illustrated in FIG. 12, where
the arrow 121 indicates a graphite particle surrounded by a ZrC
layer, and arrow 122 indicates a ZrC particle.
[0093] XRD scans of the three types of composites, all displaying
an amorphous background signal from the matrix, can be seen in FIG.
13. The figure shows XRD scans of the three kinds of composites,
all with 7 vol. % graphite, namely (from the bottom up), Vit
105-graphite composite, composite with interfacial ZrC formation
and three-phase composite with both interfacial ZrC formation and
ZrC particles in the matrix. The composites with interfacial ZrC
and additionally ZrC in the matrix both show ZrC peaks of similar
intensity compared to the standard graphite-reinforced BMG sample
which displays almost no carbide formation. Samples displaying no
carbide formation and the three-phase composites were used for
tribological testing.
[0094] FIG. 14 illustrates the setup used for tribological testing.
A steel ball is moved in circles over the sample surface with a
predetermined pressing force (load), thus creating a wear
track.
[0095] An SEM image of a sample containing 8 vol. % graphite is
shown in FIG. 15 after tribological testing. This SEM image shows
an overview of wear tracks made at different parameters, where the
first parameter indicates the vertical load and the second
parameter indicates the number of revolutions of the steel ball. A
homogeneous particle distribution is visible, which was found in
all composites.
[0096] XRD scans performed before and after tribological testing on
the sample plates of both types of BMG composites displayed no
significant changes, as can be seen in FIG. 16. One must however
consider that only about 5% of the sample surface was affected by
the tribology tests.
[0097] A comparison of coefficient of friction (COF) and wear trace
depth for monolithic samples and composites in 1000 revolution
tests conducted at a load of 1 N was performed. Amorphous Vit 105
displayed a much higher COF and much higher fluctuations of COF
than the fully crystallized alloy. The COF of the amorphous sample
dropped slightly until it stabilized after about 300 revolutions,
whereas the crystalline sample displayed a constant COF throughout
the test. All the composites displayed two significant levels of
COF. At the beginning of the test, they showed a stable COF which
is significantly lower than in the monolithic matrix material.
After >100 revolutions they jumped to an even lower level of COF
where some composites stay whereas other make jumps back up to the
higher level where they stay for <100 revolutions. This behavior
was especially prominent in samples with significant ZrC content.
It was also found that a higher reinforcement content led to less
fluctuations in COF. All samples tested for 1000 revolutions showed
a more or less linear increase in wear track depth except for the
amorphous monolithic Vit 105 which exhibited significant pin
lifting.
[0098] At a force of 1 N the amorphous monolithic alloy showed a
COF of about 0.8 which is comparable to the value measured for the
hardened bearing steel (0.78). Crystallized Vit 105 displayed a COF
of 0.6. It was found that the reinforcement of the monolithic glass
with very low volume contents of graphite leads to a significant
decrease in COF as can be seen in FIG. 17. Composites with
additional ZrC reinforcement show an even larger decrease in COF
especially in the lower COF regime seen in the 1000 revolution
runs. No significant: difference in the effect of graphite
reinforcement was found for tests run at a 5 N load for 100
revolutions; however, ZrC did not lead to quite as large of a
decrease in COF as for samples run at a 1 N load.
[0099] Some shear bands were found at the edge of the wear traces
after 1000 revolutions at 1 N. The shear bands run about 25.degree.
to the sliding direction and give evidence for high enough stresses
to lead to inhomogeneous flow. In some composite samples, smeared
matrix material was found in the wear trace. This is expected to
come from deformation in the undercooled liquid region.
[0100] In general very little smearing of graphite was observed in
the composite samples. It was found by SEM investigation that whole
graphite particles were ripped out of the matrix during the wear
tests between 100 and 1000 revolutions. Composites samples
containing ZrC showed several channels in their wear tracks. The
depth of the channels was estimated to be about 3 .mu.m. On the
steel tip of a composite sample containing ZrC, lots of small
particles with a particle size between 50 and 500 nm were found on
the surface, which are thought to be ZrC debris.
[0101] Comparing the width of the wear tracks and pin depths gives
a very qualitative approximation of the wear rate. The thinnest
wear track after 1000 revolutions at 1 N was found for crystalline
Vit 105, the thickest for the hardened steel with about 50 .mu.m
and 200 .mu.m respectively. No significant difference was seen
between the amorphous alloy and the composites, which all displayed
trace widths of about 120 .mu.m. The depths of the wear tracks are
difficult to compare because of the observed pin lifting phenomena.
It was found that the wear tracks of the BMG composites after 1000
revolutions at a load of 1 N are slightly less deep than in the
bearing steel. The depths and widths of the wear tracks do not
correlate with the hardness of the monolithic materials which are
846 HV for the bearing steel, 547 HV and 478 HV for crystalline and
amorphous Vit 105 respectively.
Discussion
[0102] In discussing the above results, it should first be
emphasized that a homogeneous particle distribution, as shown in
FIG. 1, might have been only achievable because of the good wetting
behavior between graphite and Vit 105 and due to the multiple-step
induction mixing procedure proposed. Preliminary attempts to
produce composites with non-wetting particles indeed led to
particle agglomerations. As for the DSC results (FIG. 2), the
addition of graphite apparently improved the thermal stability of
the composites compared to the monolithic alloy, such that these
composites may also be used for superplastic forging. The
phenomenon of improved thermal stability was also observed for SiC
reinforcement particles in a Zr-based BMG composite. The effect may
be due either to the change in thermal conductivity compared to the
monolithic BMG or because of a slight shift in matrix composition
resulting from the interfacial carbide reaction. The latter may
also be responsible for the slight change in crystallization
behavior.
[0103] The combination of fracture strength and plasticity found in
the present composites appears to be the highest ever recorded for
foreign-particle-reinforced BMGs, as is shown in FIG. 6. While
similar plasticity has been seen in 50% Nb-reinforced Zr-based
BMGs, in those alloys the yield strength dropped drastically to 30%
of the strength of the monolithic alloy due to the high
reinforcement volume fraction deployed. Our key to success was the
use of "soft" reinforcement particles, i.e. a reinforcement
material with a much lower Young's modulus (ca. 15 GPa in graphite)
than that of the Vit 105 matrix (Young's modulus, Ez.apprxeq.100
GPa). In contrast, all the other reinforcement materials shown in
FIG. 6 have a higher Young's modulus than the matrix material: for
the refractory metals Nb, Ta, and Mo, for example, E ranges from
105 to 327 GPa. The lower Young's modulus of the graphite leads to
local compressive stress concentrations in the matrix material
close to the particle-matrix interface, while tensile stresses in
the matrix are expected to occur for "hard" particle reinforcement.
Thus, on the one hand graphite may act as a typical reinforcement
particle, splitting shear bands (such splitting and the
particle-shear band interaction is shown in FIGS. 7C and 7D); on
the other hand, the graphite particles are expected to halt the
propagation of shear bands by reducing the stress at their tips
when they run onto the soft material. This is actually shown in
FIGS. 7A and 7B, where the reinforcement particle clearly hinders
the matrix flow during deformation. Because of their low strength
the graphite particles may act in a way similar to pores in
amorphous alloys. During compression testing the first shear band
may be initiated as soon as the stress in the "soft" particle (or
pore) reaches a critical value. After initiation of this shear band
the stress around this particle decreases, while other shear bands
are initiated at the particles which reach critical stress
concentrations. Thus, multiple shear bands nucleate, run through
the material and cross, and thus hinder, each other--which leads to
enhanced plasticity. These results--enhanced plasticity and
strength in combination with reproducibility of microstructure--are
expected to have a significant influence on the field of bulk
metallic glass strengthening, as did earlier results in the field
of nanostructured metals. There, an improvement was achieved with
the development of a twophase material of micrometer-sized grains
embedded in a matrix of grains with sizes <300 nm.
[0104] Apparently the combination of the above mentioned effects
(shear-band splitting, impairment of propagation, and shear-band
initiation) generates a great increase in plasticity at very low
graphite content, which in turn leads to only minimal decrease in
yield strength compared to the monolithic alloy. The regions around
the particles display a very small shear-band spacing in the
micrometer or even sub-micrometer range (FIGS. 7A to 7E), in
contrast to what has been reported for monolithic Zr-based BMGs.
Indeed, a direct correlation between reduced shear-band spacing and
enhanced plasticity has been reported for metallic glass ribbons of
varying thickness and BMG composites. As can be seen in FIG. 4,
there is no significant plasticity benefit if the graphite content
is increased from 5 to 10 vol. %. Once the inter-particle distance
is small enough to generate a homogeneous high shear-band density
in the matrix during deformation, more reinforcement particles will
not further improve plasticity in any significant way. Indeed,
doubling the reinforcement concentration will reduce the
inter-particle distance by only about 1/3.
[0105] The foreign-particle reinforced composites also provide the
advantage that their mechanical properties can be tailored by
tuning the carbide formation. As can be seen in the XRD scans of
FIGS. 3 and 8, the amount of graphite that trans-forms into ZrC can
be adjusted by altering the casting temperature. To prevent brittle
fracture behavior, however, it may be beneficial to keep the
carbide content fairly low, and the Zrc content only appears
dominant in the XRD scans because of Zr's high atomic mass. On the
other hand, EDX shows that the graphite particles did not transform
fully into carbides, and the optical micros-, copy image in FIGS.
10A and 10B provides evidence that only an interfacial carbide
layer formed even at a high casting temperature. Indeed, if much
carbide had formed in the composites, the matrix composition would
have shifted and glass-forming ability would have decreased. The
softening of the composite with increasing reinforcement volume
fraction (inset to FIG. 4) also suggests that only a very small
fraction of the graphite reacted to carbide. A similar minor
carbide formation in the matrix-particle interface has previously
also been observed In carbon-fiber-reinforced Zr-based BMG
composites. In contrast, however, most other studies have shown a
complete transformation of the graphite particles into ZrC. For
example, carbon or graphite particles of .ltoreq.10 .mu.m have been
used to process BMG-ZRC composites in situ, and a distinct increase
in hardness has been observed. This apparent contradiction of
carbide formation, or lack of it, can be explained by the fact that
larger graphite particles of .gtoreq.25 .mu.m were used in this
study, and special care was taken to heat moderately during
processing.
[0106] In the present case the carbides start growing in the
matrix-particle interface forming a hard shell around the graphite
particles. Interfacial carbide formation is favored because of the
short diffusion paths necessary. The formed ZrC layer acts as a
diffusion barrier and slows the carbide formation in the interface
controlling the reaction. It is expected that strong stirring of
the meft could lead to complete reaction of graphite to ZrC because
the evolving ZrC would be separated from the graphite allowing
further carbide reaction in the interface. The thin interfacial
carbide layer leads to a significant increase in hardness compared
to the standard graphite composites. If the graphite particles are
considered as spheres of 35 .mu.m and the graphite layer as an
interfacial layer of 1.5 .mu.m, composites containing 5 vol. %
graphite contain less than 0.7 vol. % ZrC. Due to this layer, an
increase in hardness of about 16% is observed, compared to the
composite with 5 vol. % graphite and minimal carbide formation.
This phenomenon cannot be explained by Ashby's rules of mixing
because of the geometrical particularities of the carbide
surrounding the graphite particles. If the graphite particle with
the hard carbide shell around it is considered as a monolithic
reinforcement particle, it will display similar mechanical
properties like a hard-boiled egg. At low stress, it will be very
stiff. If higher stress is applied, the shell will break and it
will act like a soft particle. The stress value necessary to "crack
the shell" is of course determined by the thickness of the carbide
layer but also by the shape and size of the graphite particle. At
very low reinforcement, graphite leads to a strong decrease in
hardness whereas the composites with the carbide layers show a
slight increase, as one would expect for a matrix reinforced with
hard particles. At higher reinforcement contents, the hardness of
the composites with an interfacial carbide layer also starts to
decrease and the soft graphite seems to become dominant. It is
expected that when hardness testing is conducted on samples with
low graphite volume content the hard particles are pushed into the
soft matrix relieving stress on the particles, whereas at high
volume contents they hinder each other and are exposed to enough
stress to crack the carbide shells and the graphite becomes
dominant.
[0107] The thickness of the interfacial carbide layer also has a
significant influence on the stress-strain behavior of the
composites. Soft graphite particles with a low Young's modulus are
favorable in achieving high plasticity in compression. The thicker
the interfacial carbide layer around a graphite particle, the more
it will act like a hard ceramic particle instead of a soft graphite
particle. As apparent from FIG. 9, the thicker the interfacial
carbide layer, the more brittle the material becomes. The carbide
layer that has a Young's modulus of about 400 GPa compared to 100
GPa of the matrix material leads to tensile stress concentrations
close to the particle matrix interfaces in such a way that
propagating shear bands are led around the reinforcement particles,
hindering shear band-particle interaction. In addition, the high
hardness of the carbide layer (about 2500 HV compared to 15 HV of
graphite) hinders the absorption of approaching shear bands but
deflects them, leading to fracture on one or few bands.
[0108] If plasticity is desired to be maximized, the carbide layer
should therefore be kept as thin as possible. Even in samples
processed at the lowest possible energy input where casting is
still possible, some ZrC was detected in XRD. While it might thus
be impossible to fully eliminate the carbide layer, it is still
possible to weaken the strength of a carbide shell by increasing
the particle size. If once again the particles are approximated as
a soft sphere with a hard shell around It, a shell of the same
thickness will carry less load if the sphere is larger. As can be
seen in FIG. 5, larger graphite particles lead to very high
plasticity at low reinforcement volume fractions. Samples with 3.5
vol. % particles of 25-44 .mu.m display only about 7% plasticity
compared to up to 18.5% with 45-75 .mu.m particles. This low volume
content of graphite leads to only a slight reduction of yield
strength compared to the monolithic alloy.
[0109] It was also shown above that graphite reinforcement of Vit
105 leads to a significant change in tribological behavior. The
reinforcement reduces stick-slip and leads to a significant
decrease in COF. Graphite is known to lower the COF of tribological
partners in metallic and polymer-based materials. In metallic
systems, graphite sticks especially well on oxidized surfaces (as
are present on the surface of Zr-based BMGs), which, on a
microscopic scale, can lead to graphite sliding on graphite, which
displays a very low COF. In addition it was also observed in the
present study that debris was pushed into the soft graphite
particles or the holes from ripped out particles, which is expected
to lead to less abrasive wear and additionally lower the COF.
[0110] It was further observed that the newly developed three-phase
composite shown in FIG. 12 resulted if samples were remelted and
suction cast multiple times. This probably led to ZrC breaking off
the graphite particles and distributing itself in the melt. An even
more prominent drop in the COF was found for these three-phase
composites especially in their lower regime of COF (see FIG.
17).
[0111] The jumps in COF observed in the composites containing ZrC
are clear evidence for two different sliding mechanisms. In the
regime of high COF strong fluctuations of COF are observed as is
common for the monolithic matrix material and also composites
reinforced only with graphite. This leads to expect a similar wear
mechanism as in the graphite-reinforced composites with the
exception of the effect of the hard particles in the matrix.
However, once the COF drops to the lower regime the fluctuations of
COF also decrease drastically, which, in combination with the
observed channel-like morphology of the wear track, gives evidence
for a new sliding mechanism.
[0112] The channel-like morphology of the wear tracks observed in
composites with in situ formed ZrC stands in contrast to the
relatively smooth wear tracks found in graphite-reinforced Vit 105
or the monolithic matrix material. Due to the very high strain
rates achieved on the micro-scale during sliding (.about.10.sup.5
s.sup.-1) it is unlikely that the channels are formed by
inhomogeneous deformation of the matrix material but much more by
local abrasion of the matrix material by ZrC debris. Very small
particles which are expected to be ZrC debris were found on the
steel ball used for such tests. It is expected that larger ones
were also present but fell off due to their lower surface-to-weight
ratio. Once a shallow channel has formed, debris will remain in the
channel and lead to local abrasion deepening it.
[0113] These findings indicate that the wear behavior of this
three-phase composite is based mainly on the geometrical
particularities of the wear traces. A possible explanation for the
observed tribological particularities may be that during the low
regime of the COF debris is pushed in front of the steel ball,
leading to the observed channels in the wear tracks. A steady state
is achieved where the steel ball slides on top of the channels,
leading to a slightly lower COF than in composites reinforced only
with graphite, due to the smaller contact surface leading to less
adhesive friction. The transition from the high to the low regime
of COF takes place within very few revolutions and is accompanied
by a significant increase in pin height of about 2 .mu.m. This is
thought to be due to the channels filling up with debris and the
steel ball being lifted onto the debris and sliding on top of it.
The debris, which seems to be quite round, is thought to reduce the
frictional forces by rolling underneath the steel ball in the wear
channels. Jumps back up to the higher regime of COF may take place
when a graphite particle is ripped out of the wear track and large
amounts of debris is pushed into the hole which leads to the ball
dropping back down onto the channel walls and sliding on them.
[0114] The fact that very low reinforcement contents are sufficient
for lowering the COF in both composites gives further evidence that
the mechanisms are not based on the static pairing of the two
materials but much more on the influence of debris on the dynamic
contact during sliding. In the one case graphite acts as a
lubricant and debris trap, in the other, ZrC leads to channel
formation changing the topography of the wear track. Especially
once the wear track has reached a certain depth, debris will stay
in the track and not be pushed out over the edge.
[0115] In contrast to previous accounts, no hint for
crystallization of the wear track or debris during wear was found;
however, some smearing of matrix material was observed. This
smearing is expected to take place in the undercooled liquid
region, which is quite large in the matrix material. Due to the
very low sliding speed the smeared material is probably cooled by
heat transfer to the bulk of the sample before the next revolution
takes place. This cooling is thought to be fast enough to hinder
local crystallization. The local stress on the sample during the
tribology tests can be quite large. If one considers the contact
surface to be a circle with the diameter of the wear track (120
.mu.m) the global stress would be about 90 MPa at a 1 N load. If
one, however, considers that very small particles of hard debris
might be between the ball and the sample the contact surface will
decrease significantly leading to very high local stress which
could easily be above the flow stress of the matrix material (about
1.9 GPa) which would also explain the observed local shear banding
on the edge of the wear track.
[0116] As far as wear rate is concerned, the present observations
are very qualitative. However, the width of the wear tracks and the
measured depths of the wear traces give evidence that the wear rate
of the composites is lower than in amorphous Vit 105 and commercial
bearing steel in the used testing set up. This is thought to be at
least partially due to the self-lubricating effect of graphite and
the holes of torn out graphite particles acting as traps for
debris, both leading to less abrasive wear.
CONCLUSION
[0117] In conclusion, the BMG-graphite composites developed in this
study constitute a very promising material for structural
applications due to their high plasticity, comparable to that of
crystalline alloys, combined with the high yield strength typical
of metallic glasses. The matrix-particle interface, particularly
its hardness, has a major influence on the mechanical properties of
these composites. Since the microstructure of these
foreign-particle reinforced composites can be tailored and easily
reproduced for specific applications, one may expect that these new
composites will have a great impact on research efforts in the
entire field of amorphous structural materials.
[0118] Furthermore, the tribological properties of
graphite-reinforced and of a newly developed graphite- and
carbide-reinforced BMGs were compared to those of amorphous and
crystalline alloys as well as bearing steel. It was found that
graphite and especially carbide reinforcement leads to a
significant decrease in COF. The carbide-containing composites
displayed two regimes of COF. The very low COF found in the lower
regime is thought to be due to the geometry of the wear tracks
formed by carbide particles. In this wear regime the COF is up to
four times lower than in the monolithic alloy. Crystallization of
the wear tracks or debris after tribological testing was not
observed. A qualitative comparison of the wear rate gives evidence
that the newly developed composites may show even lower wear rates
than the 100Cr6 bearing steel used as a reference material. These
tribological properties combined with the high yield strength of
the composites make them an interesting candidate for a dry
frictional bearing material.
[0119] While the above examples related to composites based on Vit
105, it is to be expected that similar results can be achieved also
with composites based on BMGs with a different composition, and the
invention is in no way limited to composites based on Vit 105.
While a homogeneous particle distribution is more easily achieved
if the wettability of the particle surface by the matrix alloy is
good, composites can also be produced for which wettability is
poor. Likewise, it is not necessary that a reaction can occur
between the matrix alloy and graphite. However, the above-described
experiments suggest that Zr-based glass-forming alloys are
particularly good matrix materials for the composites of the
present invention.
Methods
(a) Sample Preparation
[0120] Pre-alloys with the atomic composition
Zr.sub.52.5Cu.sub.17.9Ni.sub.14.6Al.sub.10Ti.sub.5 (Vit 105) were
prepared in a Buhler AM system by arc melting the high-purity
elements (>99.95%) in a 300 mbar Ar 6.0 atmosphere and casting
the molten alloy ingot into a Cu mould of 13 mm in diameter and 40
mm in length. The subsequent composite preparation took place in a
1200 mbar Ar 6.0 atmosphere. 2-20 vol. % conducting-grade graphite
with a particle size of 25-44 .mu.m or 44-75 .mu.m was mixed with
the matrix material by induction melting of the alloy on top of the
graphite powder in a water-cooled silver boat. After the powder was
picked up the sample was remelted in the silver boat to achieve a
homogeneous particle distribution. The crystalline composites were
then suction-cast into 3 mm rods with a length of 30 mm (for
compression testing, thermophysical characterization and imaging)
or into 2 mm.times.7 mm.times.30 mm plates (for tribology
measurements) in a Buhler MAM1 arc melter. 5-mm-long slices were
cut from the 3 mm rods for compression testing. Thinner slices were
cut for thermophysical invesligation. Tribology samples were first
ground and then polished with a 0.05 .mu.m Al.sub.2O.sub.3
dispersion.
[0121] Standard samples were cast at an arc power setting of 1
(corresponding to 0.35 kW power input), while a setting of 2.5 (1
kW) and 4 (2.1 kW) was used to induce interfacial ZrC formation. If
necessary, samples were remelted several times to initiate a
stronger carbide formation.
[0122] Monolithic BMG samples were prepared without the induction
mixing step, one sample was fully crystallized by annealing at
430.degree. C. for 75 min. Hardened 100Cr6 bearing steel was used
as a reference sample for tribological testing.
(b) Structural and Thermophysical Characterization
[0123] XRD was performed on polished samples with a PANalytical
X'Pert diffractometer using Cu-K.sub..alpha. radiation. A Seiko DSC
220CU system and a Setaram Labsys system were used for calorimetric
analysis. Calorimetric measurements were performed using a sample
weight of approximately 20 mg at a heating rate of 20 K/min. A
CamScan scanning electron microscope (SEM) equipped with a Noran
Energy Dispersive X-ray (EDX) detector was used for elemental
analysis. Samples for optical microscopy were polished with a 0.05
.mu.m Al.sub.2O.sub.3 suspension and etched with a solution of 30
ml HNO.sub.3 in 70 ml of distilled water. A Reichert-Jung Polyvar
Met microscope combined with a Leica camera was used to create the
optical microscopy images.
(c) Mechanical Characterization
[0124] Hardness measurements were performed on a Gnehm Brickers 220
instrument at a setting of HV 30 with an impression time of 6 s.
Compression tests were conducted on a Schenk Trebel tensile tester
combined with Merlin software at a strain rate of 10.sup.-3
s.sup.-1. A high-resolution Zeiss Gemini 1530 FEG scanning electron
microscope was used for microstructure investigation.
(d) Tribological Characterization
[0125] The tribological properties of the material were
investigated on a CETR microtribometer, where the sample was paired
against a bearing steel ball with a diameter of 2 mm at a constant
sliding speed of 100 mm/min without lubrication. All tests were run
at room temperature and a relative humidity of about 40%. After the
ball was run in for 100 revolutions at a 5 N load a first test was
performed with the same parameters, followed by 100, 10 and 1000
(not performed on all samples) revolution tests at a 1 N load. The
ball was run in at a radius of 2.9 mm and the radius was reduced by
0.4 mm for each of the following tests. The high regime of the COF
was determined by linear approximation of the force data obtained
in the 100 revolution tests. In the steel sample the 1000
revolution data was used because COF was not yet in equilibrium
after 100 revolutions due to the oxide layer. In samples displaying
two regimes of COF in the 1000 revolution tests, the lower regime
of COF was determined by averaging the values of the lower shelves.
Graphite volume content of the tribology surfaces was determined by
investigation of optical micrographs with Leica QWin software.
LIST OF ABBREVIATIONS
[0126] BMG bulk metallic glass COF coefficient of friction DSC
differential scanning calorimetry EDX energy-dispersive X-ray
diffraction XRD X-ray diffraction SEM scanning electron microscopy
HV30 Vickers hardness, measured at 30 N impression force
* * * * *