EP1957686B1 - Verbundwerkstoffe aus metallischem massivglas und graphit - Google Patents

Verbundwerkstoffe aus metallischem massivglas und graphit Download PDF

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EP1957686B1
EP1957686B1 EP06775159A EP06775159A EP1957686B1 EP 1957686 B1 EP1957686 B1 EP 1957686B1 EP 06775159 A EP06775159 A EP 06775159A EP 06775159 A EP06775159 A EP 06775159A EP 1957686 B1 EP1957686 B1 EP 1957686B1
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graphite
composite material
particles
phase
alloy
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EP1957686A1 (de
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Jörg F. LÖFFLER
Marco Siegrist
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Eidgenoessische Technische Hochschule Zurich ETHZ
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Eidgenoessische Technische Hochschule Zurich ETHZ
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C1/00—Making non-ferrous alloys
    • C22C1/10—Alloys containing non-metals
    • C22C1/1036—Alloys containing non-metals starting from a melt
    • C22C1/1068—Making hard metals based on borides, carbides, nitrides, oxides or silicides
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C16/00—Alloys based on zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C32/00—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ
    • C22C32/0084—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ carbon or graphite as the main non-metallic constituent
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C45/00—Amorphous alloys
    • C22C45/10—Amorphous alloys with molybdenum, tungsten, niobium, tantalum, titanium, or zirconium or Hf as the major constituent

Definitions

  • 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.
  • BMGs bulk metallic glasses
  • These bulk metallic glasses 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.
  • 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.
  • 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.
  • 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.
  • 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.
  • ZrC zirconium carbide
  • the composite material comprises
  • 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.
  • 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.
  • 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.
  • the particles used in the present invention may have a broad size range. There is lower limit to the size of 1 micrometer; 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.
  • volume fraction of the second (reinforcing) phase is possible. Significant effects are expected if the second phase locally occupies at least 3 volume percent of the total volume of the composite material.
  • 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 between 1 volume percent and 20 volume percent. The most preferred range depends on the average particle size. For larger particles, a lower volume fraction appears to be advantageous.
  • a volume fraction of the second phase in the range between 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.
  • the preferred volume ratio is between 3 and about 6 volume percent.
  • 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.
  • 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.
  • the graphite particles 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.
  • the properties of the second phase shape and size distribution of the graphite particles, volume fraction etc.
  • the shear bands are preferably substantially uniformly (homogeneously) distributed over regions substantially larger than an average graphite particle, preferably over the whole matrix.
  • the matrix alloy in the liquid state, is preferably capable of wetting the particles of the second phase.
  • 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°.
  • 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.
  • 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.
  • 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:
  • 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.
  • a plasticity of up to 18.5% under compression and a yield strength of up to 1.85 GPa has been achieved.
  • 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.
  • 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 ).
  • 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.
  • 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.
  • the graphite particles 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.
  • 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.
  • 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.
  • these particles may be nanocrystals with a mean size below about 1 micrometer.
  • 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.
  • the third phase may comprise carbide particles.
  • 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.
  • the carbide particles consist essentially of zirconium carbide.
  • the carbide particles preferably have a size of less than or equal to about 10 micrometers.
  • such particles can readily be formed in situ by a complete transformation of similarly small graphite particles by a reaction with the matrix alloy.
  • 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.
  • 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 parties 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.
  • 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.
  • 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.
  • the composite material of the present invention may be prepared by various methods.
  • 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:
  • the alloy may be heated above its liquidus temperature by induction melting on top of the graphite powder.
  • 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.
  • 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.
  • 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.
  • porous BMGs display a combination of high plasticity and yield strength, achieving a homogeneous pore distribution is very difficult.
  • Monolithic BMGs with high poisson 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)
  • MMCs metal-matrix composites
  • 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.
  • Figure 1 shows, as an example, the graphite particle distribution in the Vit 105 matrix for a composite containing 5 vol.% graphite at 25-44 ⁇ 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.
  • Figure 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 ⁇ 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 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.
  • FIG. 3 shows XRD scans for a Vit 105 composite with 3.5 vol.% graphite at 25-44 ⁇ m particle size at different casting temperatures.
  • the two clearly seen amorphous humps result from the glassy Vit 105 matrix, while 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. 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.
  • EDX energy-dispersive x-ray diffraction
  • 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 ⁇ m particle size.
  • FIG. 5 show the stress-strain behavior of a composite containing 3.5 vol.% graphite particles with a particle size of 44 to 75 ⁇ 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.
  • the hardness 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 ⁇ 10 vol.%.
  • composites displaying more carbides in XRD displayed a higher hardness of up to 550 HV30.
  • Figure 6 shows the yield strength and plasticity of the 5 vol.% (25-44 ⁇ m particle size), 10 vol.% (25-44 ⁇ m), and optimized (3 vol.%, 44-75 ⁇ m) graphite-reinforced BMG composites in comparison to other particle-reinforced BMG composites found in literature (accuracy of literature values: ⁇ 10%).
  • the graphite-reinforced BMG composites represent a step improvement in their combination of fracture strength and plasticity.
  • Figures 7A - 7E show representative SEM images of fracture surfaces and particle-shear band interactions for these graphite-BMG composites ( Figures 7A to 7D : 25-44 ⁇ m particle size; Figure 7E : 44-75 ⁇ m particle size).
  • Figure 7A shows a fracture surface where a high density of vein patterns in the Vit 105 matrix is observed aroung 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).
  • Figure 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 ⁇ 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 ⁇ m.
  • sample 1, 2 and 3 Three samples designated as sample 1, 2 and 3 were prepared with 25-44 ⁇ 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 Bühler MAM1 system (corresponding to 0.35 kW, 0.9 kW and 2.1 kW of power input).
  • 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.
  • the mechanical properties of the presently proposed Bulk Metallic Glass composites can be varied by merely varying the processing parameters.
  • Figure 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.
  • FIGS 10A and 10B show optical microscopy images of samples 1 and 3, respectively.
  • 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.
  • 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 ⁇ 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.
  • the graphite particles have not completely transformed into carbides, and a carbide layer surrounding the graphite particles is found.
  • 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.
  • 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 displayed a hardness of 476 HV30 (comparable to the monolithic alloy), while sample 1 showed a hardness of 432 HV30.
  • 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.
  • 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.
  • Vit 105 Zr 52.5 Cu 17.9 Ni 14.6 Al 10 Ti 5
  • the invention is not limited to this base alloy.
  • 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.
  • 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 ( ⁇ 1.8 GPa) of these novel composites make them potential candidates for self-lubricating friction bearing materials.
  • XRD scans of the three types of composites 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.
  • 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.
  • FIG. 15 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.
  • 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 .
  • 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.
  • shear bands were found at the edge of the wear traces after 1000 revolutions at 1 N. The shear bands run about 25° 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.
  • 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.
  • Figs. 7A and 7B where the reinforcement particle clearly hinders the matrix flow during deformation.
  • the graphite particles may act in a way similar to pores in amorphous alloys.
  • 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.
  • multiple shear bands nucleate, run through the material and cross, and thus hinder, each other - which leads to enhanced plasticity.
  • the foreign-particle reinforced composites also provide the advantage that their mechanical properties can be tailored by tuning the carbide formation.
  • the amount of graphite that transforms into ZrC can be adjusted by altering the casting temperature.
  • EDX shows that the graphite particles did not transform fully into carbides, and the optical microscopy image in Figs. 10A and 10B provides evidence that only an interfacial carbide layer formed even at a high casting temperature.
  • 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 melt 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.
  • graphite particles are considered as spheres of 35 ⁇ m and the graphite layer as an interfacial layer of 1.5 ⁇ 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.
  • 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.
  • 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.
  • the hardness of the composites with an interfacial carbide layer also starts to decrease and the soft graphite seems to become dominant.
  • 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.
  • 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.
  • 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.
  • 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.
  • 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 ( ⁇ 10 5 s -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.
  • 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.
  • 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.
  • Pre-alloys with the atomic composition Zr 52.5 Cu 17.9 Ni 14.6 Al 10 Ti 5 were prepared in a Bühler 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 ⁇ m or 44-75 ⁇ 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.
  • 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 x 7 mm x 30 mm plates (for tribology measurements) in a Bühler MAM1 arc melter. 5-mm-long slices were cut from the 3 mm rods for compression testing. Thinner slices were cut for thermophysical investigation. Tribology samples were first ground and then polished with a 0.05 ⁇ m Al 2 O 3 dispersion.
  • 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.
  • Monolithic BMG samples were prepared without the induction mixing step, one sample was fully crystallized by annealing at 430°C for 75 min. Hardened 100Cr6 bearing steel was used as a reference sample for tribological testing.
  • XRD XRD was performed on polished samples with a PANalytical X'Pert diffractometer using Cu-K ⁇ 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 ⁇ m Al 2 O 3 suspension and etched with a solution of 30 ml HNO 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.
  • Hardness measurements were performed on a Gappel 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 -3 s -1 . A high-resolution Zeiss Gemini 1530 FEG scanning electron microscope was used for microstructure investigation.
  • 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.

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Claims (23)

  1. Verbundmaterial, aufweisend:
    - eine amorphe erste Phase, die eine kontinuierliche Matrix bildet, wobei die erste Phase aus einer Legierung besteht; und
    - eine zweite Phase, die in der Matrix eingebettet ist, wobei die zweite Phase Graphitteilchen mit einer Grösse im Bereich zwischen 1 und 100 Mikrometer enthält, wobei die zweite Phase zwischen 3 Volumenprozent und 20 Volumenprozent des Verbundmaterials einnimmt.
  2. Verbundmaterial nach Anspruch 1, wobei die Graphitteilchen eine Grösse zwischen 25 und 75 Mikrometer aufweisen.
  3. Verbundmaterial nach einem der vorhergehenden Ansprüche, wobei die zweite Phase derart gewählt ist, dass sie unter kompressiver Deformation des Verbundmaterials bis zum Fliessen eine Verteilung von Scherbändern induziert, wobei um die Graphitteilchen herum die Scherbänder um weniger als ungefähr 5 Mikrometer voneinander entfernt sind.
  4. Verbundmaterial nach einem der vorhergehenden Ansprüche, wobei die Legierung der ersten Phase, wenn sie sich im flüssigen Zustand befindet, in der Lage ist, die zweite Phase zu benetzen.
  5. Verbundmaterial nach einem der vorhergehenden Ansprüche, wobei die Legierung mindestens ca. 40 Atom-Prozent eines Metalls enthält, welches eine negative Bildungsenthalpie für die Reaktion mit Graphit zur Bildung eines Metallkarbids aufweist.
  6. Verbundmaterial nach einem der vorhergehenden Ansprüche, wobei die Legierung mindestens ca. 40 Atom-Prozent Zirkonium enthält.
  7. Verbundmaterial nach einem der vorhergehenden Ansprüche, wobei die Legierung aus Zr52.5Cu17.9Ni14.6Al10Ti5 besteht.
  8. Verbundmaterial nach einem der vorhergehenden Ansprüche, wobei zumindest ein Teil der Graphitteilchen in der zweiten Phase einen Kern aufweisen, der aus Graphit besteht, sowie eine Oberflächenschicht aufweisen, die mindestens ein Metallkarbid enthält.
  9. Verbundmaterial nach Anspruch 8, wobei die Oberflächenschicht eine Dicke von mindestens 100 Nanometer aufweist.
  10. Verbundmaterial nach Anspruch 8 oder 9, wobei die Graphitteilchen eine Grösse von mindestens ca. 25 Mikrometer aufweisen.
  11. Verbundmaterial nach einem der Ansprüche 8 bis 10, wobei die Oberflächenschicht mindestens ein Metallkarbid enthält, dass in situ durch eine Reaktion von Graphit mit der Legierung gebildet wurde.
  12. Verbundmaterial nach einem der Ansprüche 8 bis 11, wobei die Oberflächenschicht aus Zirkoniumkarbid besteht.
  13. Verbundmaterial nach einem der vorhergehenden Ansprüche, welches ausserdem eine dritte Phase aufweist, die in der Matrix eingebettet ist, wobei die dritte Phase kristalline Teilchen enthält.
  14. Verbundmaterial nach Anspruch 13, wobei die dritte Phase kristalline Teilchen enthält, die aus denselben Elementen wie die Legierung der ersten Phase zusammengesetzt sind.
  15. Verbundmaterial nach Anspruch 13, wobei die dritte Phase Karbidteilchen enthält.
  16. Verbundmaterial nach Anspruch 15, wobei die Karbidteilchen mindestens ein Metallkarbid enthalten, dass in situ durch eine Reaktion von Graphit mit der Legierung gebildet wurde.
  17. Verbundmaterial nach Anspruch 15 oder 16, wobei die Karbidteilchen aus Zirkoniumkarbid bestehen.
  18. Verbundmaterial nach einem der Ansprüche 15 bis 17, wobei die Karbidteilchen eine Grösse von kleiner oder gleich 10 Mikrometer aufweisen.
  19. Verwendung eines Verbundmaterials nach einem der Ansprüche 1 bis 18 zur Herstellung eines Objektes zur Verwendung in einer Vorrichtung ausgewählt aus: ein Reiblager, ein Gelenk, eine Feder.
  20. Verfahren zur Herstellung eines Verbundmaterials, aufweisend:
    - Erhitzen einer Legierung über ihre Liquidus-Temperatur, um eine flüssige Legierung zu bilden;
    - Dispergieren von Graphitpulver in der flüssigen Legierung, um eine fein dispergierte Mischung zu bilden;
    - Abkühlen der Mischung unter ihre Glasübergangstemperatur, wobei dies genügend schnell erfolgt, um ein Verbundmaterial zu bilden, welches eine amorphe erste Phase, die eine kontinuierliche Legierungsmatrix bildet, und eine zweite Phase, die in der Matrix eingebettet ist und Graphitteilchen mit einer Grösse im Bereich zwischen 1 und 100 Mikrometer enthält, aufweist, wobei die zweite Phase zwischen 3 Volumenprozent und 20 Volumenprozent des Verbundmaterials einnimmt.
  21. Verfahren nach Anspruch 20, wobei die Legierung durch Induktionsschmelzen über dem Graphitpulver über ihre Liquidus-Temperatur erhitzt wird.
  22. Verfahren nach Anspruch 20 oder 21, wobei die Mischung mindestens einmal umgeschmolzen wird, und zwar für eine genügend lange Zeit, damit sich eine wahrnehmbare Karbidschicht auf der Oberfläche der Graphitteilchen bildet.
  23. Verfahren nach einem der Ansprüche 20 bis 22, wobei die Mischung mindestens einmal umgeschmolzen wird, und zwar für eine genügend lange Zeit, damit ein Teil der Graphitteilchen mit mindestens einer Metallkomponente der Legierung reagiert, um Metallkarbid-Teilchen zu bilden.
EP06775159A 2005-10-03 2006-08-29 Verbundwerkstoffe aus metallischem massivglas und graphit Not-in-force EP1957686B1 (de)

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Families Citing this family (45)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CH698962B1 (fr) 2008-06-10 2014-10-31 Rolex Sa Ressort de barillet et procédé pour sa mise en forme.
JP5273012B2 (ja) * 2009-10-26 2013-08-28 Ｔｄｋ株式会社 希土類磁石
CN102713770B (zh) * 2009-12-09 2015-11-25 劳力士有限公司 用于制造钟表用弹簧的方法
EP2395125A1 (de) * 2010-06-08 2011-12-14 The Swatch Group Research and Development Ltd. Verfahren zur Herstellung eines Teils aus amorphem beschichteten Metall
EP2585882B1 (de) * 2010-06-22 2021-02-24 The Swatch Group Research and Development Ltd. Stossdämpfungssystem für eine uhr
EP2632809B1 (de) 2010-10-27 2015-11-18 Intercontinental Great Brands LLC Magnetisch verschliessbare verpackung für ein produkt
JP5973790B2 (ja) * 2012-05-28 2016-08-23 株式会社中山アモルファス 耐食性、導電性、成形性に優れた薄板およびその製造方法
WO2014004704A1 (en) 2012-06-26 2014-01-03 California Institute Of Technology Systems and methods for implementing bulk metallic glass-based macroscale gears
US20140010259A1 (en) * 2012-07-04 2014-01-09 Joseph Stevick Temperature tuned failure detection device
WO2014058498A2 (en) 2012-07-17 2014-04-17 California Institute Of Technology Systems and methods for implementing bulk metallic glass-based macroscale compliant gears
US9328813B2 (en) * 2013-02-11 2016-05-03 California Institute Of Technology Systems and methods for implementing bulk metallic glass-based strain wave gears and strain wave gear components
US20140342179A1 (en) 2013-04-12 2014-11-20 California Institute Of Technology Systems and methods for shaping sheet materials that include metallic glass-based materials
US9610650B2 (en) 2013-04-23 2017-04-04 California Institute Of Technology Systems and methods for fabricating structures including metallic glass-based materials using ultrasonic welding
US10081136B2 (en) 2013-07-15 2018-09-25 California Institute Of Technology Systems and methods for additive manufacturing processes that strategically buildup objects
US9868150B2 (en) 2013-09-19 2018-01-16 California Institute Of Technology Systems and methods for fabricating structures including metallic glass-based materials using low pressure casting
US9963395B2 (en) 2013-12-11 2018-05-08 Baker Hughes, A Ge Company, Llc Methods of making carbon composites
CN103938131B (zh) * 2014-03-27 2016-01-20 安徽华兴金属有限责任公司 一种非晶合金组合物及制备方法
CN104087876B (zh) * 2014-06-19 2016-06-08 郑州大学 金属玻璃/石墨烯复合材料的制备方法
US9325012B1 (en) 2014-09-17 2016-04-26 Baker Hughes Incorporated Carbon composites
US10315922B2 (en) 2014-09-29 2019-06-11 Baker Hughes, A Ge Company, Llc Carbon composites and methods of manufacture
US10480288B2 (en) 2014-10-15 2019-11-19 Baker Hughes, A Ge Company, Llc Articles containing carbon composites and methods of manufacture
US9962903B2 (en) * 2014-11-13 2018-05-08 Baker Hughes, A Ge Company, Llc Reinforced composites, methods of manufacture, and articles therefrom
US11097511B2 (en) 2014-11-18 2021-08-24 Baker Hughes, A Ge Company, Llc Methods of forming polymer coatings on metallic substrates
US10300627B2 (en) 2014-11-25 2019-05-28 Baker Hughes, A Ge Company, Llc Method of forming a flexible carbon composite self-lubricating seal
US9726300B2 (en) * 2014-11-25 2017-08-08 Baker Hughes Incorporated Self-lubricating flexible carbon composite seal
US10487934B2 (en) 2014-12-17 2019-11-26 California Institute Of Technology Systems and methods for implementing robust gearbox housings
US10151377B2 (en) 2015-03-05 2018-12-11 California Institute Of Technology Systems and methods for implementing tailored metallic glass-based strain wave gears and strain wave gear components
US10174780B2 (en) 2015-03-11 2019-01-08 California Institute Of Technology Systems and methods for structurally interrelating components using inserts made from metallic glass-based materials
US10155412B2 (en) 2015-03-12 2018-12-18 California Institute Of Technology Systems and methods for implementing flexible members including integrated tools made from metallic glass-based materials
US9840887B2 (en) 2015-05-13 2017-12-12 Baker Hughes Incorporated Wear-resistant and self-lubricant bore receptacle packoff tool
US10968527B2 (en) 2015-11-12 2021-04-06 California Institute Of Technology Method for embedding inserts, fasteners and features into metal core truss panels
US10125274B2 (en) 2016-05-03 2018-11-13 Baker Hughes, A Ge Company, Llc Coatings containing carbon composite fillers and methods of manufacture
US10344559B2 (en) 2016-05-26 2019-07-09 Baker Hughes, A Ge Company, Llc High temperature high pressure seal for downhole chemical injection applications
CN106011696B (zh) * 2016-05-27 2017-08-11 中国科学院金属研究所 一种高强高弹高塑性的块体纳米金属玻璃材料及制备方法
WO2018165662A1 (en) 2017-03-10 2018-09-13 California Institute Of Technology Methods for fabricating strain wave gear flexsplines using metal additive manufacturing
EP3630395A4 (de) 2017-05-24 2020-11-25 California Institute of Technology Hypoeutektische amorphe metallbasierte materialien zur generativen fertigung
US11014162B2 (en) 2017-05-26 2021-05-25 California Institute Of Technology Dendrite-reinforced titanium-based metal matrix composites
EP3630397A4 (de) 2017-06-02 2020-11-11 California Institute of Technology Hochfeste verbundstoffe auf metallglasbasis zur generativen fertigung
US11859705B2 (en) 2019-02-28 2024-01-02 California Institute Of Technology Rounded strain wave gear flexspline utilizing bulk metallic glass-based materials and methods of manufacture thereof
US11680629B2 (en) 2019-02-28 2023-06-20 California Institute Of Technology Low cost wave generators for metal strain wave gears and methods of manufacture thereof
US11400613B2 (en) 2019-03-01 2022-08-02 California Institute Of Technology Self-hammering cutting tool
US11591906B2 (en) 2019-03-07 2023-02-28 California Institute Of Technology Cutting tool with porous regions
CN110129691A (zh) * 2019-06-23 2019-08-16 重庆师范大学 一种大块石墨烯金属玻璃复合材料及其制备方法
CN110981474B (zh) * 2019-12-13 2021-08-31 中钢集团洛阳耐火材料研究院有限公司 一种氧化锆轻质保温耐火制品
CN112553569B (zh) * 2021-01-08 2021-08-17 吉林大学 纳秒激光碳化提高锆基非晶合金表面硬度的方法

Family Cites Families (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US5127969A (en) * 1990-03-22 1992-07-07 University Of Cincinnati Reinforced solder, brazing and welding compositions and methods for preparation thereof
JPH05117823A (ja) * 1991-10-22 1993-05-14 Takeshi Masumoto 繊維強化金属複合材料
US5735975A (en) * 1996-02-21 1998-04-07 California Institute Of Technology Quinary metallic glass alloys
WO2002027055A1 (fr) * 2000-09-25 2002-04-04 Tohoku Techno Arch Co., Ltd. Alliage amorphe et procede de preparation dudit alliage

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