US20060213586A1 - Metal composites and methods for forming same - Google Patents

Metal composites and methods for forming same Download PDF

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US20060213586A1
US20060213586A1 US11/088,106 US8810605A US2006213586A1 US 20060213586 A1 US20060213586 A1 US 20060213586A1 US 8810605 A US8810605 A US 8810605A US 2006213586 A1 US2006213586 A1 US 2006213586A1
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alloy
metal
metal composite
combinations
spinodal
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Hin-Wing Kui
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Priority to US11/088,106 priority Critical patent/US20060213586A1/en
Priority to CN200680017144.1A priority patent/CN101203622A/zh
Priority to PCT/CN2006/000483 priority patent/WO2006099808A1/fr
Priority to EP06722135A priority patent/EP1861515A4/fr
Priority to JP2008502227A priority patent/JP2008537763A/ja
Priority to TW095110087A priority patent/TW200641151A/zh
Publication of US20060213586A1 publication Critical patent/US20060213586A1/en
Priority to US13/472,801 priority patent/US20130098510A1/en
Abandoned legal-status Critical Current

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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
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • 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
    • C22C1/02—Making non-ferrous alloys by melting
    • 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
    • C22C1/06—Making non-ferrous alloys with the use of special agents for refining or deoxidising
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working

Definitions

  • This invention relates generally to a metal composite and methods for forming same and, more particularly, to a metal composite having a network structure with at least one ductile sub-network.
  • Nanostructured materials are defined as materials having a grain size diameter of between and including I nm and 1,000 nm. The presence of nanostrutures in a material improves mechanical properties when compared to the same material formed without nanostructures.
  • Nanostructured materials have typically been synthesized by powder sintering or thermal annealing of glassy metals, and fluxing. In powder sintering, nanometer size powders are compacted together, and may be annealed, to produce nanostructured materials. Nanostructured materials produced by powder sintering are typically disk shaped, having a diameter of about 1 cm, and a thickness of 1 mm to 2 mm. The powdered sintered materials are typically brittle and exhibit voids and non-uniform grain growth.
  • thermal annealing of an amorphous metal produces nanocrystals in an amorphous matrix.
  • a metallic melt is quenched into a glass or amorphous metal, which is then annealed at a temperature near its glass transition temperature, resulting in uniformly distributed nanocrystals in an amorphous matrix.
  • the amorphous matrix also crystallizes, a crystalline nanostructure may be produced.
  • a fluxing technique has more recently been used to prepare nanostructured materials.
  • liquid-state spinodal decomposition occurs resulting in brittle solid spinodals.
  • a typical Pd-Si nanostructure is shown in FIG. 1 .
  • discrete Pd precipitates shown in white
  • the discrete Pd precipitates are at a negligible volume fraction so that the solid spinodals remain brittle.
  • conventional nanostructures are brittle, and have a broad grain size distribution.
  • Conventional nanostructures also have low strength and low impact fracture energy.
  • conventional amorphous metals typically have a small overall dimension, such as ribbon or foil forms having a thickness of 30-50 microns that renders them inappropriate for commercial applications.
  • the present invention is directed to a method comprising forming an alloy with the constituents, purifying the alloy, and forming a network structure of the alloy comprising at least one ductile sub-network structure.
  • Another embodiment is directed to a metal composite comprising a ductile spinodal structure. Another embodiment is directed to a metal article comprising the metal composite. In another embodiment, the metal article is a nanostructure composite.
  • Another embodiment is directed to a method of forming a metal composite comprising forming an alloy, purifying the alloy, forming one or more spinodals, and heating the one or more spinodals causing at least one of the one or more spinodals to transform into one or more ductile phases.
  • FIG. 1 is a TEM micrograph of a conventional Pd-Si nanostructure at a magnification of 75,000X.
  • FIG. 2 is a SEM micrograph showing the microstructure of a conventional steel ball at a magnification of 9,500X.
  • FIG. 3A is a SEM micrograph showing a portion of a two phase spinodal microstructure of Fe 80 Cl 5 Si 5 at a magnification of 35, 000 X.
  • FIG. 3B is a SEM micrograph showing another portion of the two phase spinodal microstructure of FIG. 3A at a magnification of 6,00OX.
  • FIG. 3C is a SEM micrograph showing another portion of the two phase spinodal microstructure of FIG. 3A at a magnification of 34,000X.
  • FIG. 3D is a SEM micrograph showing another fracture surface of the two phase spinodal microstructure of FIG. 3A at a magnification of 8,500X.
  • FIG. 3E is a SEM micrograph showing a fracture surface of the two phase spinodal microstructure of FIG. 3A at a magnification of 9,400X.
  • FIG. 4A is a SEM micrograph showing a portion of a spinodal structure of Fe 405 Co 405 Cl 4 Si 5 at magnification of 34, 000 X.
  • FIG. 4B is a SEM micrograph showing another portion of the spinodal structure of FIG.4A at a magnification of 5,400X.
  • FIG. 4C is a SEM micrograph showing a fracture surface of the spinodal structure of FIG. 4A at a magnification of 13,500X.
  • FIG. 4D is a SEM micrograph showing another fracture surface of the spinodal structure of FIG. 4A at a magnification of 4,1000X.
  • FIG. 5A is a SEM micrograph showing a portion of a spinodal structure of Co 75 Si 15 B, 0 at a magnification of 34,000X.
  • FIG. 5B is a SEM micrograph showing another portion of the spinodal structure of FIG. 5A at a magnification of 2,000X.
  • FIG. 5C is a SEM micrograph showing another portion of the spinodal structure of FIG. 5A at a magnification of 5,400X.
  • FIG. 5D is a SEM micrograph showing a fracture surface of the spinodal structure of FIG. 5A at a magnification of 1,400X.
  • FIG. 6A is a SEM micrograph showing a portion of a spinodal structure of Fe 82 Cl 8 at a magnification of 4,100X.
  • FIG. 6B is a SEM micrograph of a fracture surface of the spinodal structure of FIG. 6A at a magnification of 9,500X.
  • the present invention provides a metal composite comprising a network of at least two spinodals, or sub-networks, and methods for forming same.
  • spinodal and “sub-network” are used interchangeably herein, to define a solid morphology of isolated clusters and/or interconnected regions following liquid state spinodal decomposition and subsequent solidification.
  • spinodal decomposition and subsequent solidification of a binary alloy results in two sub-networks.
  • spinodal decomposition of a tertiary alloy results in two or three sub-networks.
  • Sub-networks may include discrete precipitates, but not in an appreciable amount.
  • At least one spinodal or sub-network of the metal composite is a ductile phase.
  • ductile phase is defined as a malleable phase.
  • the ductile spinodal may be isolated clusters, partially interconnected, substantially interconnected, and combinations thereof. That is to say that the degree of interconnection of the ductile spinodal may vary throughout the metal composite.
  • the metal composite comprises a ductile spinodal and a brittle spinodal.
  • the metal composite comprises a first ductile spinodal and a second ductile spinodal.
  • relative volume fractions of the ductile and brittle phases may be changed to vary any desired property, such as but not limited to, hardness, fatigue strength, compressive yield strength, wear resistance, and maximum operating temperature.
  • the metal composite may be formed from any alloy for which a predominant constituent has a positive heat of mixing with at least one other constituent in the liquid state.
  • heat of mixing is defined as an enthalpy change when I mol of a mixture is formed from its pure components at a temperature, T.
  • the term “predominant” is defined as an intended primary constituent of the alloy.
  • the liquid state may be either stable or metastable. Examples of alloys having constituents with a positive heat of mixing of include: monotectic, eutectic, and peritectic alloys. Other examples of alloys having constituents with a positive heat of mixing are noted in the text “Thermodynamics of Solids,” by Richard A. Swalin, Published by John Wiley & Sons, Inc.(1962), which is incorporated herein by reference for all purposed.
  • the alloy may comprise a metal and a metalloid.
  • the overall heat of mixing of all constituents in the alloy may be positive.
  • the overall heat of mixing in the liquid state of all constituents in the alloy may be negative.
  • an alloy comprised of first, second, and third components may have an overall negative heat of mixing in the liquid state, although the first constituent (the major constituent) has a positive heat of mixing in the liquid state with the second component, and also has a positive heat of mixing in the liquid state with the third component.
  • an overall negative heat of mixing may arise from a very negative heat of mixing in the liquid state for the second and third constituents, which overwhelms the positive heats of mixing in the liquid state for the first major constituent with the second and third constituents.
  • the metal may be a Group VIII metal, such as Fe, Co, Ni, Ru, Rh, Pd, Os, Ir, Pt, and combinations thereof.
  • the metal may be selected from the group consisting of Fe, Co, Ni, Cu, Pd, Pt, Mn, Al, T, Zr, Cr, W, and combinations thereof.
  • the metal is Co.
  • the metal is Fe.
  • the metal is Ni.
  • any metalloid may be used to obtain desired composite properties, as long as the selected metalloid has a positive heat of mixing when combined with other predominant constituents in the liquid state.
  • the metalloid may be any of B, C, Si, As, Sb, Te, Po and combinations thereof.
  • the metalloid may be any of B, C, Si, and combinations thereof.
  • the metalloid is C.
  • the alloy comprises Fe and at least one of Si and C.
  • Fe may range between and including about 70 atomic % and about 92 atomic %
  • Si may range between and including about 0 atomic % and about 20 atomic %
  • C may range between and including about 0 atomic % and about 30 atomic %.
  • the ranges used represent minimums and maximums of individual components, though it is understood that individual components are combined in such a way to result in an atomic percent of the alloy of 100 %. It is believed that these ranges of Fe, Si, and C will result in a ductile spinodal network when processed according to the methods disclosed herein. Although compositions outside these ranges may produce the desired ductile nanostructure, it is believed these ranges may be more effective under processing conditions.
  • the Fe—C—Si composition may range between and including Fe 76 C 24 Si 0 ; Fe 81 C 19 Si 0 ; Fe 85.5 C 0 Si14.5; Fe 88.5 C 0 Si 11.5 and all compositions in between. It has been found that this range, and all points therein, result in a nanostructure comprising ductile spinodals upon air cooling, and substantial undercooling.
  • the Fe—C—Si composition may range between and including Fe 81 C 19 Si 0 ; Fe 84 C 16 Si 0 ; Fe 90 C 0 Si 10 ; Fe 88.5 C 0 Si 11.5 and all points in between.
  • the Fe—C—Si composition of another embodiment ranges between and including Fe 73 C 27 Si 0 ; Fe76C24Si 0 ; Fe84C 0 Si16; Fe 85.5 C 0 S 14.5 and all points therein. It has been found that within these ranges, and all points there between, portions of the composite formed from these compositions form nanostructures comprising ductile spinodals upon air cooling. However, increased cooling rates may be used so that the entire sample of these compositions may form the desired nanostructures.
  • the Fe—C—Si composition may range between and including Fe 84 Cl 6 Si 0 ; Fe 87 Cl 3 SiO; FegoCoSilo; Fe 92 COSi 8 and all points in between.
  • the Fe—C—Si composition of another embodiment ranges between and including Fe 70 C 30 Si 0 ; Fe 73 C 27 Sio; Fe 82 C 0 Si 18 ; Fe 84 C 0 Si 16 and all points therein. It is believed that with these ranges, the desired nanostructures may form in a drop tower filled with a gaseous medium.
  • B may be added to the Fe—C—Si alloy.
  • B may be added in a range of between and including about 0 atomic % and about 5 atomic % without significantly impacting the formation of the spinodal structures.
  • the alloy may comprise any non-metal selected from the group Ge, P, S and combinations thereof.
  • Ge may replace, or be used in addition to Si, without significantly affecting the nanostructure, However, in some instances, the presence of Ge resulted in the formation of voids.
  • P may be added to the alloy to increase the formation of spinodal structures. For example, between and including about 0.5 atomic % and about 4 atomic % of P may be added to the Fe—C—Si alloy so that the entire composite may comprise a spinodal structure. It has been found that the addition of P in the Fe—C—Si alloy reduces or eliminates the presence of eutectic structures, thereby increasing the amount of spinodal structures.
  • Ni may be added to the Fe—C—Si ally to increase the volume fraction of the ductile phase.
  • about 1 atomic % to more that about 3 atomic % Ni may be added to the F 3 —C—Si alloy to increase the volume fraction of the ductile phase.
  • the alloy may have a glass forming ability (GFA) defined herein as a ratio of a glass transition temperature (T g ) to a liquidus temperature (T 1 ), greater than or equal to about 0.35 so that the nanostructures formed are sufficiently large to impart the desired physical properties to the composite.
  • GFA glass forming ability
  • the GFA is greater than or equal to about 0.49.
  • Fe 825 B 17.5 has a GFA of about 0.35
  • Fe 80 B 20 has a GFA of about 0.49
  • Fe 80 C 7 P 13 has a GFA of about 0.58
  • Fe 79 Si 10 B 11 has a GFA of about 0.58
  • Co 75 Si 15 B 10 has a GFA of about 0.56.
  • the alloy may be formed by heating a selected composition of constituents in a desired proportion. Heating may be carried out under normal alloying conditions with conventional equipment, such as a radio frequency induction furnace or high temperature furnace.
  • the formed alloy may be divided into smaller portions of the alloy for further processing.
  • the alloy may be placed in a first portion of a container which also has a second portion smaller than the first portion.
  • a vacuum may drawn on the container.
  • the container may be heated allowing the alloy to melt forming a first molten alloy.
  • the first molten alloy may be forced into the second portion of the container with a pressurized gas.
  • the container and the first molten alloy may be cooled forming a solidified alloy.
  • the solidified alloy may be removed from the second portion of the container and may be apportioned into a desired size or mass.
  • the first portion of the container may have a cross sectional area smaller than a cross sectional area of the second portion. Accordingly, the second portion may be sufficiently long to accommodate the entirety of the molten alloy in the narrower cross sectional area.
  • the solidified alloy may be cut into various thicknesses, suitable for a desired application.
  • the alloy not further divided or the solidified alloy further divided into smaller portions may be further processed to remove impurities according to conventional methods.
  • the solidified alloy may be heated to a temperature greater than or equal to its liquidus (T 1 ) in the presence of a flux to form a second molten alloy.
  • the solidified alloy and flux material may be heated to a temperature greater than about 1,000° C.
  • the flux may be boron oxide, glass, calcium oxide, barium oxide, aluminum oxide, magnesium oxide, lithium oxide, and mixtures thereof.
  • the flux is glass.
  • the flux is boron oxide.
  • Boron oxide (B 2 0 3 ), nominally anhydrous, is available from Atomergic Chemetals Corporation (Farmingdale, N.Y.).
  • the second molten alloy may then be cooled to a temperature sufficient to form a second solidified alloy.
  • the second molten alloy may be undercooled by cooling the second molten alloy below its liquidus.
  • the second molten alloy may be cooled to or below a critical temperature, Tc, typically below the liquidus to allow liquid state spinodal decomposition, thereby forming liquid spinodals.
  • Tc critical temperature
  • the second molten alloy is a metastable liquid that undergoes spinodal decomposition upon entering the metastable miscibility gap (critical temperature T C ), which is often substantially below the liquidus of the alloy.
  • the molten alloy splits into a number of metastable liquid spinodals having a liquid phase wavelength ⁇ .
  • the wavelength ( ⁇ ) is defined as a lateral dimension, or diameter, of the spinodal.
  • the metastable liquid may have a liquid phase ⁇ of less than about 300 nm, preferably less than about 100 nm.
  • the second molten alloy solidifies to form an undercooled specimen having a spinodal or sub-network structure.
  • the solid spinodals may be crystalline, amorphous, quasi-crystalline, and mixtures thereof.
  • the alloy is cooled to a ⁇ T of about 100° K to about 500° K.
  • AT is defined as the difference between the liquidus temperature and actual temperature (T 1 -T).
  • the solid phase wavelengths ( ⁇ ) formed by the liquid state spinodal decomposition of the second molten alloy typically range from microns to nanometers.
  • the resultant composite may have a fine microstructure defined as a material having a grain size diameter, or wavelength, of between and including 1 nm and 100,000 nm.
  • the composite may include nanostructures, wherein one physical dimension of one constituent phase is about 1,000 nm or less.
  • Each spinodal, or sub-network, within the entire networked structure my have solid phase wavelengths that are similar to or differing from one another.
  • the spinodal or sub-network structure formed may have a solid phase ⁇ of less than about 50 microns. In another embodiment, the spinodal or sub-network structure formed may have a solid phase k of about 10 microns or less. In yet another embodiment, the spinodal structure formed has a solid phase k of about 300 nm or less, preferably less than about 100 nm.
  • the solid phase wavelength may vary throughout a specimen.
  • the solid phase wavelength may change during crystallization.
  • a crystallization front moves across the molten so that the solid phase ⁇ increases, effectively creating a coarser spinodal structure. Therefore, the solid phase wavelength may be smallest at a site where crystallization is initiated, and increase as crystallization proceeds further from the initiation site.
  • the wavelength of the solid spinodal at the site where crystallization is initiated may be similar to the wavelength of the liquid spinodal.
  • the spinodal morphology at a distance from the site where crystallization is initiated may be replaced (partially or entirely) by other structures, including dendrite and eutectic.
  • the liquid spinodals may solidify into amorphous spinodals upon further cooling.
  • the solidification may occur homogeneously, that is, hardening does not begin at any single location.
  • the hardening continues on further cooling, until all the liquid spinodals become amorphous solids.
  • the solid phase k is substantially uniform in this mode of solidification, so that the liquid phase ⁇ 60 of the liquid spinodals may be similar to the solid phase ⁇ of the amorphous spinodals.
  • Solid spinodals may also form as a mixture of amorphous and crystalline spinodals, so that ⁇ may vary through out the solid.
  • coherent grain boundary is defined as a coherent interface and/or a semi-coherent interface.
  • a coherent interface has interfacial energies of about 10 to about 100 mJ/m 2 , and occurs when two crystals perfectly match at an interface plane so that the two lattices are continuous across the interface.
  • a semi coherent interface occurs when the interface has a series of edge or screw dislocations and has an interfacial energy of about 200 mJ/m 2 to about 500 mJ/m 2 .
  • any brittle spinodals present in a specimen may be further treated to undergo phase transformation into one or more ductile spinodals.
  • Fe 3 Si decomposes into Fe and graphite, thereby further increasing the strength and impact fracture of the specimen.
  • one or more brittle spinodals may be annealed to form ductile phases.
  • the resultant ductile phases may be isolated clusters, partially or substantially completely interconnected, and combinations thereof.
  • the ductile phases may be interconnected with other ductile phases, and or with brittle phases.
  • the metal composite may be a bulk material having any shape suitable for a particular purpose.
  • the phrase “bulk material” is defined as a material having a shape with a cross sectional dimension greater than or equal to about 1 mm in all directions.
  • the composite may be a sphere, cone, pyramid, square, rectangle or irregular shape.
  • the composite is a sphere.
  • the bulk material has at least one cross sectional dimension of about 2.54 cm, preferably about 1 cm.
  • the metal composite may be a sphere having any diameter suitable for a particular purpose.
  • the metal composite may be a sphere having a diameter less than about 1 inch, less than or equal to about 2 cm, less than or equal to about 1.0 cm, and less than or equal to about 5 mm respectively.
  • the spherical metal composite may have a diameter of about 0.1 mm.
  • a desired composition of raw materials selected from Fe, Co, Ni, C, Si, B, Ge, and P, were alloyed in an RF induction furnace at a minimum temperature of about 1,000° C.
  • the alloy was air cooled to solidify, and then positioned in a large portion of a fused silica tube.
  • the fused silica tube comprised a large portion with an inner diameter of about 10 mm to 30 mm, and a long small portion having an inner diameter of about 2 mm to about 8 mm. The small portion had a length of about 10 mm to about 600 mm.
  • the silica tube containing the alloy was evacuated by a mechanical pump and placed in a furnace at a sufficient temperature and for a sufficient time to melt the alloy.
  • Each disk was positioned with anhydrous B 2 O 3 in individual fused silica tubes, having inner diameters between about 3 mm and about 15 mm, and a length of about 10 mm to about 100 mm.
  • a number of fused silica tubes containing an alloy disk and anhydrous B 2 O 3 were placed in a larger fused silica tube having a diameter of about 20 mm to about 100 mm.
  • a vacuum was drawn on the larger silica tube, which created a vacuum in the individual tubes containing each alloy disk and anhydrous B 2 O 3 .
  • the large tube was then heated at sufficient temperature and for a specified time period, ranging from about 15 minutes to about 8 hours to completely melt the alloys.
  • the molten alloys were cooled and allowed to crystallize at the AT temperatures noted below.
  • This alloy was prepared by purifying molten Fe 80 C 15 Si 5 in a flux above its liquidus, and subsequent undercooling below its liquidus.
  • the Fe—C—Si system was formed into a precision ball bearing with some of its properties listed in Table 1. Also listed in Table are comparative results for conventional chrome steel balls (available from FAG Bearing, Danbury, Conn.). TABLE 1 Chrome steel balls (available from FAG Property Fe—C—Si system Bearing, Danbury, CT) Hardness (HV) 750-850 Fatigue strength >3,600** ⁇ 2,600*** (Maximum compressive pressure (MPA))* Compressive yield About 7,000 About 3,600 strength (MPA) Maximum operating 550° C. 150° C. temperature *Fatigue strength was determined by using a cyclic compressive force, with a minimum compressive pressure of bout 0 MPA and a maximum cyclic compressive pressure of about 3,600 MPA for 10 7 cycles. **Samples remained in tact after completion of the test. ***Samples did not survive the test.
  • the Fe—C—Si system had about the same wear resistance and hardness as the conventional steel ball.
  • balls made from the Fe—C—Si system exhibited a compressive yield strength almost double that of the conventional steel ball (7,000 MPA compared to 3,600 MPA).
  • the balls made from the Fe—C—Si system also exhibited a greater fatigue strength of 3,600 MPA without failure, compared to the conventional balls which failed at 2,600 MPA.
  • the balls of the Fe—C—Si system exhibited a greater Young's modulus and greater thermal stability than the conventional steel balls.
  • the Fe—C—Si system balls exhibited an impact fracture energy that approached that of a ceramic precision ball bearing made of SiN.
  • FIG. 3A is a SEM of a two phase spinodal microstructure of the Fe 80 Cl 5 Si 5 ball.
  • the Fe 80 Cl 5 Si 5 was analyzed at various locations within the ball.
  • the micrograph in FIG. 3A shows a two-phase spinodal structure having an interconnected microstructure. Both phases are crystalline, and have an average wavelength of about 300 nm. The randomness of the microstructure indicates that this is the site where crystallization was initiated.
  • FIG. 3B is a micrograph of a center of the specimen of FIG. 3A .
  • FIG. 3B shows alignments and differing orientation of the alignments of the spinodal structure near the center of the specimen.
  • FIG. 3C is a micrograph of the specimen at an end opposite to the site of initiation.
  • FIG. 3C shows the formation of eutectic structures at the end opposite to the site of initiation.
  • FIG. 3D shows the fracture behavior of this system, in which the fracture surface is scale-like or cloud-like.
  • the white curves in FIG. 3D are ridges, illustrating that ductile fracture has occurred.
  • FIG. 3E is another micrograph of the specimen of FIG. 3A showing the microstructure on a fracture surface of the specimen.
  • the two solid spinodals in the metal composite are Fe 3 Si and body centered cubic (BCC) Fe, or solid solution of Fe.
  • BCC body centered cubic
  • the melt split into two liquid spinodals. These two liquid spinodal were metastable and were, therefore, prone to crystallization.
  • Crystallization started at a point on the surface of the molten specimen (called the site of initial crystallization). The crystallization front then spread out until the entire molten specimen became a crystal.
  • heat was released (heat of crystallization) that may have been partly responsible for the change in microstructures at increasing distances from the site of initial crystallization.
  • the morphology of the liquid spinodals may have been similar to the morphology of the solid spinodals. Moving away from the site of initial crystallization, the spinodal morphology evolved.
  • FIG. 2 is a SEM of a microstructure of a conventional steel ball, typically a mixture of austenite and martensite, produced by conventional methods.
  • steel in the form of a wire coil still austenite, which is soft
  • Short cylindrical pieces are cut form the wire and cold forged in a heading machine.
  • the surface of the headed balls are ground by a flashing machine.
  • the balls are then hardened in a furnace, transforming more than half of the austentite into martensite, which is hard. After hardening, the balls go though two more grinding processes (grinding and lapping) to produce a desire surface finish.
  • the balls are then cleaned and polished.
  • Drawbacks to the conventional method of preparing steel balls include the necessity of using a high quality wire coil, and ensuring the heat treatment does not convert all the austenite into martensite. Moreover, because not all the austenite is converted, the conventional steel balls may not be suitable for high temperature application, such as those greater than 150° C.
  • the microstructure including the ductile phase of the present invention provides a micro-or nano-structured object with unique physical properties when compared to similar objects made with conventional methods, such as: greater compressive yield strength, greater Young's modulus, greater fatigue resistance, and greater thermal stability, while maintaining similar wear resistance and hardness, while also exhibiting similar impact fracture energy similar to that of ceramic balls.
  • FIGS. 4 A-4D The microstructures of an as formed Fe 40.5 CO 40.5 C 14 Si 5 ingot (the specimen) are shown in FIGS. 4 A-4D.
  • FIG. 4A shows the microstructure of a portion of the specimen that crystallized first, locating the free surface of the specimen. In FIG. 4A , there are two solid spinodals, of different phases, which together form a random network.
  • FIG. 4B shows the microstructure of a portion of the specimen that crystallized last, at the opposite end of the specimen. As seen in FIG. 4 b , the microstructure is a mixture of spinodal morphology and eutectic structure.
  • FIG. 4A shows the microstructure of a portion of the specimen that crystallized first, locating the free surface of the specimen. In FIG. 4A , there are two solid spinodals, of different phases, which together form a random network.
  • FIG. 4B shows the microstructure of a portion of the specimen that crystallized last, at the opposite end of
  • 4 c shows the fracture surface of the specimen, in which bright ridges indicate ductile fracture had take place in one of the two spinodals.
  • the distribution of the ridges is consistent with the prediction of a spinodal mechanism.
  • the eutectic morphology is exhibited on the fracture surface shown in FIG. 4D .
  • FIGS. 5A-5D The microstructures of an as formed CO 75 Si 15 B 10 ingot (specimen) are shown in FIGS. 5A-5D .
  • FIG. 5A shows the microstructures of a portion of the specimen which crystallized first, located at the free surface of the ingot.
  • the network morphology changes as the location in the specimen moves away from the site where crystallization was initiated.
  • Toward the center of the specimen, one solid spinodal shown apparently breaks into elongated grains, as shown in FIG. 5B .
  • the elongated grains may be over about 20 microns.
  • a layer of multiple phases separates the elongated grains.
  • FIG. 5C shows microstructures at a location further from the site of initial crystallization. As shown in FIG. 5C , the layer surrounding the elongated grains of the portion of the specimen that crystallized last appear to be substantially homogeneous.
  • FIG. 5D shows a fracture surface of the specimen which illustrates that the layers surrounding the elongated grains are ductile. Without being bound to any particular theory, it is believed that the ductile layers provide the specimen with impact fracture energy and strength.
  • FIG. 6A The microstructure of an as prepared ingot of Fe 82 C 18 specimen is shown in FIG. 6A .
  • FIG. 6A there are many domains, each of which is occupied by a network-like structure that is aligned.
  • the boundaries between the domains are relatively flat, and the alignments of the network close to the boundaries are sharp.
  • FIG. 6B shows the microstructure of a fracture surface of the specimen of FIG. 6A .
  • two vertical lines are found near the center of the micrograph, and represent the boundaries described above. Attached to these boundaries are aligned features, similar to that of a dendrite. Further away from the boundaries, scale-like or cloud-like structures again dominate the microstructure.
  • the lighted parts of FIG. 6B represent ridges which have undergone a ductile fracture.

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CN200680017144.1A CN101203622A (zh) 2005-03-23 2006-03-23 金属复合材料及其形成方法
PCT/CN2006/000483 WO2006099808A1 (fr) 2005-03-23 2006-03-23 Composites métalliques et méthodes de fabrication desdits composites
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WO2010048060A1 (fr) * 2008-10-21 2010-04-29 The Nanosteel Company, Inc. Mécanisme de formation structurelle pour des composites de verre métallique présentant une ductilité
US20110186259A1 (en) * 2010-02-02 2011-08-04 The Nanosteel Company, Inc. Utilization of Carbon Dioxide And/Or Carbon Monoxide Gases in Processing Metallic Glass Compositions
WO2012061282A3 (fr) * 2010-11-02 2013-11-07 The Nanosteel Company, Inc. Nanomatériaux vitreux
US20150096652A9 (en) * 2013-01-07 2015-04-09 Glassimetal Technology, Inc. Bulk nickel-silicon-boron glasses bearing iron
WO2014012113A3 (fr) * 2012-07-13 2015-05-07 The Nanosteel Company, Inc. Stratifié de fibre métallique vitreuse
US9556504B2 (en) 2012-11-15 2017-01-31 Glassimetal Technology, Inc. Bulk nickel-phosphorus-boron glasses bearing chromium and tantalum
EP2393950A4 (fr) * 2009-02-03 2017-09-13 The Nanosteel Company, Inc. Procédé et produit de découpe de matériaux
US9816166B2 (en) 2013-02-26 2017-11-14 Glassimetal Technology, Inc. Bulk nickel-phosphorus-boron glasses bearing manganese
US9863024B2 (en) 2012-10-30 2018-01-09 Glassimetal Technology, Inc. Bulk nickel-based chromium and phosphorus bearing metallic glasses with high toughness
US9863025B2 (en) 2013-08-16 2018-01-09 Glassimetal Technology, Inc. Bulk nickel-phosphorus-boron glasses bearing manganese, niobium and tantalum
US9920410B2 (en) 2011-08-22 2018-03-20 California Institute Of Technology Bulk nickel-based chromium and phosphorous bearing metallic glasses
US9920400B2 (en) 2013-12-09 2018-03-20 Glassimetal Technology, Inc. Bulk nickel-based glasses bearing chromium, niobium, phosphorus and silicon
US9957596B2 (en) 2013-12-23 2018-05-01 Glassimetal Technology, Inc. Bulk nickel-iron-based, nickel-cobalt-based and nickel-copper based glasses bearing chromium, niobium, phosphorus and boron
US10000834B2 (en) 2014-02-25 2018-06-19 Glassimetal Technology, Inc. Bulk nickel-chromium-phosphorus glasses bearing niobium and boron exhibiting high strength and/or high thermal stability of the supercooled liquid
US10287663B2 (en) 2014-08-12 2019-05-14 Glassimetal Technology, Inc. Bulk nickel-phosphorus-silicon glasses bearing manganese
US10458008B2 (en) 2017-04-27 2019-10-29 Glassimetal Technology, Inc. Zirconium-cobalt-nickel-aluminum glasses with high glass forming ability and high reflectivity
US11371108B2 (en) 2019-02-14 2022-06-28 Glassimetal Technology, Inc. Tough iron-based glasses with high glass forming ability and high thermal stability
US11377720B2 (en) 2012-09-17 2022-07-05 Glassimetal Technology Inc. Bulk nickel-silicon-boron glasses bearing chromium
US11905582B2 (en) 2017-03-09 2024-02-20 Glassimetal Technology, Inc. Bulk nickel-niobium-phosphorus-boron glasses bearing low fractions of chromium and exhibiting high toughness

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US20100154942A1 (en) * 2008-10-21 2010-06-24 The Nanosteel Company, Inc. Mechanism of Structural Formation For Metallic Glass Based Composites with Enhanced Ductility
US8882941B2 (en) 2008-10-21 2014-11-11 The Nanosteel Company, Inc. Mechanism of structural formation for metallic glass based composites with enhanced ductility
WO2010048060A1 (fr) * 2008-10-21 2010-04-29 The Nanosteel Company, Inc. Mécanisme de formation structurelle pour des composites de verre métallique présentant une ductilité
EP2393950A4 (fr) * 2009-02-03 2017-09-13 The Nanosteel Company, Inc. Procédé et produit de découpe de matériaux
US20110186259A1 (en) * 2010-02-02 2011-08-04 The Nanosteel Company, Inc. Utilization of Carbon Dioxide And/Or Carbon Monoxide Gases in Processing Metallic Glass Compositions
WO2011097239A1 (fr) * 2010-02-02 2011-08-11 The Nanosteel Company, Inc. Utilisation des gaz dioxyde de carbone et/ou monoxyde de carbone dans le traitement des compositions de verre métallique
CN102803168A (zh) * 2010-02-02 2012-11-28 纳米钢公司 加工金属玻璃组合物中二氧化碳和/或一氧化碳气体的利用
US8807197B2 (en) 2010-02-02 2014-08-19 The Nanosteel Company, Inc. Utilization of carbon dioxide and/or carbon monoxide gases in processing metallic glass compositions
WO2012061282A3 (fr) * 2010-11-02 2013-11-07 The Nanosteel Company, Inc. Nanomatériaux vitreux
CN103492591A (zh) * 2010-11-02 2014-01-01 纳米钢公司 玻璃状纳米材料
US8986472B2 (en) 2010-11-02 2015-03-24 The Nanosteel Company, Inc. Glassy nano-materials
US9920410B2 (en) 2011-08-22 2018-03-20 California Institute Of Technology Bulk nickel-based chromium and phosphorous bearing metallic glasses
WO2014012113A3 (fr) * 2012-07-13 2015-05-07 The Nanosteel Company, Inc. Stratifié de fibre métallique vitreuse
US11377720B2 (en) 2012-09-17 2022-07-05 Glassimetal Technology Inc. Bulk nickel-silicon-boron glasses bearing chromium
US9863024B2 (en) 2012-10-30 2018-01-09 Glassimetal Technology, Inc. Bulk nickel-based chromium and phosphorus bearing metallic glasses with high toughness
US9556504B2 (en) 2012-11-15 2017-01-31 Glassimetal Technology, Inc. Bulk nickel-phosphorus-boron glasses bearing chromium and tantalum
US20150096652A9 (en) * 2013-01-07 2015-04-09 Glassimetal Technology, Inc. Bulk nickel-silicon-boron glasses bearing iron
US9534283B2 (en) * 2013-01-07 2017-01-03 Glassimental Technology, Inc. Bulk nickel—silicon—boron glasses bearing iron
US9816166B2 (en) 2013-02-26 2017-11-14 Glassimetal Technology, Inc. Bulk nickel-phosphorus-boron glasses bearing manganese
US9863025B2 (en) 2013-08-16 2018-01-09 Glassimetal Technology, Inc. Bulk nickel-phosphorus-boron glasses bearing manganese, niobium and tantalum
US9920400B2 (en) 2013-12-09 2018-03-20 Glassimetal Technology, Inc. Bulk nickel-based glasses bearing chromium, niobium, phosphorus and silicon
US9957596B2 (en) 2013-12-23 2018-05-01 Glassimetal Technology, Inc. Bulk nickel-iron-based, nickel-cobalt-based and nickel-copper based glasses bearing chromium, niobium, phosphorus and boron
US10000834B2 (en) 2014-02-25 2018-06-19 Glassimetal Technology, Inc. Bulk nickel-chromium-phosphorus glasses bearing niobium and boron exhibiting high strength and/or high thermal stability of the supercooled liquid
US10287663B2 (en) 2014-08-12 2019-05-14 Glassimetal Technology, Inc. Bulk nickel-phosphorus-silicon glasses bearing manganese
US11905582B2 (en) 2017-03-09 2024-02-20 Glassimetal Technology, Inc. Bulk nickel-niobium-phosphorus-boron glasses bearing low fractions of chromium and exhibiting high toughness
US10458008B2 (en) 2017-04-27 2019-10-29 Glassimetal Technology, Inc. Zirconium-cobalt-nickel-aluminum glasses with high glass forming ability and high reflectivity
US11371108B2 (en) 2019-02-14 2022-06-28 Glassimetal Technology, Inc. Tough iron-based glasses with high glass forming ability and high thermal stability

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WO2006099808A1 (fr) 2006-09-28
TW200641151A (en) 2006-12-01
CN101203622A (zh) 2008-06-18
US20130098510A1 (en) 2013-04-25

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