US20120067464A1 - Co-ni-based alloy, method of controlling crystal of co-ni-based alloy, method of producing co-ni-based alloy, and co-ni-based alloy having controlled crystallinity - Google Patents

Co-ni-based alloy, method of controlling crystal of co-ni-based alloy, method of producing co-ni-based alloy, and co-ni-based alloy having controlled crystallinity Download PDF

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US20120067464A1
US20120067464A1 US13/231,539 US201113231539A US2012067464A1 US 20120067464 A1 US20120067464 A1 US 20120067464A1 US 201113231539 A US201113231539 A US 201113231539A US 2012067464 A1 US2012067464 A1 US 2012067464A1
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based alloy
heat treatment
alloy
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crystal
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Akihiko Chiba
Takuma OTOMO
Yasunori Akasaka
Tomoo Kobayashi
Ryo Sugawara
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Tohoku University NUC
Seiko Instruments Inc
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Assigned to TOHOKU UNIVERSITY, SEIKO INSTRUMENTS INC. reassignment TOHOKU UNIVERSITY ASSIGNMENT OF ASSIGNORS INTEREST (SEE DOCUMENT FOR DETAILS). Assignors: KOBAYASHI, TOMOO, SUGAWARA, RYO, AKASAKA, YASUNORI, OTOMO, TAKUMA, CHIBA, AKIHIKO
Publication of US20120067464A1 publication Critical patent/US20120067464A1/en
Priority to US14/966,817 priority Critical patent/US10808306B2/en
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    • 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
    • C22F1/10—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of nickel or cobalt or alloys based thereon
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D7/00—Casting ingots, e.g. from ferrous metals
    • B22D7/005—Casting ingots, e.g. from ferrous metals from non-ferrous metals
    • 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
    • C21D1/26—Methods of annealing
    • 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
    • C22C1/02—Making non-ferrous alloys by melting
    • C22C1/023—Alloys based on nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C19/00—Alloys based on nickel or cobalt
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C19/00—Alloys based on nickel or cobalt
    • C22C19/03—Alloys based on nickel or cobalt based on nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C19/00—Alloys based on nickel or cobalt
    • C22C19/03—Alloys based on nickel or cobalt based on nickel
    • C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
    • C22C19/051—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W
    • C22C19/055—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 20% but less than 30%
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C19/00—Alloys based on nickel or cobalt
    • C22C19/03—Alloys based on nickel or cobalt based on nickel
    • C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
    • C22C19/051—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W
    • C22C19/056—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 10% but less than 20%
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C19/00—Alloys based on nickel or cobalt
    • C22C19/07—Alloys based on nickel or cobalt based on cobalt
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C30/00—Alloys containing less than 50% by weight of each constituent
    • 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
    • C22F1/16—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of other metals or alloys based thereon
    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01F—MAGNETS; INDUCTANCES; TRANSFORMERS; SELECTION OF MATERIALS FOR THEIR MAGNETIC PROPERTIES
    • H01F1/00—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties
    • H01F1/01—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials
    • H01F1/03—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity
    • H01F1/12—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity of soft-magnetic materials
    • H01F1/14—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity of soft-magnetic materials metals or alloys
    • H01F1/147—Alloys characterised by their composition
    • 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
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Definitions

  • the present invention relates to a Co—Ni-based alloy, a method of controlling a crystal of a Co—Ni-based alloy, a method of producing a Co—Ni-based alloy, and a Co—Ni-based alloy having controlled crystallinity.
  • the present invention has been made in view of the conventional actual circumstances described above, and has an object to provide a Co—Ni-based alloy in which a crystal is easily controlled, a method of controlling a crystal of a Co—Ni-based alloy, a method of producing a Co—Ni-based alloy, and a Co—Ni-based alloy having controlled crystallinity.
  • the present invention has adopted the following constitution in order to solve the above-mentioned problem.
  • a Co—Ni-based alloy according to a first aspect of the present invention includes Co, Ni, Cr, and Mo, in which the Co—Ni-based alloy has a crystal texture in which a Goss orientation is a main orientation.
  • a Co—Ni-based alloy according to a second aspect of the present invention includes Co, Ni, Cr, and Mo, in which the Co—Ni-based alloy has a fine region and a deformation twin, the deformation twin being separated by the fine region.
  • a Co—Ni-based alloy according to a third aspect of the present invention includes Co, Ni, Cr, and Mo, in which the Co—Ni-based alloy has a dislocation density of 10 15 m ⁇ 2 or more.
  • the Co—Ni-based alloy according to a fourth aspect of the present invention preferably has a composition including, in terms of mass ratio: 28 to 42% of Co, 10 to 27% of Cr, 3 to 12% of Mo, 15 to 40% of Ni, 0.1 to 1% of Ti, 1.5% or less of Mn, 0.1 to 26% of Fe, 0.1% or less of C, and an inevitable impurity; and at least one kind selected from the group consisting of 3% or less of Nb, 5% or less of W, 0.5% or less of Al, 0.1% or less of Zr, and 0.01% or less of B.
  • the Co—Ni-based alloy according to a fifth aspect of the present invention preferably has a crystal texture in which a Goss orientation accounts for 35 to 55% of all orientations.
  • the Co—Ni-based alloy of the present invention is preferably produced by performing cold rolling at a reduction ratio of 15% or more.
  • a main orientation of the crystal texture after heat treatment is also preferably identical to a main orientation of the crystal texture before heat treatment.
  • a crystal texture is also preferably converted to a texture in which a plurality of regions each having a low dislocation density are present in a region having a high dislocation density, by performing heat treatment.
  • a method of controlling a crystal of a Co—Ni-based alloy according to the present invention includes: producing the Co—Ni-based alloy according to any one of the first to fifth aspects of the present invention by performing cold rolling at a reduction ratio of 15% or more to an alloy including Co, Ni, Cr, and Mo; and applying heat treatment to the Co—Ni-based alloy, thereby converting a texture of the Co—Ni-based alloy to a texture in which a plurality of regions each having a low dislocation density are present in a region having a high dislocation density so that a main orientation of a crystal texture after the heat treatment is identical to a main orientation of a crystal texture before the heat treatment.
  • the Co—Ni-based alloy preferably has a crystal texture in which a Goss orientation is a main orientation.
  • the applying of the heat treatment is preferably performed at temperature of 350° C. or more.
  • the applying of the heat treatment may also be performed at temperature of 350° C. to 750° C.
  • a method of producing a Co—Ni-based alloy having controlled crystallinity according to the present invention includes using the above-mentioned method of controlling a crystal of a Co—Ni-based alloy.
  • the present invention also provides a Co—Ni-based alloy having controlled crystallinity, which is produced by using the above-mentioned method of controlling a crystal of a Co—Ni-based alloy.
  • the present invention can provide a Co—Ni-based alloy having high mechanical strength, having excellent corrosion resistance, and being excellent as an elastic material.
  • the Suzuki effect is expressed by performing heat treatment to a Co—Ni-based alloy.
  • the Co—Ni-based alloy is recrystallized so as to have a texture in which a plurality of regions in which dislocations are extended and locked owing to the Suzuki effect, thereby having a low dislocation density are present in a region having a high dislocation density.
  • Such dislocation locking due to the Suzuki effect as described above delays dislocation recovery, and hence the main orientation of the crystal texture can remain unchanged.
  • the method of controlling a crystal of a Co—Ni-based alloy can provide a Co—Ni-based alloy which is not softened rapidly even if thermal history such as annealing is applied, has a high recrystallization temperature, and includes recrystallized grains each having a small diameter.
  • the main orientation of the crystal texture is identical to the main orientation of the crystal texture before heat treatment, which indicates that crystals are controlled.
  • the Co—Ni-based alloy obtained by the method of controlling a crystal of a Co—Ni-based alloy according to the present invention recrystallized grains grow slowly, and hence the Co—Ni-based alloy is formed by fine recrystallized grains.
  • a Co—Ni-based alloy in which characteristics such as workability are improved.
  • the Suzuki effect is expressed by heat treatment, thereby causing dislocation locking, resulting in resisting slip.
  • FIG. 1A is a schematic view illustrating dislocations pinned by solute atoms and by dislocations of different slip planes
  • FIG. 1B is a schematic view illustrating how dislocations are locked like a plane because of the extension of the dislocations due to the Suzuki effect;
  • FIG. 2A is a schematic view illustrating how a Co—Ni-based alloy according to this embodiment is recrystallized by heat treatment
  • FIG. 2B is a schematic view illustrating how a general alloy is recrystallized by heat treatment
  • FIG. 3 is a graph illustrating temperature dependence of stacking fault energy of the Co—Ni-based alloy according to this embodiment and a Co—Ni alloy;
  • FIG. 4A is a pole figure of a rolling texture (111) of a Co—Ni-based alloy in Example 1
  • FIG. 4B is a pole figure of a rolling texture (111) of a Co-35Ni alloy in Comparative Example 1;
  • FIG. 5A is a graph prepared by plotting the peak intensity of each of a Goss orientation, a Copper orientation, and a Brass orientation in the pole figure illustrated in FIG. 4A
  • FIG. 5B is a graph prepared by plotting the peak intensity of each of a Goss orientation, a Copper orientation, and a Brass orientation in the pole figure illustrated in FIG. 4B ;
  • FIG. 6 is a TEM bright-field image photograph of the Co—Ni-based alloy in Example 1;
  • FIG. 7 includes partially magnified photographs of the TEM bright-field image photograph shown in FIG. 6 ;
  • FIG. 8 includes TEM bright-field image photographs of the Co-35Ni alloy in Comparative Example 1;
  • FIG. 9 includes TEM bright-field image photographs of a Co-35Ni alloy in Comparative Example 2.
  • FIG. 10 is a graph prepared by plotting a relationship between the dislocation density and crystallite of the Co—Ni-based alloy for each cold rolling reduction ratio
  • FIG. 11A is a graph illustrating the peak intensity ratio of each orientation component in the pole figure of the rolling texture (111) of the Co—Ni-based alloy for each cold rolling reduction ratio
  • FIG. 11B is a graph illustrating the peak intensity ratio of each orientation component in the pole figure of the rolling texture (111) of the Co-35Ni alloy for each cold rolling reduction ratio;
  • FIG. 12 includes ODF maps of a Co—Ni-based alloy processed at a reduction ratio of 70% and a Co-35Ni alloy processed at a reduction ratio of 70%;
  • FIG. 13 includes measurement results by XRD or EBSD of a crystal texture before heat treatment and a crystal texture after heat treatment of an alloy in each of Example 3 and Comparative Example 3;
  • FIG. 14A is a pole figure of a crystal texture (111) before heat treatment of an alloy in Example 4
  • FIG. 14B is a pole figure of the crystal texture (111) after heat treatment of the alloy in Example 4
  • FIG. 14C is a pole figure of a crystal texture (111) before heat treatment of an alloy in Comparative Example 4
  • FIG. 14D is a pole figure of the crystal texture (111) after heat treatment of the alloy in Comparative Example 4;
  • FIG. 15 is a line graph illustrating grain growth depending on the heat treatment time of an alloy in each of Example 5 and Comparative Example 5;
  • FIG. 16 is an isothermal recrystallization curve of the alloy in each of Example 5 and Comparative Example 5;
  • FIG. 17A includes photographs showing measurement results of KAM images by EBSD of the alloy in Example 5 for respective heat treatment times
  • FIG. 17B includes photographs showing measurement results of KAM images by EBSD of the alloy in Comparative Example 5 for respective heat treatment times;
  • FIG. 18 is a graph illustrating a change of hardness before and after heat treatment of an alloy in each of Example 6 and Comparative Example 6;
  • FIG. 19A is a TEM bright-field image photograph of the alloy before heat treatment in Example 6
  • FIG. 19B is a TEM bright-field image photograph of the alloy after heat treatment in Example 6
  • FIG. 19C is a TEM bright-field image photograph of the alloy before heat treatment in Comparative Example 6
  • FIG. 19D is a TEM bright-field image photograph of the alloy after heat treatment in Comparative Example 6;
  • FIG. 20 is a graph prepared by plotting a change of hardness depending on a heat treatment time of an alloy in each of Example 7 and Comparative Example 7;
  • FIG. 21 is a graph prepared by plotting a change of hardness depending on a heat treatment temperature of an alloy in each of Example 8 and Comparative Example 8.
  • a Co—Ni-based alloy according to this embodiment includes Co, Ni, Cr, and Mo, in which the Co—Ni-based alloy has a crystal texture in which a Goss orientation ⁇ 110 ⁇ 001> (hereinafter, simply referred to as Goss orientation) is a main orientation.
  • the crystal texture of the Co—Ni-based alloy according to this embodiment mainly includes, as orientation factors, in addition to the Goss orientation, a Brass orientation ⁇ 110 ⁇ 112> (hereinafter, simply referred to as Brass orientation) and a Copper orientation ⁇ 211 ⁇ 111> (hereinafter, simply referred to as Copper orientation).
  • the main orientation of a crystal texture can be decided by determining the orientations of crystal grains based on three stereographic projection views such as (111), (001), and (110). For example, by comparing the peak intensity of each orientation in the pole figure of the crystal texture (111), the orientation that exhibits the highest peak intensity can be determined as the main orientation of the crystal texture.
  • the Goss orientation preferably accounts for 35 to 55% of all orientation factors.
  • the Co—Ni-based alloy according to this embodiment is preferably subjected to cold rolling at a reduction ratio of 15% or more, and is more preferably subjected to cold rolling at a reduction ratio of 15 to 90%.
  • the Co—Ni-based alloy can have a Goss orientation as the main orientation of its crystal texture.
  • a Brass orientation sometimes develops, and hence the reduction ratio is preferably controlled to 90% or less.
  • the Co—Ni-based alloy according to this embodiment preferably has a composition including, in terms of mass ratio: 28 to 42% of Co, 10 to 27% of Cr, 3 to 12% of Mo, 15 to 40% of Ni, 0.1 to 1% of Ti, 1.5% or less of Mn, 0.1 to 26% of Fe, 0.1% or less of C, and an inevitable impurity; and at least one kind selected from the group consisting of 3% or less of Nb, 5% or less of W, 0.5% or less of Al, 0.1% or less of Zr, and 0.01% or less of B.
  • the reason why the composition is limited to such range is described below.
  • Co per se has a large work-hardening capability, and hence Co has a reducing effect on the fragility of edge cutting, an increasing effect on the fatigue strength, and an increasing effect on the high-temperature strength.
  • the content of Co is less than 28%, those effects are weakly exhibited.
  • the content of Co is more than 42% in this composition, a matrix becomes too hard, with the result that working on the alloy becomes difficult and a face-centered cubic lattice phase becomes unstable with respect to a hexagonal close-packed lattice phase.
  • the content of Co was set to 28 to 42%.
  • Cr is an essential component for ensuring the corrosion resistance and has a reinforcing effect on a matrix. However, if the content of Cr is less than 10%, the effect of imparting excellent corrosion resistance is weakly exhibited. If the content of Cr is more than 27%, the workability on and toughness of the alloy sharply lower. Thus, the content of Cr was set to 10 to 27%.
  • Mo has a reinforcing effect on a matrix by forming a solid solution with the matrix, an increasing effect on the work-hardening capability, and an enhancing effect on the corrosion resistance in the coexistence with Cr.
  • the content of Mo is less than 3%, desired effects are not provided.
  • the content of Mo is more than 12%, the workability sharply lowers and a fragile a phase is apt to be generated.
  • the content of Mo was set to 3 to 12%.
  • Ni has a stabilizing effect on a face-centered cubic lattice phase, a maintaining effect on the workability, and an enhancing effect on the corrosion resistance.
  • the content of Ni is less than 15%, providing a stabilized face-centered cubic lattice phase is difficult. If the content of Ni is more than 40%, the mechanical strength lowers. Thus, the content of Ni was set to 15 to 40%.
  • Ti has strong effects of deoxidation, denitrification, and desulfurization, and has a miniaturizing effect on an ingot texture. However, if the content of Ti is less than 0.1%, those effects are weakly exhibited. If the content is, for example, 1%, no problem occurs. If the content of Ti is too large, the amount of inclusions increases in the alloy, or an ⁇ phase (Ni 3 Ti) is precipitated, resulting in a reduction in toughness. Thus, the content of Ti was set to 0.1 to 1%.
  • Mn has the effects of deoxidation and desulfurization, and a stabilizing effect on a face-centered cubic lattice phase. However, if the content of Mn is too large, the corrosion resistance and the oxidation resistance deteriorate. Thus, the content of Mn was set to 1.5% or less.
  • the content of Fe is too large, the oxidation resistance lowers.
  • priority was given to the reinforcing effect on a matrix by forming a solid solution with the matrix rather than the oxidation resistance, and hence the upper limit of the content of Fe was set to 26%.
  • the content of Fe was set to 0.1 to 26%.
  • C forms a solid solution with a matrix and, in addition, has a preventing effect on grain coarsening by forming carbides with Cr, Mo, Nb, W, or the like.
  • the content of C was set to 0.1% or less.
  • Nb has a reinforcing effect on a matrix by forming a solid solution with the matrix and an increasing effect on the work-hardening capability.
  • a ⁇ phase or a ⁇ phase (Ni 3 Nb) is precipitated, resulting in a reduction in toughness.
  • the content of Nb, if any, was set to 3% or less.
  • W has a reinforcing effect on a matrix by forming a solid solution with the matrix and a significant increasing effect on the work-hardening capability.
  • the content of W is more than 5%, a G phase is precipitated, resulting in a reduction in toughness.
  • Al has the effect of deoxidation and an enhancing effect on the oxidation resistance. However, if the content of Al is too large, the corrosion resistance deteriorates, for example. Thus, the content of Al, if any, was set to 0.5% or less.
  • Zr has an enhancing effect on the hot workability by increasing the strength of a crystal grain boundary at high temperatures.
  • the content of Zr was set to 0.1% or less.
  • B has an improving effect on the hot workability. However, if the content of B is too large, the hot workability lowers in reverse, resulting in easy break of the alloy. Thus, the content of B, if any, was set to 0.01% or less.
  • the Co—Ni-based alloy according to this embodiment more preferably includes 0.1 to 3% of Fe and 3% or less of Nb selected as the at least one kind. That is, more preferred is a Co—Ni-based alloy having a composition including, in terms of mass ratio, 28 to 42% of Co, 10 to 27% of Cr, 3 to 12% of Mo, 15 to 40% of Ni, 0.1 to 1% of Ti, 1.5% or less of Mn, 0.1 to 3% of Fe, 0.1% or less of C, 3% or less of Nb, and an inevitable impurity.
  • the Co—Ni-based alloy having the composition described above by setting the upper limit of Fe to 3%, the oxidation resistance can be prevented from lowering more effectively.
  • a face-centered cubic lattice (fcc) alloy undergoes some processing, a Brass orientation usually develops in the crystal texture of the alloy rather than a Goss orientation. Further, it is known that the recrystallization of the alloy after heat treatment generally results in the change of its crystal texture. Thus, the change of the crystal texture described above made it difficult to control the crystals of the alloy.
  • the recrystallized texture probably has a certain orientation in its core, and hence the main orientation of its crystal texture is maintained.
  • the crystals of the alloy are controlled, it is not necessary to consider the change of the crystal texture, and it is enough to consider only the parameters of a heat treatment temperature and time, and hence the crystals of the alloy can easily be controlled.
  • the reason why the main orientation of the crystal texture of the Co—Ni-based alloy according to this embodiment does not change by heat treatment is that the Co—Ni-based alloy according to this embodiment is an alloy which expresses the Suzuki effect by undergoing heat treatment.
  • the Suzuki effect is one of the interactions between a dislocation and a solute atom.
  • Dislocations in a face-centered cubic lattice (fcc) alloy and a hexagonal close-packed lattice (hcp) alloy are extended dislocations in many cases, and hence an extended dislocation portion has a different energy state to a certain extent from a surrounding portion, and solute atoms are segregated in the extended dislocation portion.
  • fcc face-centered cubic lattice
  • hcp hexagonal close-packed lattice
  • the interaction between a dislocation and a solute atom is generally called a chemical interaction or the Suzuki effect.
  • the expression of the Suzuki effect contributes to improving mechanical characteristics such as the hardness and tensile strength of an alloy.
  • dislocations when dislocations are locally pinned by solute atoms and by dislocations of different slip planes, the dislocations do not slip easily because of the pinning, but the dislocations may project between pinned portions and the dislocations are not locked very strongly.
  • the inventors of the present invention have made various studies. As a result, the inventors have found that the Suzuki effect can be expressed in the Co—Ni-based alloy according to this embodiment, and have found that a Co—Ni-based alloy having excellent characteristics can be provided by taking advantage of the Suzuki effect.
  • formed in the alloy is a texture in which, in region A having a high dislocation density (high-density dislocation region), there exist a plurality of regions B each having a low dislocation density (low-density dislocation regions) in which extended dislocations are induced by the Suzuki effect, forming a state in which the dislocations do not recover.
  • region A having a high dislocation density high-density dislocation region
  • regions B each having a low dislocation density (low-density dislocation regions) in which extended dislocations are induced by the Suzuki effect forming a state in which the dislocations do not recover.
  • the dislocations recover slowly, and hence recrystallized grains grow slowly.
  • Dislocations are not extended in the high dislocation density region A, and hence this region serves as a grain core for developing recrystallization, but the plurality of low-density dislocation regions B contribute to preventing the rapid progress of grain growth, thereby forming fine recrystallized grains. Further, in the regions B in which dislocations are extended owing to the Suzuki effect and the dislocations recover slowly, even if recrystallization develops, recrystallized grains grow while the crystal texture of the regions B is maintained. Thus, in the Co—Ni-based alloy according to this embodiment, the main orientation of the crystal texture before heat treatment can remain the same even after recrystallization is caused by the heat treatment.
  • the Co—Ni-based alloy according to this embodiment include Co, Ni, Cr, and Mo, have fine regions a and deformation twins b, and have a crystal texture in which the deformation twins b are separated by the fine regions a.
  • FIG. 6 shows a transmission electron microscope (TEM) bright-field image photograph of the crystal texture of a Co—Ni-based alloy sample obtained in an Example described below.
  • FIG. 7 shows TEM bright-field image photographs taken by partially magnifying the crystal texture shown in FIG. 6 . Note that the Co—Ni-based alloy shown in FIG. 6 and FIG. 7 is formed by applying cold rolling at a reduction ratio of 70%. In regions A and B in FIG.
  • each diffraction pattern shows a ring shape, and hence it is found that these regions are polycrystalline fine regions a having various orientations.
  • regions C and D looking like lines in FIG. 7 each diffraction pattern shows that diffraction spots have a dot shape or regularity.
  • the region is formed of deformation twins b having the same orientation in view of the positional relationship between the diffraction spots.
  • the Co—Ni-based alloy according to this embodiment has, as shown in the wide view photograph of FIG. 6 , a crystal texture in which deformation twins b looking like lines are separated like a grid by fine regions a represented by broken lines.
  • the Co—Ni-based alloy according to this embodiment has, as shown in FIG. 6 and FIG. 7 deformation twins b separated like a grid, the deformation twins being very fine deformation twins compared with those of general binary alloys such as a Co-35Ni alloy.
  • the fact that such very fine deformation twins are preferentially formed and large shear bands are not formed, thereby delaying the development of a Brass orientation, is also probably a cause for a Goss orientation to be maintained. That is, the regions represented by broken lines in FIG. 6 and the regions A and B in FIG.
  • the Co—Ni-based alloy according to this embodiment has a feature that its dislocation density is 10 15 m ⁇ 2 or more.
  • General alloys each have a dislocation density of about 10 10 to 10 12 m ⁇ 2 after usual heat treatment, and have a dislocation density of about 10 12 to 10 14 m ⁇ 2 even after cold rolling processing is performed.
  • the Co—Ni-based alloy according to this embodiment has a relatively high dislocation density compared with dislocation densities of general alloys, and moreover, the Co—Ni-based alloy has such polycrystalline fine regions and fine deformation twins as described above. Thus, dislocations are formed in the Co—Ni-based alloy more easily than in general alloys, probably resulting in its higher dislocation density.
  • FIG. 14A illustrates a pole figure of the crystal texture (111) of the Co—Ni-based alloy (a cold rolling reduction ratio of 90%) according to this embodiment.
  • FIG. 14B illustrates a pole figure of the crystal texture (111) of the Co—Ni-based alloy having been subjected to heat treatment at 1,050° C. for 1 hour. As illustrated in FIG. 14A and FIG.
  • an alloy including the composition described above is subjected to vacuum melting in a vacuum melting furnace, followed by furnace cooling to produce an ingot.
  • the resultant ingot is subjected to hot casting by a general method, followed by annealing.
  • cold rolling is performed at a reduction ratio of 15% or more, thereby producing the Co—Ni-based alloy according to this embodiment.
  • cold rolling is performed at a reduction ratio of 15% or more, it is possible to obtain a Co—Ni-based alloy having a Goss orientation as the main orientation of its crystal texture.
  • cold rolling is performed at a reduction ratio of more than 90%, a Brass orientation tends to develop easily, and hence cold rolling is preferably performed at a reduction ratio of 90% or less. Note that the crystal texture of the present invention is not formed after hot casting and annealing.
  • Heat treatment conditions can be altered arbitrarily. Heat treatment is preferably performed at temperature of 350° C. or more because the Suzuki effect is expressed, thereby extending and locking dislocations, the recovery of the dislocations is delayed, the main orientation of the crystal texture of the Co—Ni-based alloy remains unchanged, and hence a Goss orientation can be still maintained as the main orientation after the heat treatment. Further, as the Suzuki effect is expressed in the early stage of heating, the upper limit of heat treatment temperature is not particularly limited. The main orientation of the crystal texture can remain unchanged even at as high a temperature as, for example, about 1,050° C., but recrystallization is apt to be more dominant at 800° C.
  • the temperature of the heat treatment is more preferably in the range of 350° C. to 750° C.
  • the Suzuki effect can be effectively expressed, thereby allowing the main orientation of the crystal texture to remain unchanged.
  • the time of the heat treatment can be altered arbitrarily depending on the temperature of the heat treatment, and is set to preferably 0.5 hour or more and 3.5 hours or less, more preferably 0.5 hour or more and 1.5 hours or less.
  • a Co—Ni-based alloy can be produced while the crystals of the Co—Ni-based alloy are being controlled.
  • heat treatment does not change the main orientation of the crystal texture of the alloy, and hence it becomes possible to control the crystals of the alloy by performing the heat treatment while considering only the temperature and time of the heat treatment.
  • the Suzuki effect is expressed by performing heat treatment, thereby, as illustrated in FIG. 2A , causing the crystal texture of the Co—Ni-based alloy to recrystallize as a texture in which a plurality of regions B each having a low dislocation density are present in a region A having a high dislocation density.
  • the main orientation of the crystal texture can remain unchanged.
  • FIG. 19A shows a TEM bright-field image photograph of a Co—Ni-based alloy (a reduction ratio of 15%) having a component composition including 31 mass % of Ni, 19 mass % of Cr, 10.1 mass % of Mo, 2 mass % of Fe, 0.8 mass % of Ti, 1 mass % of Nb, and Co accounting for the balance, as one Example of the Co—Ni-based alloy according to this embodiment.
  • FIG. 19B shows a photograph of the crystal texture of the above-mentioned Co—Ni-based alloy to which heat treatment was applied at 700° C. for 1 hour. The heat treatment causes many stacking faults that look like small vertical lines to occur as shown in FIG. 19B , and it is found that the Suzuki effect has contributed to extending and locking dislocations.
  • FIG. 19C shows a TEM bright-field image photograph of a Co-35Ni alloy (a reduction ratio of 15%).
  • FIG. 19D shows a TEM bright-field image photograph of the above-mentioned Co-35Ni alloy to which heat treatment was applied at 350° C. for 1 hour. As shown in FIG. 19D , in the Co-35Ni alloy, dislocations that look like lines are decreased by performing the heat treatment, indicating the recovery of the dislocations.
  • FIG. 17A shows electron backscatter diffraction (EBSD) images of the Co—Ni-based alloy (a reduction ratio of 70%) according to this embodiment to which alloy heat treatment was applied at 800° C. for treatment times of 5 minutes, 20 minutes, and 60 minutes, respectively.
  • EBSD electron backscatter diffraction
  • the main orientation of its crystal texture is identical to the main orientation of the crystal texture before heat treatment, which indicates that crystals are controlled.
  • X-ray diffraction measurement was carried out using an X-ray diffractometer “monochromator” manufactured by Koninklijke Philips Electronics N.V.
  • a dislocation density was calculated by using a modified Warren-Averbach method (J. Phys. Chem. Sol., 62, 2001, 1935-1941) which was established by introducing a contrast factor C (constant for crystal face dependence of strain sensitivity) to the Warren-Averbach method proposed by T. Unger.
  • the X-ray diffraction profile of each sample is measured, and the background is subtracted from the raw profile. After that, measurement error factors are corrected, the Fourier transform is performed, and a Fourier coefficient A(L) corresponding to a Fourier length (L) is obtained from each diffraction profile. Then, the dislocation density and attribute parameter of the texture can be calculated by using the Warren-Averbach calculating formulae represented by the Equation (1) to Equation (3) described below.
  • Equation (1) to Equation (3) b represents a Burgers vector
  • R e represents the size of a strain field caused by dislocation
  • ⁇ represents a dislocation density
  • K 2 sin ⁇ / ⁇
  • O represents a constant based on a distance between dislocations
  • a s (L) represents a Fourier coefficient based on a crystal grain diameter
  • L represents a distance satisfying a coherent diffraction condition (Fourier length).
  • Equation (2) shows, X(L) is a coefficient of a linear term of Equation (1), and Equation (2) can be modified to Equation (3).
  • the dislocation density ⁇ can be determined.
  • an X-ray diffractometer “monochromator” manufactured by Koninklijke Philips Electronics N.V. was used to measure an X-ray diffraction profile, and Origin (manufactured by OriginLab Corporation) was used as analysis software.
  • K represents a Scherrer constant
  • ⁇ represents the wavelength of an X-ray used
  • ⁇ represents the half-value width of an X-ray diffraction peak
  • ⁇ represents an X-ray incident angle 2 ⁇ . Note that the crystallite size refers to the size of a subgrain.
  • FIG. 3 illustrates the stacking fault energy (SFE) of alloy systems, the energy being necessary for causing phase transformation from a ⁇ phase, which has a face-centered cubic lattice (fcc) structure, to an ⁇ phase, which has a hexagonal close-packed lattice (hcp) structure.
  • SFE stacking fault energy
  • ⁇ G ⁇ represents a Gibbs energy change associated with ⁇ transformation
  • ⁇ ⁇ / ⁇ represents the interface energy of a ⁇ / ⁇ boundary
  • Used for ⁇ G ⁇ was a value calculated by using Thermo-Calc (manufactured by Thermo-Calc Software: ver. 4.1.3.41, database: FE ver. 6).
  • the temperature dependence of the interface energy in Equation (4) is small and the value of the temperature dependence is constant in transition metal irrespective of temperature.
  • 2 ⁇ ⁇ / ⁇ 15 mJm ⁇ 2 , which is a surface energy term, was used to make a calculation.
  • a high-frequency vacuum induction melting furnace was used to blend and melt the following each element, with a component composition of 31 mass % of Ni, 19 mass % of Cr, 10.1 mass % of Mo, 2 mass % of Fe, 0.8 mass % of Ti, 1 mass % of Nb, and Co accounting for the balance, followed by furnace cooling.
  • the resultant ingot was subjected to hot casting at 100° C. and then subjected to annealing at 1,050° C., providing an alloy material (hereinafter, referred to as “alloy material for examples”), which was used to produce each Co—Ni-based alloy.
  • a high-frequency vacuum induction melting furnace was used to blend and melt the following each element, with a component composition of 35 mass % of Ni and Co accounting for the balance, followed by furnace cooling.
  • the resultant ingot was subjected to hot casting at 100° C. and then subjected to annealing at 1,000° C., providing an alloy material (hereinafter, referred to as “alloy material for comparative examples”), which was used to produce each Co-35Ni alloy.
  • heat treatment in the following examples and comparative examples was performed in a vacuum at a temperature rise speed of 8° C./second, and at a cooling speed of 12° C./second.
  • a Co—Ni-based alloy was produced by applying cold rolling to the alloy material for examples at a reduction ratio of 70%.
  • a Co-35Ni alloy was produced by applying cold rolling to the alloy material for comparative examples at a reduction ratio of 70%.
  • a Co-35Ni alloy was produced by applying cold rolling to the alloy material for comparative examples at a reduction ratio of 50%.
  • FIG. 4A is a pole figure of the rolling texture (111) of the Co—Ni-based alloy in Example 1
  • FIG. 4B is a pole figure of the rolling texture (111) of the Co-35Ni alloy in Comparative Example 1.
  • FIG. 5A is a graph prepared by plotting the peak intensity of each of a Goss orientation, a Copper orientation, and a Brass orientation in the pole figure illustrated in FIG. 4A
  • FIG. 5B is a graph prepared by plotting the peak intensity of each of a Goss orientation, a Copper orientation, and a Brass orientation in the pole figure illustrated in FIG. 4B .
  • FIG. 6 is a TEM bright-field image photograph of the Co—Ni-based alloy in Example 1
  • FIG. 7 includes magnified photographs of the TEM bright-field image photograph of FIG. 6 .
  • each diffraction pattern shows a ring shape, and hence it is found that these regions are polycrystalline fine regions having various orientations.
  • each diffraction pattern shows that diffraction spots have a dot shape or regularity.
  • the regions are formed of deformation twins having the same orientation in view of the positional relationship between the diffraction spots.
  • deformation twins b looking like lines were separated like a grid by fine regions a represented by broken lines.
  • FIG. 8 includes TEM bright-field image photographs of the Co-35Ni alloy in Comparative Example 1
  • FIG. 9 includes TEM bright-field image photographs of the Co-35Ni alloy in Comparative Example 2.
  • the crystal texture of the Co-35Ni alloy in each of Comparative Examples 1 and 2 was a texture including large deformation twins having a wide band shape, and such deformation twins separated like a grid as found in the Co—Ni-based alloy in Example 1 were not found.
  • FIG. 10 illustrates a graph prepared by plotting a relationship between the dislocation density and crystallite of the Co—Ni-based alloy for each cold rolling reduction ratio. Note that, in FIG. 10 , the circular symbols each represents a point plotted for a crystallite size and the rhombic symbols each represents a point plotted for a dislocation density.
  • ODFs 3-D crystal orientation distribution functions
  • the Co—Ni-based alloys of No. 1 to No. 7 and the Co-35Ni alloys in comparative examples each included a Copper twin orientation and a Dillamore orientation, in addition to a Goss orientation, a Brass orientation, and a Copper orientation.
  • FIGS. 11A and 11B each illustrate a graph prepared by plotting the peak intensity ratios of the Goss orientation, the Copper orientation, and the Brass orientation out of the various orientation components obtained.
  • FIG. 11A illustrates the peak intensity ratio of each orientation component of the Co—Ni-based alloys of No. 1 to No. 7, and FIG. 11B illustrates the peak intensity ratio of each orientation component of the Co-35Ni alloys in comparative examples. From the results of FIG.
  • FIG. 12 includes ODF maps of the Co—Ni-based alloy processed at a reduction ratio of 70% and the Co-35Ni alloy processed at a reduction ratio of 70%.
  • FIG. 13 illustrates the EBSD result of the Co—Ni-based alloy to which heat treatment at 800° C. for 5 minutes was applied and the EBSD result of the Co—Ni-based alloy to which heat treatment at 800° C. for 60 minutes was applied, together with a (111) pole figure of the Co—Ni-based alloy before heat treatment. Consequently, as illustrated in FIG. 13 , even though the Co—Ni-based alloy in Example 3 was subjected to the heat treatment at 800° C.
  • the crystal texture of the Co—Ni-based alloy in Example 3 mainly included a Goss orientation after the heat treatment, indicating that the main orientation of the crystal texture before the heat treatment remained unchanged.
  • FIG. 13 illustrates the EBSD result of the Co-35Ni alloy to which heat treatment at 350° C. for 5 minutes was applied and the EBSD result of the Co-35Ni alloy to which heat treatment at 350° C. for 60 minutes was applied, together with a (111) pole figure of the Co-35Ni alloy before heat treatment. Consequently, as illustrated in FIG. 13 , by applying heat treatment at 350° C.
  • FIG. 14A is a (111) pole figure of the Co—Ni-based alloy before the heat treatment
  • FIG. 14B is a (111) pole figure of the Co—Ni-based alloy after the heat treatment.
  • FIG. 14A and FIG. 14B even though the Co—Ni-based alloy in Example 4 was subjected to the heat treatment at 1,050° C. for 1 hour, there was no significant change between both peaks in the (111) pole figures, and Goss ⁇ 110 ⁇ 001> was present in the rolling direction RD.
  • the crystal texture of the Co—Ni-based alloy in Example 4 mainly included a Goss orientation after the heat treatment, indicating that the main orientation of the crystal texture before the heat treatment remained unchanged.
  • FIG. 14C is a (111) pole figure of the SUS316L-CR before the heat treatment
  • FIG. 14D is a (111) pole figure of the SUS316L-CR after the heat treatment.
  • FIG. 14C and FIG. 14D by applying the heat treatment at 1,050° C. for 1 hour to the SUS316L, there was a remarkable change between the (111) pole figures.
  • the main orientation of the crystal texture of the SUS316L was changed by performing the heat treatment.
  • FIG. 15 is a graph prepared by plotting, with respect to the heat treatment times, the average diameters of recrystallized grains determined based on the results of the EBSD measurements.
  • FIG. 16 is a graph prepared by plotting, with respect to the heat treatment times, the fractions of a recrystallization region determined based on the results of the EBSD measurements.
  • FIG. 17A shows kernel average misorientation (KAM) images by EBSD of the Co—Ni-based alloys to which heat treatment at 800° C. was performed for treatment times of 5 minutes, 20 minutes, and 60 minutes.
  • KAM kernel average misorientation
  • FIG. 15 plots, with respect to the heat treatment times, the average diameters of recrystallized grains determined based on the results of the EBSD measurements.
  • FIG. 16 plots, with respect to the heat treatment times, the fractions of a recrystallization region determined based on the results of the EBSD measurements.
  • FIG. 17B shows KAM images by EBSD of the Co-35Ni alloys to which heat treatment at 350° C. was performed for treatment times of 0.5 minute, 2.5 minutes, and 60 minutes.
  • FIG. 15 , FIG. 16 , and FIG. 17 show that, in the Co-35Ni alloy in Comparative Example 5, recrystallization and grain growth progressed rapidly, and recrystallized grains grew so as to each have a diameter of about 10 ⁇ m by the heat treatment at 350° C. for 60 minutes.
  • the diameter of each recrystallized grain changed only slightly by changing the heat treatment time, and the diameter of the each recrystallized grain was about 2 ⁇ m even after the heat treatment at 800° C. for 60 minutes was performed. From this result, it is estimated that, in the Co—Ni-based alloy in Example 5, as illustrated in FIG.
  • FIG. 18 is a graph prepared by plotting the hardness [HV] of the Co—Ni-based alloy before the heat treatment and the hardness [HV] of the Co—Ni-based alloy after the heat treatment.
  • FIG. 19A is a TEM bright-field image photograph of the Co—Ni-based alloy before the heat treatment
  • FIG. 19B is a TEM bright-field image photograph of the Co—Ni-based alloy after the heat treatment.
  • FIG. 18 plots the hardness [HV] of the Co-35Ni alloy before the heat treatment and the hardness [HV] of the Co-35Ni alloy after the heat treatment.
  • FIG. 19C is a TEM bright-field image photograph of the Co-35Ni alloy before the heat treatment
  • FIG. 19D is a TEM bright-field image photograph of the Co-35Ni alloy after the heat treatment.
  • FIG. 18 shows that, by applying the heat treatment at 350° C. for 1 hour to the Co-35Ni alloy in Comparative Example 6, the hardness of the alloy remarkably lowered.
  • the heat treatment at 700° C. for 1 hour to the Co—Ni-based alloy in Example 6 the hardness of the alloy improved. From this result, it is estimated that, in the Co—Ni-based alloy in Example 6, the Suzuki effect caused by heat treatment induced the locking of dislocations, thereby preventing easy slip of the dislocations, and resulting in the improvement of the hardness.
  • FIG. 19A and FIG. 19B by applying the heat treatment at 700° C.
  • FIG. 20 shows that, by performing heat treatment to the Co-35Ni alloy, its hardness remarkably lowered.
  • heat treatment at 800° C. for one minute induced the expression of the Suzuki effect, leading to the improvement of the hardness, and then, the Suzuki effect delayed dislocation recovery, resulting in a gradual change in the hardness.
  • FIG. 21 shows that, by performing heat treatment to the Co-35Ni alloy, its hardness remarkably lowered.
  • heat treatment at 350° C. or more improved the hardness, and hence the heat treatment at 350° C. or more was confirmed to induce the expression of the Suzuki effect.
  • the Suzuki effect is expressed in the early stage of heating, the heating temperature may be 1,050° C., but recrystallization becomes dominant at 800° C. or more over dislocation locking induced by the Suzuki effect, resulting in the reduction of the hardness of the Co—Ni-based alloy.
  • the heating temperature was confirmed to be more preferably in the range of 350° C. to 750° C.

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Cited By (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20150337694A1 (en) * 2012-07-06 2015-11-26 Man Diesel & Turbo, Filial Af Man Diesel & Turbo Se, Tyskland Exhaust valve spindle for an exhaust valve in an internal combustion engine
US20180272451A1 (en) * 2017-03-21 2018-09-27 Kennametal Inc. Imparting wear resistance to superalloy articles
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US11591684B2 (en) 2018-12-13 2023-02-28 Unison Industries, Llc Nickel-cobalt material and method of forming
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Citations (10)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3811872A (en) * 1971-04-21 1974-05-21 Int Nickel Co Corrosion resistant high strength alloy
US3816106A (en) * 1972-08-25 1974-06-11 Int Nickel Co Strong, corrosion resistant alloy
US4877461A (en) * 1988-09-09 1989-10-31 Inco Alloys International, Inc. Nickel-base alloy
JPH08311587A (ja) * 1995-03-14 1996-11-26 Seiko Instr Inc 超塑性Ｎｉ−Ｃｏ基合金およびその超塑性加工方法
JPH10140279A (ja) * 1996-09-13 1998-05-26 Seiko Instr Inc Ｃｏ−Ｎｉ基合金
US20040025989A1 (en) * 2000-09-19 2004-02-12 Akihiko Chiba Co-ni base heat-resistant alloy and method for producing thereof
US20040033158A1 (en) * 2002-07-05 2004-02-19 Akihiko Chiba Precipitation hardened Co-Ni based heat-resistant alloy and production method therefor
US20050016645A1 (en) * 2003-03-26 2005-01-27 Osamu Takahashi Co-Ni-base alloy
US20080166258A1 (en) * 2006-09-25 2008-07-10 Nippon Seisen Co., Ltd. Heat-resistant alloy spring and Ni-based alloy wire therefor
US20100048322A1 (en) * 2008-08-21 2010-02-25 Ryo Sugawara Golf club head, face of the golf club head, and method of manufacturing the golf club head

Family Cites Families (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3767385A (en) * 1971-08-24 1973-10-23 Standard Pressed Steel Co Cobalt-base alloys
JPS60187652A (ja) 1984-03-05 1985-09-25 Seiko Electronic Components Ltd 高弾性合金の製造方法
EP0365716A1 (en) * 1984-08-08 1990-05-02 Latrobe Steel Company Nickel-cobalt base alloys
JP3041585B2 (ja) * 1995-04-28 2000-05-15 セイコーインスツルメンツ株式会社 ぜんまいの製造方法
JP2004292918A (ja) * 2003-03-27 2004-10-21 Japan Science & Technology Agency 高強度Ｃｏ−Ｎｉ基合金の製造方法
JP4510542B2 (ja) * 2004-07-30 2010-07-28 セイコーインスツル株式会社 刃物及びその製造方法
JP5103107B2 (ja) * 2007-09-18 2012-12-19 セイコーインスツル株式会社 高弾性合金

Patent Citations (10)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3811872A (en) * 1971-04-21 1974-05-21 Int Nickel Co Corrosion resistant high strength alloy
US3816106A (en) * 1972-08-25 1974-06-11 Int Nickel Co Strong, corrosion resistant alloy
US4877461A (en) * 1988-09-09 1989-10-31 Inco Alloys International, Inc. Nickel-base alloy
JPH08311587A (ja) * 1995-03-14 1996-11-26 Seiko Instr Inc 超塑性Ｎｉ−Ｃｏ基合金およびその超塑性加工方法
JPH10140279A (ja) * 1996-09-13 1998-05-26 Seiko Instr Inc Ｃｏ−Ｎｉ基合金
US20040025989A1 (en) * 2000-09-19 2004-02-12 Akihiko Chiba Co-ni base heat-resistant alloy and method for producing thereof
US20040033158A1 (en) * 2002-07-05 2004-02-19 Akihiko Chiba Precipitation hardened Co-Ni based heat-resistant alloy and production method therefor
US20050016645A1 (en) * 2003-03-26 2005-01-27 Osamu Takahashi Co-Ni-base alloy
US20080166258A1 (en) * 2006-09-25 2008-07-10 Nippon Seisen Co., Ltd. Heat-resistant alloy spring and Ni-based alloy wire therefor
US20100048322A1 (en) * 2008-08-21 2010-02-25 Ryo Sugawara Golf club head, face of the golf club head, and method of manufacturing the golf club head

Cited By (8)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20150337694A1 (en) * 2012-07-06 2015-11-26 Man Diesel & Turbo, Filial Af Man Diesel & Turbo Se, Tyskland Exhaust valve spindle for an exhaust valve in an internal combustion engine
US9714724B2 (en) * 2012-07-06 2017-07-25 Man Diesel & Turbo. Filial Af Man Diesel & Turbo Se. Tyskland Exhaust valve spindle for an exhaust valve in an internal combustion engine
US20180272451A1 (en) * 2017-03-21 2018-09-27 Kennametal Inc. Imparting wear resistance to superalloy articles
US11117208B2 (en) * 2017-03-21 2021-09-14 Kennametal Inc. Imparting wear resistance to superalloy articles
US11591684B2 (en) 2018-12-13 2023-02-28 Unison Industries, Llc Nickel-cobalt material and method of forming
EP3708687A1 (en) * 2019-03-14 2020-09-16 Unison Industries, LLC Thermally stabilized nickel-cobalt materials and methods of thermally stabilizing the same
US11807929B2 (en) 2019-03-14 2023-11-07 Unison Industries, Llc Thermally stabilized nickel-cobalt materials and methods of thermally stabilizing the same
CN115725887A (zh) * 2022-11-21 2023-03-03 广东省科学院中乌焊接研究所 一种中熵合金及其制备方法和应用

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JP5736140B2 (ja) 2015-06-17
US20160097115A1 (en) 2016-04-07
CN102400014B (zh) 2016-02-24

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