WO2013118285A1 - ALLIAGE À BASE DE Co-Cr-Mo ET PROCÉDÉ DE PRODUCTION D'ALLIAGE À BASE DE Co-Cr-Mo - Google Patents

ALLIAGE À BASE DE Co-Cr-Mo ET PROCÉDÉ DE PRODUCTION D'ALLIAGE À BASE DE Co-Cr-Mo Download PDF

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WO2013118285A1
WO2013118285A1 PCT/JP2012/053053 JP2012053053W WO2013118285A1 WO 2013118285 A1 WO2013118285 A1 WO 2013118285A1 JP 2012053053 W JP2012053053 W JP 2012053053W WO 2013118285 A1 WO2013118285 A1 WO 2013118285A1
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mass
based alloy
grain
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hot
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PCT/JP2012/053053
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千葉　晶彦
謙太 山中
真奈美 森
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国立大学法人東北大学
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    • 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
    • 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
    • C—CHEMISTRY; METALLURGY
    • C25—ELECTROLYTIC OR ELECTROPHORETIC PROCESSES; APPARATUS THEREFOR
    • C25F—PROCESSES FOR THE ELECTROLYTIC REMOVAL OF MATERIALS FROM OBJECTS; APPARATUS THEREFOR
    • C25F3/00—Electrolytic etching or polishing
    • C25F3/16—Polishing
    • C25F3/22—Polishing of heavy metals
    • C25F3/26—Polishing of heavy metals of refractory metals

Definitions

  • the present invention relates to a Co—Cr—Mo base alloy and a method for producing a Co—Cr—Mo base alloy.
  • Biomaterials are mainly divided into metal materials, polymer materials, and ceramics, but metal materials are superior in balance of strength and toughness compared to polymers and ceramics, so they are used as skeletal substitute materials that support the body. ing.
  • the main metal-based biomaterials currently in practical use include pure Ti / Ti alloys, austenitic stainless steels, and Co—Cr—Mo alloys. Among them, Co—Cr—Mo alloys are other metal materials. Therefore, it plays an important role as a material for artificial joints.
  • the casting material is not limited to the Co—Cr—Mo alloy, and the structure is coarse and often contains casting defects. Therefore, the casting material has poor strength and ductility. Therefore, in order to meet the recent demand for higher durability of biomedical devices such as artificial hip joints, the mechanical properties of Co—Cr—Mo alloys are improved by microstructure control using plastic working and heat treatment. There is a need.
  • Co-29Cr-6Mo (mass%) alloy which is a representative composition of the ASTM standard, easily undergoes martensitic transformation during rapid cooling after solution heat treatment, and therefore is metastable with a face-centered cubic structure (fcc) at room temperature.
  • fcc face-centered cubic structure
  • hcp hexagonal close-packed
  • the reason why Co—Cr—Mo alloys are classified as difficult-to-work alloys is due to the low plastic workability resulting from such microstructures. This has been a problem when applied to thin plates, thin pipes, or stents that require processing into wires.
  • Co—Cr—Mo base alloy when a Co—Cr—Mo base alloy is subjected to plastic working, generally 10 to 40% of Ni is added to stabilize the ⁇ phase (for example, Patent Documents). 1). Actually, Co—Ni—Cr—Mo (ASTM F562) alloy and Co—Cr—Ni—W (ASTM F90) alloy are used as the Co—Cr based alloy for stents.
  • Ni causes carcinogenicity and metal allergy
  • conventional alloys containing Ni as described in Patent Document 1 may adversely affect the human body when used as a biomaterial.
  • Ni causes carcinogenicity and metal allergy
  • conventional alloys containing Ni as described in Patent Document 1 may adversely affect the human body when used as a biomaterial.
  • the stent cannot be taken out once it is placed in the body and is implanted in the body for a long period of time, it is not desirable to add Ni, which is concerned about safety to the living body.
  • the present invention has been made by paying attention to such problems, and has high safety as a biomaterial, high strength, high ductility, and high fatigue strength. It aims at providing the manufacturing method of a base alloy.
  • the Co—Cr—Mo base alloy according to the present invention contains Cr: 25 to 35% by mass, Mo: 3 to 8% by mass, impurities: 3.2% by mass or less, and the balance being It is characterized by comprising Co.
  • the impurities preferably include N: 0.08 to 0.8 mass%.
  • the impurities are: Zr: 0 to 0.1% by mass, Si: 1% by mass or less, Mn: 1% by mass or less, C: 0.01 to 0% It may contain at least one selected from the group consisting of 3% by mass.
  • the manufacturing method of the Co—Cr—Mo base alloy according to the present invention is as follows: Cr: 25 to 35 mass%, Mo: 3 to 8 mass%, Zr: 0 to 0.1 mass%, Si: 1 mass% or less Mn: 1 mass% or less, C: 0.01 to 0.3 mass%, N: at least one impurity selected from the group consisting of 0.08 to 0.8 mass%, and the balance It is characterized by performing hot working on a raw material made of Co at 1000 ° C. or higher.
  • the hot working has a rolling reduction of 60 to 90%, and at least one of rolling, forging, swaging, and groove roll rolling. Preferably it consists of one.
  • the Co—Cr—Mo base alloy according to the present invention is preferably manufactured by the method for manufacturing a Co—Cr—Mo base alloy according to the present invention.
  • the Co—Cr—Mo base alloy according to the present invention is highly safe as a biomaterial because Ni, which is supposed to have an adverse effect on the human body, is reduced to a level at which it is inevitably contained as an impurity.
  • Co—Cr—Mo alloys have a face-centered cubic structure ⁇ phase stably in the high temperature range (1000 ° C. or higher), but the ⁇ phase has a hexagonal close packed structure due to nonthermal martensitic transformation during cooling. To metamorphosis. If an ⁇ phase formed by martensitic transformation is present in the constituent phase, stress concentration is likely to occur at the ⁇ / ⁇ interface, which becomes a starting point of fracture, and causes a decrease in room temperature ductility including cold workability.
  • the method for producing a Co—Cr—Mo base alloy and a Co—Cr—Mo base alloy according to the present invention non-thermal martensitic transformation that occurs during cooling is suppressed by containing 0.08% or more of N. And a metastable ⁇ -phase single-phase structure can be obtained even at room temperature.
  • the ⁇ -phase single-phase structure exhibits higher ductility than that having a two-phase structure of the ⁇ phase and the ⁇ phase, and can increase the fatigue strength.
  • the ⁇ phase having a face-centered cubic structure is preferably 90% or more, and more preferably 95% or more. Note that when N is added in an amount of 0.8% or more, N exists as a large amount of nitride or blowhole, which adversely affects the mechanical characteristics.
  • the method for producing a Co—Cr—Mo base alloy according to the present invention can be applied to a Co—Cr—Mo base alloy produced at a relatively low rolling reduction of 60 to 90% by 1 ⁇ 10 15 to 6 ⁇ 10 15 m.
  • a huge dislocation density of ⁇ 2 can be introduced. This is because the stacking fault energy of the Co—Cr—Mo alloy is remarkably low at about 30 mJm ⁇ 2 even in the processing temperature region (around 1000 to 1200 ° C.).
  • the stacking fault energy increases with increasing temperature, but the stacking fault energy of Co—Cr—Mo alloys is lower than that of metals / alloys (eg, austenitic stainless steel) that have been considered to have low stacking fault energy.
  • the structure after hot working includes high density (partial) dislocations, stacking faults, and deformation twins.
  • the Co—Cr—Mo based alloy according to the present invention has an equiaxed crystal grain structure, an average crystal grain size measured excluding the ⁇ 3 grain boundary is 0.1 to 100 ⁇ m, and a dislocation density of 1 ⁇ It is preferably 10 15 to 5 ⁇ 10 15 m ⁇ 2 , and the appearance ratio of the grain boundary shifted by 10 ° at the maximum from the ideal orientation relation of the ⁇ 3 grain boundary is preferably 5 to 60%.
  • the hot working is performed in multiple stages with a reduction rate of less than 20% at a time so that the total reduction rate is 60% or more. Also good.
  • it has an equiaxed crystal grain structure, the average crystal grain size measured excluding the ⁇ 3 grain boundary is 1 to 100 ⁇ m, and the dislocation density and crystallite size are 1 ⁇ 10 15 to 4 ⁇ 10 15, respectively. It is possible to produce a Co—Cr—Mo-based alloy with m ⁇ 2 , 10 to 25 nm, and a grain boundary appearance ratio of 10 to 60% that is shifted by 10 ° from the ideal orientation relationship of the ⁇ 3 grain boundary. .
  • the dislocation density in the crystal grains is further increased, the average crystal grain size measured excluding the ⁇ 3 grain boundary is 1 to 100 ⁇ m, and the dislocation density and crystallite size are 3.5 ⁇ 10 15 to 5 ⁇ 10 15 m ⁇ 2 and 5 to 15 nm, respectively, the proportion of small-angle grain boundaries with a crystal orientation difference of 15 ° or less is 30 to 60%, and the ideal orientation of the ⁇ 3 grain boundary A Co—Cr—Mo base alloy having a grain boundary appearance rate of 5 to 40% shifted by up to 10 ° from the relationship can be produced.
  • the hot working may be performed with a rolling reduction of 20% or more at a time so that the total rolling reduction is 60% or more.
  • the dynamic recrystallization phenomenon that occurs during hot working it has an equiaxed crystal grain structure and the average crystal grain size measured excluding the ⁇ 3 grain boundary is 0.1-5 ⁇ m.
  • the dislocation density and the crystallite size are 1 ⁇ 10 15 to 4 ⁇ 10 15 m ⁇ 2 and 5 to 20 nm, the proportion of small-angle grain boundaries with a crystal orientation difference of 15 ° or less is 20% or less, and the ⁇ 3 grain boundary
  • a Co—Cr—Mo base alloy having a grain boundary appearance ratio of 5 to 40% deviating from an ideal orientation relationship of 10 ° at the maximum can be produced.
  • GN boundary, GNB dislocation boundary
  • the ⁇ phase has a fiber texture including ⁇ 111> and ⁇ 001> in the length direction, and is rolled.
  • the ⁇ -phase has a texture with an orientation in which the Brass orientation ⁇ 110 ⁇ ⁇ 112> and the Goss orientation ⁇ 110 ⁇ ⁇ 001> are mixed.
  • heat treatment may be performed at 900 to 1200 ° C. after the hot working.
  • the heat treatment is preferably performed for a short time as long as no significant grain growth occurs.
  • the average crystal grain size measured excluding the ⁇ 3 grain boundary is 0.1 to 10 ⁇ m
  • the dislocation density is 2 ⁇ 10 15 m ⁇ 2 or less
  • Co-Cr-Mo in which the proportion of small-angle grain boundaries with a misorientation of 15 ° or less is 20% or less and the appearance rate of grain boundaries shifted by 10 ° at the maximum from the ideal orientation relationship of ⁇ 3 grain boundaries is 5-30%
  • a base alloy can be manufactured. The heat treatment decreases the 0.2% proof stress, but the ductility can be remarkably improved.
  • the method for producing a Co—Cr—Mo base alloy according to the present invention performs high-temperature processing such as forging and rolling at a high temperature of 1000 ° C. or higher, thereby achieving a high density of 1 ⁇ 10 15 m ⁇ 2 or higher. Dislocations can be introduced, and a Co—Cr—Mo based alloy according to the present invention having excellent room temperature strength can be produced. Further, by setting the equiaxed crystal grain structure and the ratio of the ⁇ phase to 90% or more, it is possible to suppress a decrease in ductility accompanying an increase in strength.
  • the method for producing a Co—Cr—Mo base alloy according to the present invention can control the structure by effectively utilizing the deformation characteristics inherent to the alloy and the characteristics of deformation structure formation.
  • the Co—Cr—Mo base alloy according to the present invention having high strength and high fatigue strength can be obtained by utilizing grain boundary strengthening and dislocation strengthening as the strengthening mechanism.
  • the Co—Cr—Mo base alloy according to the present invention introduces a large amount of dislocations by its manufacturing method and can ensure sufficient ductility even when the strength is increased. Is possible.
  • the method of calculating by observing a dislocation and a domain structure directly using a transmission electron microscope (TEM), and X-ray diffraction (XRD) It is roughly classified into methods for analyzing the obtained line profile.
  • TEM transmission electron microscope
  • XRD X-ray diffraction
  • the structure state is defined using the dislocation density and crystallite size obtained by XRD line profile analysis.
  • the XRD line profile analysis can acquire data that can be analyzed by a laboratory XRD apparatus.
  • the XRD line profile analysis includes the Williamson-Hall method, the Warren-Averbach method, the Garrod method, the modified Williamson-Hall method, the modified Warren-Averbach method, and the various multiple-Pulp method.
  • the Co—Cr—Mo base alloy according to the present invention is not limited to a specific analysis method, and any analysis method may be used.
  • the appearance frequency of the small-angle grain boundary and the ⁇ 3 grain boundary can be obtained by using an electron beam backscatter diffraction (EBSD) method.
  • EBSD electron beam backscatter diffraction
  • the texture can be determined by EBSD method or XRD measurement.
  • the present invention it is possible to provide a Co—Cr—Mo base alloy and a method for producing a Co—Cr—Mo base alloy having high safety as a biomaterial and having high strength, high ductility, and high fatigue strength. .
  • FIG. 1 is an X-ray diffraction pattern of (a) before hot rolling, (b) after 60% hot rolling, and (c) 92.8% hot rolling of the Co—Cr—Mo based alloy shown in FIG. 2 is a graph of a nominal stress-nominal strain curve obtained by a tensile test of the Co—Cr—Mo based alloy shown in FIG. 1 obtained by EBSD measurement of the Co—Cr—Mo base alloy shown in FIG.
  • FIG. 3 is a (111), (110) and (100) pole figure of a phase.
  • 2 is a graph showing the distribution of crystal orientation difference obtained by EBSD measurement of the Co—Cr—Mo based alloy shown in FIG. The graph showing the deviation from the ideal orientation difference of the ⁇ 3 twin interface obtained by EBSD measurement of the Co—Cr—Mo base alloy shown in FIG. 1, and (inset) the ratio of the ⁇ 3 twin interface to the rolling reduction It is a graph which shows a change.
  • FIG. 3 shows the XRD line profile analysis results of 331, 420 and 422 diffractions of the ⁇ phase after (a) hot rolling and (b) 92.8% hot rolling of the Co—Cr—Mo based alloy shown in FIG. It is a graph. It is an optical microscope photograph which shows the structure of (a) CCM alloy and (b) CCMN alloy before hot forging of the Co—Cr—Mo base alloy of the second example of the present invention. 9 is an X-ray diffraction pattern before and after hot forging of the (a) CCM alloy and (b) CCMN alloy of the Co—Cr—Mo based alloy shown in FIG. FIG.
  • FIG. 9 is an image quality (IQ) map after hot forging obtained by EBSD measurement of the Co—Cr—Mo base alloy shown in FIG. 9 is a graph showing the distribution of crystal orientation difference before and after hot forging obtained by EBSD measurement of the Co—Cr—Mo base alloy shown in FIG. 9 is a (111), (110) and (100) pole figure obtained by EBSD data analysis of a CCMN alloy subjected to hot forging at a reduction rate of 83% of the Co—Cr—Mo based alloy shown in FIG. .
  • FIG. 9 is a graph of a nominal stress-nominal strain curve obtained by a tensile test of the (a) CCM alloy and (b) CCMN alloy before and after hot forging of the Co—Cr—Mo based alloy shown in FIG. 8. It is Image quality (IQ) map after hot forging obtained by EBSD measurement of the Co—Cr—Mo base alloy of the third example of the present invention.
  • FIG. 15 is a graph of a nominal stress-nominal strain curve obtained by a tensile test of the Co—Cr—Mo based alloy shown in FIG. 14.
  • initial structure This hot forged material (initial structure) was cut into a size of 12.5 mm in thickness, 80 mm in width, and 60 mm in length using a wire electric discharge machine, and used as a sample for hot rolling.
  • the heating temperature is 1200 ° C.
  • the rolling rate is 0.2 to 0.7 mm per pass (the rolling rate is less than 20%)
  • the plate thickness is 5.0 mm (the rolling rate is 60.0%)
  • heating was repeated at 1200 ° C. for 180 seconds for each pass.
  • the obtained hot rolled sheet was subjected to structure observation, EBSD measurement, X-ray diffraction, and tensile test.
  • Microscopic observation was performed by a backscattered electron diffraction (EBSD) method using an optical microscope and a field emission scanning electron microscope (FESEM: “XL30S-FEG” manufactured by Philips) on the rolling surface of the test piece.
  • the EBSD measurement was performed at an acceleration voltage of 15 kV.
  • the average crystal grain size was calculated by the intercept method, excluding the annealing twins contained in the crystal grains and the grain boundaries resulting from them.
  • the constituent phases were identified by X-ray diffraction (XRD: “X'Pert MPD” manufactured by Philips).
  • the tensile test piece had a length between ratings of 11.5 mm, a width of 1.6 mm, and a thickness of 1 mm, and an Instron type tensile tester was used at an initial strain rate of 1.45 ⁇ 10 ⁇ 4 s ⁇ 1 .
  • the tensile load axis was taken from the test piece so as to be parallel to the rolling direction (hereinafter referred to as RD) and perpendicular to the rolling direction (hereinafter referred to as TD).
  • FIG. 1 shows optical microscope structures before and after hot rolling.
  • the hot forged material initial structure
  • FIGS. 1B and 1C the same contrast was confirmed in the structure after hot rolling.
  • Table 1 shows the average crystal grain size of the test pieces before and after hot rolling.
  • the average crystal grain size in the hot forged material (initial structure) is 80.6 ⁇ m, but the crystal grains are refined to 50.7 ⁇ m after 92.8% hot rolling.
  • the structures before and after hot rolling were all equiaxed structures.
  • FIG. 2 shows the XRD measurement results of the hot forged material (initial structure) and the hot rolled material.
  • the diffraction pattern of the hot forging material (initial structure) showed a slight ⁇ phase, but it was confirmed that the ⁇ phase was stabilized even at room temperature by addition of nitrogen.
  • FIGS. 2 (b) and (c) there is almost no change in the constituent phase due to hot rolling, and the structures after hot rolling of 60.0% and 92.8% are almost ⁇ single phase ( ⁇ phase is 95% or more).
  • ⁇ phase and Cr 2 N and the like precipitation is expected diffraction peaks derived from these have not been confirmed. Therefore, hot rolling is performed in the ⁇ single phase region, and it is determined that precipitation of ⁇ phase did not occur during air cooling.
  • FIG. 3 shows a nominal stress-nominal strain curve obtained by a tensile test of the hot forged material (initial structure) and the hot rolled material.
  • Table 2 shows the tensile properties of each test piece. As shown in FIG. 3 and Table 2, tensile strength sufficient as a biomaterial was obtained in any of the test pieces, but the strength increased remarkably as the hot rolling rate increased, and 92.8% hot. After rolling, the 0.2% yield strength was increased to 2 times or more (1089 MPa) before rolling (509 MPa). On the other hand, although the elongation slightly decreases with an increase in the hot rolling rate, both are about 20%, and it has been confirmed that an excellent strength-ductility balance can be obtained by hot rolling. In addition, there is no significant difference in tensile properties for both RD and TD.
  • FIG. 4 shows test pieces determined by EBSD measurement (a) hot forged material (initial structure), (b) after 60.0% hot rolling, and (c) after 92.8% hot rolling.
  • the (111), (110) and (100) pole figures of the ⁇ phase are respectively shown.
  • the EBSD measurement result is local information, it has been confirmed that the incomplete pole figure obtained by the XRD measurement shows a similar texture.
  • the hot forged material (initial structure) subjected to hot rolling has a relatively random crystal orientation distribution, whereas, as shown in FIG. In the 0% hot-rolled material, a mixed texture of Brass type ⁇ 110 ⁇ ⁇ 112> and Goss type ⁇ 110 ⁇ ⁇ 001> was confirmed.
  • Such a texture is the same as the cold rolling texture of 304 series austenitic stainless steel, which is a representative alloy with low stacking fault energy, and despite hot rolling at 1200 ° C, It can be said that such a texture is very characteristic. As shown in FIG. 4 (c), after 92.8% rolling, although the crystal rotation was slightly observed around the TD axis, the same texture as after 60.0% hot rolling was confirmed.
  • FIG. 5 shows the distribution of misorientation angle of each test piece obtained by EBSD measurement.
  • the proportion of small-angle grain boundaries with a crystal orientation difference of 15 ° or less is low, and the presence of ⁇ 3-compatible grain boundaries, that is, annealing twins. It was confirmed that the distribution around 60 ° showing Moreover, it was confirmed that with the increase in the hot rolling rate, the ratio of the ⁇ 3 twin interface greatly decreased, and instead, the small-angle grain boundaries significantly increased.
  • FIG. 6 shows the result of analysis by EBSD, focusing on the ⁇ 3 twin interface.
  • the dislocation density of this hot-rolled material was measured by XRD line profile analysis using the method by Garrod et al. (R.I.Garrod, J.H. Auld: Acta Metall. 3 (1955) 190-198).
  • the XRD line profile in addition to the influence of microcrystal distortion and diffraction that reflect mainly the dislocation density and the coherent single crystal domain, that is, the size of the crystallite, the broadening of the peak due to the optical system occurs. Therefore, from the measurement profile, the influence of the optical system was corrected using the method proposed by Stokes (A.R. Stokes: Proc. Phys. Soc. 61 (1948) ⁇ 382-391), and the sample-derived physical profile was extracted.
  • the cosine coefficient part (real part) of the normalized Fourier coefficient of h (s), g (s) and f (s) is H (n), G (n) and F ( When expressed as n), the following relationship is obtained.
  • n is an integer
  • the measurement profile is equally divided into X
  • Equations (2) and (3) are obtained.
  • Equation (4) is established in a range where L is small.
  • D is the average crystallite size
  • ⁇ 2 > is the mean square of strain
  • a is the lattice constant
  • h 0 2 h 2 + k 2 + l 2 (h, k, l: plane of diffraction peak used for analysis Index). Therefore, D and ⁇ 2 > are obtained from the intercept and slope of the plot of ⁇ lnF (L) / L. Further, the average dislocation density ⁇ can be obtained from the equation (5) using ⁇ 2 >.
  • b is the magnitude of the dislocation Burgers vector.
  • Table 3 summarizes the dislocation density and crystallite size of the samples before and after hot rolling obtained from the analysis. From the analysis results of each diffraction peak, the average value of the dislocation density is 1.1 ⁇ 10 15 m ⁇ 2 in the hot forged material (initial structure), but after 92.8% hot rolling, the average value is 4 .6 ⁇ 10 15 m -2, and the possible dislocation density by hot rolling is increased to about 4 times was observed. It was also confirmed that the crystallite size was reduced to about 1 ⁇ 2 by hot rolling. Therefore, it became clear that the remarkable increase in strength by hot rolling is due to the introduction of a high dislocation density.
  • CCM alloy an alloy containing almost no N (addition amount 0.005 mass%)
  • CCMN alloy an alloy added with 0.12 mass% N
  • a 30 kg ingot of both alloys was melted in a high frequency induction melting furnace. This ingot was subjected to a homogenization heat treatment at 1225 ° C. for 12 hours, and then heated to 1200 ° C. to produce a round bar having a diameter of 14 mm by swaging so that the rolling reduction at one time was less than 20%.
  • a cylindrical sample having a diameter of 14 mm and a height of 28 mm was cut out from the round bar.
  • the columnar sample was heated at 1200 ° C. for 5 minutes, and then subjected to hot upset forging in the atmosphere with a hydraulic servo press. Since the ⁇ ⁇ ⁇ phase transformation adversely affects hot forgeability, it is important to maintain the sample temperature at 900 ° C. or higher at which the ⁇ phase can exist stably.
  • the previous research results show that fine crystal grains can be obtained by relatively high-speed processing (strain rate 0.1 to 1.0 s ⁇ 1 ). 30 mms ⁇ 1 (initial strain rate was 1.1 s ⁇ 1 ).
  • Hot forging is a reduction ratio of 61% for CCM alloys (height after forging 11.0 mm), 78% (6.10 mm) and 83% (4.66 mm), and for CCMN alloys, The rolling reduction was 83% (4.71 mm). All samples were water quenched after high temperature forging to avoid precipitation of ⁇ phase.
  • FIG. 8A and FIG. 8B show the structures of round bars of CCM alloy and CCMN alloy before hot upset forging, respectively (hereinafter referred to as “initial structure”).
  • both alloys have an equiaxed structure, and a linear structure is observed in the crystal grains.
  • these correspond to the non-thermal ⁇ -martensite phase and annealing twins introduced by water quenching after hot working.
  • most of CCMN alloys are annealed twins.
  • the crystal grain size calculated by excluding ⁇ -martensite and annealing twins of the round bar was about 100 ⁇ m in any composition.
  • FIG. 9 shows the X-ray diffraction results of the initial structure and the hot forged material.
  • the diffraction peak of the ⁇ phase is observed in the initial structure of the CCM alloy. That is, it can be seen that the CCM alloy has a two-phase structure of a ⁇ phase and an ⁇ phase, and the proportion of the ⁇ phase is clearly below 90%.
  • the hot forged material of the CCM alloy has less ⁇ phase compared to the initial structure. This shows that non-thermal martensitic transformation is suppressed by crystal grain refinement.
  • FIG. 9 (a) shows that in the initial structure of the CCM alloy in addition to the diffraction peak of the ⁇ phase. That is, it can be seen that the CCM alloy has a two-phase structure of a ⁇ phase and an ⁇ phase, and the proportion of the ⁇ phase is clearly below 90%.
  • the hot forged material of the CCM alloy has less ⁇ phase compared to the initial structure. This shows that non-thermal martensi
  • the hot forged material of CCMN alloy shows almost no ⁇ phase peak in the XRD pattern, and has a substantially ⁇ single phase structure ( ⁇ phase is 95% or more). It was confirmed that there was. This shows that the ⁇ phase is stabilized even with a small amount of N added.
  • FIG. 10 shows the image quality (IQ) map obtained by the EBSD measurement of the hot forging material of each alloy.
  • the observation surface is a sample cross section parallel to the forging direction.
  • IQ image quality
  • FIG. 11 shows the distribution of misorientation angle of each test piece obtained by EBSD measurement.
  • the initial structure of the CCMN alloy has a low proportion of small-angle grain boundaries with a crystal orientation difference of 15 ° or less (5% or less), and a distribution near 60 ° is 50 to 60%. It can be seen that it contains crystals.
  • the ratio of small-angle grain boundaries increases by hot forging, but even in the CCM alloy subjected to hot forging with a reduction ratio of 83%, which is the highest ratio of small-angle grain boundaries, It is 15% or less of the grain boundary, and is greatly different from the structure obtained by hot rolling shown in Example 1.
  • FIG. 12 shows a pole figure created by analyzing EBSD data of a CCMN alloy subjected to hot forging with a rolling reduction of 83%. As shown in FIG. 12, it was confirmed by hot forging that a fiber texture was formed in which the forging direction and the ⁇ 110> direction of the ⁇ phase were parallel.
  • FIG. 13 shows a nominal stress-nominal strain curve when a room temperature tensile test is performed on samples before and after hot forging of (a) CCM alloy and (b) CCMN alloy.
  • the obtained tensile properties are shown in Table 5 together with the average crystal grain size.
  • the 0.2% proof stress is remarkably improved as the hot forging rate increases, but the tensile ductility increases the working rate. It was confirmed that the value decreased to 2.5% after 83% hot forging.
  • the test piece after 83% hot forging shows a high 0.2% proof stress equivalent to the CCM alloy subjected to hot forging at the same processing rate.
  • a large elongation exceeding 20% was exhibited.
  • This remarkable increase in strength is caused by dislocation strengthening due to the introduction of dislocations at a high density as the crystal grains become finer.
  • the reason why the ductility is not impaired in the CCMN alloy is that the crystal is controlled to a ⁇ single phase structure. Therefore, it was confirmed that the CCMN alloy has both high strength and high ductility.
  • Example 2 As a test material, a CCMN alloy subjected to hot forging with a reduction rate of 83% in Example 2 was used. Using this test material, heat treatment was performed at 900 to 1200 ° C. and a holding time of 1 to 5 minutes using an infrared cold image furnace (manufactured by ULVAC-RIKO). In addition, it hardened in ice water after heat processing.
  • FIG. 14 shows the structure of each test piece after the heat treatment. As shown in FIG. 14, it was confirmed that in any case, remarkable crystal grain growth did not occur as compared with the hot forged test piece, and a fine structure of about several ⁇ m or less was maintained.
  • FIG. 15 shows a nominal stress-nominal strain curve of each test piece obtained by the tensile test. For comparison, the results of the test pieces before hot forging and as hot forged are also shown. As shown in FIG. 15, it was confirmed that the specimens subjected to heat treatment at 1100 ° C. and 1200 ° C. for 1 minute had a significantly lower elongation than the tensile properties as hot forged, but the elongation was significantly increased. It was. Further, it was confirmed that the test piece subjected to the heat treatment after hot forging showed a large elongation even when the 0.2% proof stress was comparable, compared with the initial structure before hot forging. This is because the heat treatment material has finely adjusted crystal grains compared to the initial structure.

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  • Crystallography & Structural Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Chemical Kinetics & Catalysis (AREA)
  • Electrochemistry (AREA)
  • Forging (AREA)
  • Heat Treatment Of Steel (AREA)

Abstract

L'invention concerne un alliage à base de Co-Cr-Mo qui peut être utilisé comme biomatériau, est très sûr et présente une résistance élevée, une ductilité élevée et une résistance à la fatigue élevée, ledit alliage étant produit par la soumission d'un matériau brut à un formage à chaud de manière à ce que le taux total de réduction par laminage à 1000˚C ou plus atteigne 60 % ou plus, le matériau brut comprenant 25 à 35 % en masse de Cr, 3 à 8 % en masse de Mo et des impuretés contenant au moins un élément sélectionné dans le groupe constitué par 0 à 0,1 % en masse de Zr, 1 % en masse ou moins de Si, 1 % en masse ou moins de Mn, 0,01 à 0,3 % en masse de C et 0,08 à 0,8 % en masse de N, le reste étant du Co; et un procédé de production de l'alliage.
PCT/JP2012/053053 2012-02-10 2012-02-10 ALLIAGE À BASE DE Co-Cr-Mo ET PROCÉDÉ DE PRODUCTION D'ALLIAGE À BASE DE Co-Cr-Mo WO2013118285A1 (fr)

Priority Applications (2)

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JP2013557324A JP5846530B2 (ja) 2012-02-10 2012-02-10 Ｃｏ−Ｃｒ−Ｍｏ基合金およびＣｏ−Ｃｒ−Ｍｏ基合金の製造方法
PCT/JP2012/053053 WO2013118285A1 (fr) 2012-02-10 2012-02-10 ALLIAGE À BASE DE Co-Cr-Mo ET PROCÉDÉ DE PRODUCTION D'ALLIAGE À BASE DE Co-Cr-Mo

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PCT/JP2012/053053 WO2013118285A1 (fr) 2012-02-10 2012-02-10 ALLIAGE À BASE DE Co-Cr-Mo ET PROCÉDÉ DE PRODUCTION D'ALLIAGE À BASE DE Co-Cr-Mo

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Citations (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS5410224A (en) * 1977-06-23 1979-01-25 Howmedica Nitrogen containing cobalt cromium molibuden alloy
JP2008111177A (ja) * 2006-10-31 2008-05-15 Iwate Univ 塑性加工性に優れる生体用Ｃｏ基合金およびその製造方法
JP2009046760A (ja) * 2007-07-24 2009-03-05 Kobe Steel Ltd 熱間型鍛造用の生体用Ｃｏ基合金素材及びその製造方法
JP2009046758A (ja) * 2007-07-24 2009-03-05 Kobe Steel Ltd 生体用Ｃｏ基合金及びその製造方法
WO2010026996A1 (fr) * 2008-09-05 2010-03-11 国立大学法人東北大学 PROCÉDÉ DE FORMATION DE FINS GRAINS CRISTALLINS DANS UN ALLIAGE DE Co-Cr-Mo DOPÉ À L'AZOTE ET ALLIAGE DE Co-Cr-Mo DOPÉ À L'AZOTE
JP2010215960A (ja) * 2009-03-16 2010-09-30 Iwate Univ 機械部品の製造方法及び機械部品

Family Cites Families (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2011184783A (ja) * 2010-03-11 2011-09-22 Tohoku Univ 窒素添加Ｃｏ−Ｃｒ−Ｍｏ合金の結晶粒微細化方法

Patent Citations (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS5410224A (en) * 1977-06-23 1979-01-25 Howmedica Nitrogen containing cobalt cromium molibuden alloy
JP2008111177A (ja) * 2006-10-31 2008-05-15 Iwate Univ 塑性加工性に優れる生体用Ｃｏ基合金およびその製造方法
JP2009046760A (ja) * 2007-07-24 2009-03-05 Kobe Steel Ltd 熱間型鍛造用の生体用Ｃｏ基合金素材及びその製造方法
JP2009046758A (ja) * 2007-07-24 2009-03-05 Kobe Steel Ltd 生体用Ｃｏ基合金及びその製造方法
WO2010026996A1 (fr) * 2008-09-05 2010-03-11 国立大学法人東北大学 PROCÉDÉ DE FORMATION DE FINS GRAINS CRISTALLINS DANS UN ALLIAGE DE Co-Cr-Mo DOPÉ À L'AZOTE ET ALLIAGE DE Co-Cr-Mo DOPÉ À L'AZOTE
JP2010215960A (ja) * 2009-03-16 2010-09-30 Iwate Univ 機械部品の製造方法及び機械部品

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