WO2021220754A1 - Stainless steel sheet, method for producing same, edged tools and cutlery - Google Patents

Stainless steel sheet, method for producing same, edged tools and cutlery Download PDF

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WO2021220754A1
WO2021220754A1 PCT/JP2021/014829 JP2021014829W WO2021220754A1 WO 2021220754 A1 WO2021220754 A1 WO 2021220754A1 JP 2021014829 W JP2021014829 W JP 2021014829W WO 2021220754 A1 WO2021220754 A1 WO 2021220754A1
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Prior art keywords
steel sheet
rolled
hot
less
cold
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PCT/JP2021/014829
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French (fr)
Japanese (ja)
Inventor
正崇 吉野
詩乃 廣田
卓也 松本
彩子 田
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Ｊｆｅスチール株式会社
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Priority to EP21797895.6A priority Critical patent/EP4144882A1/en
Priority to CN202180030734.2A priority patent/CN115461481B/en
Priority to KR1020227034500A priority patent/KR20220147135A/en
Priority to JP2021540467A priority patent/JP7226564B2/en
Priority to MX2022013630A priority patent/MX2022013630A/en
Priority to US17/995,582 priority patent/US20230212705A1/en
Publication of WO2021220754A1 publication Critical patent/WO2021220754A1/en

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    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
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    • C21D1/18—Hardening; Quenching with or without subsequent tempering
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    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0273—Final recrystallisation annealing
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    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/18—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for knives, scythes, scissors, or like hand cutting tools
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    • 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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
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    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
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    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
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    • C—CHEMISTRY; METALLURGY
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    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel with niobium or tantalum
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Definitions

  • the present invention is a stainless steel having high hardness and good surface quality suitable for use in knives such as kitchen knives and scissors, medical scalpels, cutlery such as table knives and forks, spoons, and precision tools such as tweezers. It is about steel plates.
  • Stainless steel plates may be used as materials for knives, scissors, scalpels and other cutting tools, and precision tools such as tweezers.
  • a stainless steel plate is blanked or forged into a predetermined shape by press working or the like.
  • the stainless steel sheet processed into a predetermined shape is subjected to quenching treatment, quenching treatment and tempering treatment to harden it.
  • the hardened stainless steel sheet is subjected to blade polishing (a process of thinning the portion to be the blade edge by polishing) or the like to obtain a final product (kitchen knife).
  • Examples of stainless steel used for such high-hardness high-grade blades include 14 mass% Cr-0.5 mass% C steel conforming to the European unified standard: EN1.4116.
  • the 14 mass% Cr-0.5 mass% C steel conforming to this European unified standard: EN1.4116 is a steel in which the amount of C is increased and the hardness is increased as compared with the 13 mass% Cr-0.3 mass% C steel.
  • Patent Document 1 "C: 0.88 mass% or more and 1.2 mass% or less, Cr: 12.5 mass% or more and 16.50 mass% or less, Si: 0.05 mass% or more and 0.20 mass% or less, N: 0.001 mass% or more and 0.02 mass” % Or less, Mn: 1.0 mass% or less, Cu: 1.0 mass% or less, P: 0.03 mass% or less, S: 0.010 mass% or less, and Ni: 1.0 mass% or less, and the balance is A stainless steel strip characterized by consisting of Fe and unavoidable impurities. " Is disclosed.
  • the present invention has been developed in view of the above-mentioned current situation, has high hardness when used as a product (hereinafter, also simply referred to as high hardness), and generates streaks when processed into a product. It is an object of the present invention to provide a stainless steel sheet having a good surface quality (hereinafter, also simply referred to as a good surface quality) in which the above is suppressed. Another object of the present invention is to provide the above-mentioned method for manufacturing a stainless steel sheet. Furthermore, an object of the present invention is to provide a cutlery and a cutlery made of the above-mentioned stainless steel plate.
  • the stainless steel sheet of the present invention is intended for those having high hardness when used as products such as cutlery and cutlery. That is, the stainless steel sheet of the present invention includes not only a steel sheet after being hardened (after quenching) but also a steel sheet as a product material before being hardened (before quenching).
  • the inventors came to think as follows. That is, even if the steel sheet a and the steel sheet b are manufactured under the same manufacturing conditions due to the difference in the component composition, the precipitation state of the precipitates is significantly different. Then, due to the difference in the precipitation state of the precipitate, a streak pattern is generated on the steel plate a.
  • the inventors observed the metallographic structures of the steel plate a and the steel plate b and compared them in detail.
  • coarse Cr-based carbides are continuously present in the metal structure in the rolling direction, which causes the streaks to be generated.
  • the Cr-based carbide is harder than the base material of the stainless steel sheet (both before and after quenching). Therefore, when coarse Cr-based carbides are present in the metal structure, the amount of polishing is smaller in the portion where the Cr-based carbides are present than in other portions. As a result, after polishing, convex portions are locally generated, and these are manifested as streaks.
  • the steel sheet b in order to obtain higher hardness, the steel sheet b (13 mass% Cr-0.3 mass%) is used. It contains a large amount of C and Cr as compared with C steel). Therefore, the steel sheet b does not generate a large amount of coarse Cr-based carbides even if it is manufactured by a conventionally known method, but the steel sheet a manufactured under the same conditions produces a large amount of coarse Cr-based carbides and has a streak pattern. appear.
  • the inventors further studied and obtained the following findings. That is, the above-mentioned coarse Cr-based carbide is produced along the casting direction in the vicinity of the boundary between the columnar crystal and the equiaxed crystal in the slab cross section at the time of casting. Further, the coarse Cr-based carbon dioxide produced during casting is subjected to hot rolling, hot rolling plate annealing, cold rolling and cold rolling plate annealing steps after the casting step under general manufacturing conditions known conventionally. Still remains in the rolling direction (the same direction as the casting direction).
  • the inventors have repeatedly studied a method for preventing the formation of coarse Cr-based carbides while obtaining high hardness.
  • the following findings were obtained.
  • (1) Appropriately adjust the component composition, in particular, adjust the C content and Cr content to 0.45 to 0.60% by mass and 13.0% or more and less than 16.0%, respectively.
  • (2) Then, the heating, hot rolling and hot rolling sheet annealing conditions of the steel slab are appropriately controlled. In particular, (A) Hold the steel slab at 1200 to 1350 ° C. for 30 minutes or more, and (B) Of the rolling passes in hot rolling, the number of rolling passes having an end temperature of 1050 ° C. or higher and a rolling reduction ratio of 20% or higher shall be 3 or more.
  • the winding temperature of the hot-rolled steel sheet is set to 600 ° C. or higher. This is very important. As a result, even when the C content and the Cr content are contained in a certain amount or more, the formation of coarse Cr-based carbides can be suppressed, and the generation of streaks during polishing can be effectively prevented.
  • the inventors consider the reason why the formation of coarse Cr-based carbides is suppressed by controlling the production conditions as described above. That is, as described in (2) and (a) above, by holding the steel slab at 1200 to 1350 ° C. for 30 minutes or more, the coarse Cr-based carbide produced in the casting step is solid-solved in the austenite phase (Cr-based carbide). Is decomposed into Cr atoms, C atoms, etc. and incorporated into the austenite phase in the atomic state) is promoted. Further, in this state, as described in (2) and (b) above, by performing the rolling pass in hot rolling at a high temperature and a high rolling reduction, the solid solution of the Cr-based carbide in the austenite phase is further increased. It is promoted.
  • the gist structure of the present invention is as follows. 1.
  • C 0.45 to 0.60%
  • Si 0.05 to 1.00%
  • Mn 0.05 to 1.00%
  • P 0.05% or less
  • S 0.020% or less
  • Cr 13.0% or more and less than 16.0%
  • Ni 0.10 to 1.00%
  • N 0.010 to 0.200%
  • Particle size A stainless steel sheet having a total volume fraction of Cr-based carbides of 2.0 ⁇ m or more of 10% or less.
  • the component composition is further increased by mass%.
  • Mo 0.05-1.00%
  • Cu 0.05 to 1.00%
  • Co 0.05 to 0.50%
  • the component composition is further increased by mass%.
  • Al 0.001 to 0.100%
  • Ti 0.01 to 0.10%
  • Nb 0.01 to 0.10%
  • V 0.05 to 0.50%
  • Zr 0.01-0.10%
  • Mg 0.0002 to 0.0050%
  • B 0.0002 to 0.0050%
  • Ca 0.0003 to 0.0030%
  • REM 0.01 to 0.10%
  • a method for manufacturing the stainless steel sheet according to any one of 1 to 3 above The first step of holding the steel slab having the component composition according to any one of 1 to 3 at 1200 to 1350 ° C. for 30 minutes or more.
  • a third step of subjecting the hot-rolled steel sheet to hot-rolled annealed steel sheet to obtain a hot-rolled and annealed steel sheet is provided.
  • the rolling passes in the hot rolling of the second step the number of rolling passes having an end temperature of 1050 ° C. or higher and a reduction rate of 20% or higher is 3 or more, and the hot-rolled steel sheet is wound.
  • the taking temperature is 600 ° C or higher
  • the holding temperature in the hot-rolled sheet annealing of the third step is 750 to 900 ° C., and the holding time is 10 minutes or more. Manufacturing method of stainless steel plate.
  • the cold-rolled steel sheet is subjected to cold-rolled sheet annealing to obtain a cold-rolled annealed steel sheet, which is provided with a fifth step.
  • a sixth step of quenching the hot-rolled annealed steel sheet, the cold-rolled steel sheet, or the cold-rolled annealed steel sheet is provided.
  • a seventh step of tempering the hardened steel sheet is provided.
  • the present invention will be described based on the following embodiments. First, the component composition of the stainless steel sheet according to the embodiment of the present invention will be described.
  • the unit in the component composition is "mass%”, but hereinafter, unless otherwise specified, it is simply indicated by "%".
  • C 0.45 to 0.60%
  • C has the effect of hardening the martensite phase obtained by the quenching treatment.
  • the C content is less than 0.45%, the hardness after the quenching treatment is insufficient, and the sharpness required for a high-grade blade cannot be sufficiently obtained.
  • the C content exceeds 0.60%, even if the production conditions are appropriately controlled, the generation of coarse carbides cannot be sufficiently suppressed, and good surface quality cannot be obtained.
  • quench cracking is likely to occur during the quenching process, making it difficult to stably manufacture the cutting tool. Therefore, the C content is set in the range of 0.45 to 0.60%.
  • the C content is preferably 0.55% or less, more preferably 0.50% or less.
  • Si acts as an antacid during the melting of steel.
  • the Si content is set to 0.05% or more.
  • the Si content is in the range of 0.05 to 1.00%.
  • the Si content is preferably 0.20% or more.
  • the Si content is preferably 0.60% or less.
  • Mn 0.05 to 1.00% Mn has the effect of promoting the formation of the austenite phase and improving the hardenability. In order to obtain such an effect, the Mn content is set to 0.05% or more. However, if the Mn content exceeds 1.00%, the corrosion resistance is lowered. Therefore, the Mn content is set in the range of 0.05 to 1.00%. The Mn content is preferably 0.40% or more. The Mn content is preferably 0.80% or less.
  • P 0.05% or less
  • P is an element that promotes grain boundary fracture due to grain boundary segregation. Therefore, it is desirable to reduce P as much as possible. Therefore, the P content is set to 0.05% or less.
  • the P content is preferably 0.04% or less, more preferably 0.03% or less.
  • the lower limit of the P content is not particularly limited. However, since excessive de-P causes an increase in cost, the P content is preferably 0.005% or more.
  • S 0.020% or less
  • S is an element that exists in steel as a sulfide-based inclusion such as MnS and lowers ductility, corrosion resistance, and the like. Therefore, it is desirable to reduce S as much as possible. Therefore, the S content is set to 0.020% or less.
  • the S content is preferably 0.015% or less.
  • the lower limit of the S content is not particularly limited. However, since excessive de-S causes an increase in cost, the S content is preferably 0.0005% or more.
  • the Cr content is set to 13.0% or more.
  • the Cr content is 16.0% or more, the amount of austenite produced during heating and holding of the quenching treatment decreases. Therefore, the martensite phase obtained after the quenching treatment is reduced, and sufficient hardness cannot be obtained. Therefore, the Cr content is in the range of 13.0% or more and less than 16.0%.
  • the Cr content is preferably 14.0% or more.
  • the Cr content is preferably 15.5% or less, more preferably 15.0% or less.
  • Ni 0.10 to 1.00%
  • Ni has the effect of improving corrosion resistance and toughness after quenching.
  • the Ni content is set to 0.10% or more.
  • the Ni content is saturated.
  • the steel sheet becomes excessively hard before the quenching treatment, and it becomes difficult to obtain sufficient workability when forming a predetermined shape of a cutting tool or the like. Therefore, the Ni content is set in the range of 0.10 to 1.00%.
  • the Ni content is preferably 0.15% or more, more preferably 0.20% or more.
  • the Ni content is preferably 0.80% or less, more preferably 0.60% or less.
  • N 0.010 to 0.200%
  • N has the effect of hardening the martensite phase obtained by quenching.
  • N also has the effect of improving the corrosion resistance after the quenching treatment.
  • the N content is set to 0.010% or more.
  • the N content is set in the range of 0.010 to 0.200%.
  • the N content is preferably 0.015% or more, more preferably 0.020% or more.
  • the N content is preferably 0.150% or less, more preferably 0.100% or less.
  • the basic composition of the stainless steel sheet according to the embodiment of the present invention has been described above.
  • Mo 0.05 to 1.00%
  • Cu 0.05 to 1.00%
  • Co 0.05 to 0.50%
  • And / or Al 0.001 to 0.100%
  • Ti 0.01 to 0.10%
  • Nb 0.01 to 0.10%
  • V 0.05 to 0.50%
  • Zr 0.01 to Of 0.10%
  • Mg 0.0002 to 0.0050%
  • B 0.0002 to 0.0050%
  • Mo 0.05-1.00% Mo has the effect of improving corrosion resistance.
  • the Mo content is preferably 0.05% or more.
  • the Mo content is more preferably 0.10% or more, still more preferably 0.50% or more.
  • the Mo content is more preferably 0.80% or less, still more preferably 0.65% or less.
  • Cu 0.05-1.00%
  • the Cu content is preferably 0.05% or more. However, if the Cu content exceeds 1.00%, the corrosion resistance is lowered. Therefore, when Cu is contained, the Cu content is preferably in the range of 0.05 to 1.00%.
  • the Cu content is more preferably 0.10% or more.
  • the Cu content is more preferably 0.50% or less, still more preferably 0.20% or less.
  • Co has the effect of improving toughness.
  • the Co content is preferably 0.05% or more.
  • the Co content is more preferably 0.10% or more.
  • the Co content is more preferably 0.20% or less.
  • Al acts as an antacid in the same manner as Si.
  • the Al content is preferably 0.001% or more.
  • the Al content is preferably in the range of 0.001 to 0.100%.
  • the Al content is more preferably 0.050% or less, still more preferably 0.010% or less.
  • Ti 0.01 to 0.10% Like Cr, Ti has a high affinity for C and N and is an element that forms carbides in steel. In addition, Ti has the effect of improving tempering and softening resistance. Therefore, it is possible to improve the toughness while suppressing the softening at the time of tempering. In order to obtain such an effect, the Ti content is preferably 0.01% or more. However, when the Ti content exceeds 0.10%, the effect is saturated. Moreover, the toughness is rather lowered. Therefore, when Ti is contained, the Ti content is preferably in the range of 0.01 to 0.10%. The Ti content is more preferably 0.02% or more. The Ti content is more preferably 0.05% or less.
  • Nb 0.01 to 0.10%
  • Nb is an element that has a high affinity for C and N and forms carbides in steel.
  • Nb has the effect of improving tempering softening resistance. Therefore, it is possible to improve the toughness while suppressing the softening at the time of tempering.
  • the Nb content is preferably 0.01% or more.
  • the Nb content exceeds 0.10%, the effect is saturated.
  • the toughness may decrease due to the precipitation of intermetallic compounds. Therefore, when Nb is contained, the Nb content is preferably in the range of 0.01 to 0.10%.
  • the Nb content is more preferably 0.02% or more.
  • the Nb content is more preferably 0.05% or less.
  • V 0.05 to 0.50%
  • V has a high affinity for C and N and is an element that forms carbides in steel. Further, V has an effect of improving tempering softening resistance. Therefore, it is possible to improve the toughness while suppressing the softening at the time of tempering.
  • the V content is preferably 0.05% or more. However, when the V content exceeds 0.50%, the effect is saturated. In addition, the toughness may decrease due to the precipitation of intermetallic compounds. Therefore, when V is contained, the V content is preferably in the range of 0.05 to 0.50%.
  • the V content is more preferably 0.10% or more.
  • the V content is more preferably 0.30% or less, still more preferably 0.20% or less.
  • Zr 0.01-0.10%
  • Zr has a high affinity for C and N and is an element that forms carbides in steel.
  • Zr has the effect of improving tempering and softening resistance. Therefore, it is possible to improve the toughness while suppressing the softening at the time of tempering.
  • the Zr content is preferably 0.01% or more.
  • the Zr content exceeds 0.10%, the effect is saturated.
  • the toughness may decrease due to the precipitation of intermetallic compounds. Therefore, when Zr is contained, the Zr content is preferably in the range of 0.01 to 0.10%.
  • the Zr content is more preferably 0.02% or more.
  • the Zr content is more preferably 0.05% or less.
  • Mg 0.0002 to 0.0050%
  • Mg has the effect of improving the equiaxed crystal ratio of the slab and improving workability and toughness.
  • the Mg content is preferably 0.0002% or more.
  • the Mg content is preferably in the range of 0.0002 to 0.0050%.
  • the Mg content is more preferably 0.0010% or more.
  • the Mg content is more preferably 0.0020% or less.
  • B 0.0002 to 0.0050%
  • B has the effect of improving hot workability during casting and hot rolling. Further, B segregates at the grain boundaries of the ferrite phase and the austenite phase to increase the grain boundary strength. This suppresses the occurrence of cracks during casting and hot rolling.
  • the B content is preferably 0.0002% or more. However, if the B content exceeds 0.0050%, the workability when forming the steel sheet into a predetermined shape such as a cutting tool cannot be sufficiently obtained before the quenching treatment. It also causes a decrease in toughness. Therefore, when B is contained, the B content is preferably in the range of 0.0002 to 0.0050%. The B content is more preferably 0.0005% or more. The B content is more preferably 0.0030% or less, still more preferably 0.0020% or less.
  • Ca 0.0003 to 0.0030%
  • Ca has the effect of miniaturizing inclusions generated during smelting and continuous casting, and is particularly effective in preventing nozzle blockage during continuous casting.
  • the Ca content is preferably 0.0003% or more.
  • the Ca content is more preferably 0.0005% or more, still more preferably 0.0007% or more.
  • the Ca content is more preferably 0.0020% or less, still more preferably 0.0015% or less.
  • REM 0.01 to 0.10% REM (Rare Earth Metals) has the effect of improving hot ductility.
  • REM also has an effect of suppressing cracking and rough skin of the end face portion of the steel sheet during hot rolling.
  • the REM content is preferably 0.01% or more.
  • the REM content exceeds 0.10%, the effect is saturated.
  • REM is also an expensive element. Therefore, when REM is contained, the REM content is preferably in the range of 0.01 to 0.10%.
  • the REM content is more preferably 0.05% or less.
  • the rest of the components other than the above are Fe and unavoidable impurities.
  • the metal structure of the stainless steel sheet according to the embodiment of the present invention changes its main structure before and after the quenching treatment.
  • the steel sheet is blanked or forged into a predetermined shape by press working or the like at a stage where the steel sheet is not hardened.
  • the steel sheet processed into a predetermined shape is hardened by quenching or quenching and tempering. That is, before and after the quenching treatment, the main structure is changed, specifically, the ferrite phase is changed to the martensite phase.
  • Cr-based carbides having a particle size of 2.0 ⁇ m or more do not change much before and after the quenching treatment and are almost maintained. Therefore, in the metal structure of the stainless steel sheet according to the embodiment of the present invention, it is extremely important that the volume ratio of Cr-based carbide having a particle size of 2.0 ⁇ m or more is 10% or less regardless of before and after the quenching treatment. ..
  • volume fraction of Cr-based carbide 10% or less Cr-based carbide is harder than the base material of stainless steel sheet (both before and after quenching). Therefore, if polishing or cutting is performed in a state where a large amount of coarse Cr-based carbides, particularly Cr-based carbides having a particle size of 2.0 ⁇ m or more, are present in the metal structure, the Cr-based carbides are present. The amount of polishing is smaller in the area where the material is used than in other areas. As a result, after polishing, convex portions are locally generated, and these are manifested as streaks. Therefore, the volume fraction of Cr-based carbide having a particle size of 2.0 ⁇ m or more is set to 10% or less.
  • the volume ratio of Cr-based carbide having a particle size of 2.0 ⁇ m or more is preferably 5% or less, more preferably 2% or less.
  • the volume fraction of Cr-based carbide having a particle size of 2.0 ⁇ m or more may be 0%.
  • the volume fraction of Cr-based carbide having a particle size of less than 2.0 ⁇ m is not particularly limited.
  • the Cr-based carbide referred to here is mainly Cr 23 C 6 . Further, a part of Cr in the Cr carbide is replaced with an element such as Fe, Mn, Ti, Nb, V, Zr, and a part of C is replaced with N. It shall be contained in carbide.
  • the structure other than the Cr-based carbide in the stainless steel sheet according to the embodiment of the present invention is a metal structure in which the total volume ratio of the ferrite phase and the martensite phase is 95% or more, more preferably 98% or more.
  • the total volume fraction of the ferrite phase and the martensite phase may be 100%.
  • Remaining structures other than the ferrite phase, martensite phase and the above-mentioned Cr-based carbides include residual austenite phase and other precipitates (including Cr-based carbides having a particle size of less than 2.0 ⁇ m), inclusions (for example, Al and the like. Oxides such as Si and sulfides such as Mn) can be mentioned.
  • the volume fraction of the residual tissue is preferably 5% or less, more preferably 2% or less.
  • the volume fraction of the residual tissue may be 0%.
  • the stainless steel sheet according to the embodiment of the present invention includes both steel sheets before and after the quenching treatment, for example, a hot-rolled steel sheet, a hot-rolled tempered steel sheet, a cold-rolled steel sheet, a cold-rolled fired steel sheet, and these.
  • the steel sheet includes a steel sheet that has undergone a quenching treatment and / or a tempering treatment (a hardened steel sheet and a tempered steel sheet described later) and the like.
  • the structure other than the Cr-based carbide is mainly a ferrite phase structure.
  • the ferrite phase has a metal structure having a volume fraction of 80% or more, preferably 90% or more, more preferably 95% or more, and further preferably 98% or more.
  • the volume fraction of the ferrite phase may be 100%.
  • the volume fraction of the residual tissue is preferably 20% or less, more preferably 10% or less, still more preferably 5% or less, still more preferably 2% or less.
  • the volume fraction of the residual tissue may be 0%.
  • the hot-rolled steel sheet includes, in addition to the hot-rolled steel sheet, a steel sheet obtained by subjecting the hot-rolled steel sheet to an oxidation scale removal treatment such as pickling.
  • the steel sheet obtained by hot-rolling the hot-rolled steel sheet is further subjected to an oxidation scale removal treatment such as pickling.
  • the steel sheet obtained from the above is included.
  • the cold-rolled steel sheet includes, in addition to the cold-rolled steel sheet, a steel sheet obtained by subjecting the cold-rolled steel sheet to an oxidation scale removal treatment such as pickling.
  • the structure other than Cr-based carbides is mainly composed of martensite phase.
  • the martensite phase has a metal structure having a volume fraction of 80% or more, preferably 90% or more, more preferably 95% or more, and further preferably 98% or more.
  • the volume fraction of the martensite phase may be 100%.
  • the residual structure other than the martensite phase and the above-mentioned Cr-based carbides includes a ferrite phase, a retained austenite phase, other precipitates (including Cr-based carbides having a particle size of less than 2.0 ⁇ m), inclusions (for example, Al and the like. Oxides such as Si and sulfides such as Mn) can be mentioned.
  • the volume fraction of the residual tissue is preferably 20% or less, more preferably 10% or less, still more preferably 5% or less, still more preferably 2% or less.
  • the volume fraction of the residual tissue may be 0%. Since it is hardened by the quenching treatment, the rockwell hardness of the hardened steel sheet is HRC55 or more.
  • a steel plate obtained by tempering a hardened steel plate (hereinafter, also referred to as a tempered steel plate), martensite having a reduced dislocation density and solid-melt C and N in structures other than Cr-based carbides as compared with that after the quenching treatment.
  • the structure is mainly phase (sometimes called tempered martensite phase), and the martensite fraction before tempering is almost maintained.
  • the martensite phase has a metal structure having a volume fraction of 80% or more, preferably 90% or more, more preferably 95% or more, and further preferably 98% or more.
  • the ferrite phase has a metal structure having a volume fraction of 20% or less, preferably 10% or less, more preferably 5% or less, still more preferably 2% or less.
  • Remaining structures other than the ferrite phase, martensite phase and the above-mentioned Cr-based carbides include residual austenite phase and other precipitates (including Cr-based carbides having a particle size of less than 2.0 ⁇ m), inclusions (for example, Al and the like. Oxides such as Si and sulfides such as Mn) can be mentioned.
  • the volume fraction of the residual tissue is preferably 5% or less, more preferably 2% or less.
  • the tempering treatment is performed to adjust the hardness and durability of the steel sheet hardened by the quenching treatment, and the hardness of the tempered steel sheet is lower than that of the tempered steel sheet before the tempering treatment.
  • the Rockwell hardness is HRC40 to 50.
  • the volume fraction of Cr-based carbide having a particle size of 2.0 ⁇ m or more is measured as follows. That is, a test piece for observing the structure is collected from the central portion of the width of the steel plate used as the test material. Then, the cross section of the test piece in the rolling direction is mirror-polished and then etched with an aqueous solution of hydrochloric acid picrinate, and an optical micrograph with a magnification of 500 times is taken in 10 fields. The area of Cr-based carbides in the obtained microstructure photograph is measured by image analysis, and Cr-based carbides having a circle-equivalent diameter of 2.0 ⁇ m or more are identified.
  • the total area ratio of the Cr-based carbides having the specified circle-equivalent diameter of 2.0 ⁇ m or more is calculated, and the calculated value is used as the volume fraction of the Cr-based carbides having a particle size of 2.0 ⁇ m or more.
  • the boundary between the grain boundary of the matrix phase (ferrite phase or martensite phase) and the boundary of the precipitate is automatically detected by the contrast difference using the image analysis device for the digital data of the tissue photograph (grains).
  • the boundaries and boundaries have a linear black contrast, and the grains have a relatively bright contrast).
  • the region surrounded by the boundary line between the matrix and the precipitate is designated as a precipitate, and the area of each precipitate region is automatically measured.
  • the identification that the precipitate in the above-mentioned microstructure photograph is a Cr-based carbide is performed as follows. That is, point analysis using SEM-EDS (Scanning Electron Microscope-Energy Dispersive X-ray Spectroscopy) is performed in the same field of view in which the above-mentioned tissue photograph is taken, and the main component of the observed precipitate is measured. Specifically, the total content of Cr and Fe in the precipitate is 60% by mass or more, and the ratio of the Cr content in the precipitate to the total content of Fe and Cr in the precipitate ([Cr. When the content (% by mass)] / ([Fe content (% by mass)] + [Cr content (% by mass)]) is 0.4 or more, the precipitate is identified as a Cr-based carbide.
  • SEM-EDS Sccanning Electron Microscope-Energy Dispersive X-ray Spectroscopy
  • the volume fractions of the ferrite phase and the martensite phase are determined as follows. That is, in the above microstructure photograph, the martensite phase and the ferrite phase are distinguished from each other based on the structure shape and the etching strength (note that the martensite phase is etched deeper than the ferrite phase. Therefore, the martensite phase is more than the ferrite phase. The contrast is dark.). Then, the volume fractions of the ferrite phase and the martensite phase are calculated for each field of view by image processing. Then, the arithmetic mean value of the volume ratios of the ferrite phase and the martensite phase obtained for each field of view is calculated, and the value is used as the volume ratio of the ferrite phase and the martensite phase.
  • the thickness of the stainless steel plate according to the embodiment of the present invention is not particularly limited, but is preferably 0.1 to 5.0 mm from the viewpoint of application to kitchen knives, razors, medical blades, and the like. Is.
  • the thickness of the stainless steel sheet according to the embodiment of the present invention is more preferably 0.5 mm or more, still more preferably 1.0 mm or more.
  • the thickness of the stainless steel sheet according to the embodiment of the present invention is more preferably 4.0 mm or less, still more preferably 2.5 mm or less.
  • molten steel is melted in a melting furnace such as a converter or an electric furnace. Then, the molten steel is subjected to secondary refining by ladle refining or vacuum refining to adjust to the above-mentioned composition. Then, the molten steel is made into a steel material (steel slab) by a continuous casting method, an ingot-bulk rolling method, or the like.
  • the slab heating temperature is less than 1200 ° C., the solid solution of Cr-based carbides in the austenite phase is not sufficiently promoted. Therefore, the formation of coarse Cr-based carbides is not sufficiently suppressed, and good surface quality cannot be obtained.
  • the slab heating temperature exceeds 1350 ° C.
  • the metal structure of the steel slab becomes a two-phase structure of an austenite phase and a delta ferrite phase or a single phase structure of delta ferrite, and the Cr-based carbide solid solution to the austenite phase. Melting is not sufficiently promoted. Therefore, the formation of coarse Cr-based carbides is not sufficiently suppressed, and good surface quality cannot be obtained. Therefore, the slab heating temperature is in the range of 1200 to 1350 ° C.
  • the slab heating temperature is preferably 1300 ° C. or lower, more preferably 1250 ° C. or lower.
  • the holding time at 1200 to 1350 ° C. is set to 30 minutes or more. If the holding time exceeds 24 hours, the oxide scale generated during heating of the steel slab becomes thick, and surface defects are likely to occur. In addition, productivity is also reduced. Therefore, the holding time is preferably 24 hours or less. The holding time is more preferably 12 hours or less, still more preferably 3 hours or less.
  • Hot rolling step Then, as a second step, the steel slab is hot-rolled to obtain a hot-rolled steel sheet, and the hot-rolled steel sheet is wound up. At this time, among the rolling passes in hot rolling, the number of rolling passes having an end temperature of 1050 ° C. or higher and a reduction rate of 20% or higher is 3 or more, and the winding temperature of the hot-rolled steel sheet is 600. It is important to keep the temperature above °C.
  • the rolling reduction ratio for each rolling pass effectively promotes dynamic recrystallization and / or static recrystallization of the austenite phase.
  • rolling strain is effectively applied to the central portion of the plate thickness of the steel slab.
  • coarse Cr-based carbides formed along the casting direction in the vicinity of the boundary between the columnar crystal and the equiaxed crystal of the steel slab are more effectively eliminated. Therefore, among the rolling passes in hot rolling, the number of rolling passes at which the end temperature is 1050 ° C. or higher and the rolling reduction is 20% or higher (hereinafter, also referred to as rolling passes satisfying predetermined conditions) is 3 or more. There is a need to.
  • the upper limit of the number of rolling passes that satisfy the predetermined conditions is not particularly limited, but if it is excessively increased, a large amount of heat is required to maintain the rolling temperature, which leads to an increase in manufacturing cost.
  • the number of rolling passes that satisfy the conditions is preferably 10 or less.
  • the upper limit of the rolling reduction ratio for each rolling pass in hot rolling is not particularly limited, but if the rolling reduction ratio for each rolling pass becomes excessively large, the rolling load increases and rolling becomes difficult. Therefore, the rolling reduction ratio for each rolling pass is preferably 60% or less.
  • the rolling reduction ratio for each rolling pass is ([plate thickness of the material to be rolled (mm) at the start of the rolling pass]-[plate thickness of the material to be rolled (mm) at the end of the rolling pass]) / It was obtained as [plate thickness (mm) of the material to be rolled at the start of the rolling pass] ⁇ 100.
  • the number of rolling passes (total number) for hot rolling is preferably 8 to 20 passes. Further, hot rolling is generally composed of rough rolling and finish rolling. In this case, it is preferable that the number of rolling passes for rough rolling is 3 to 10 and the number of rolling passes for finish rolling is 5 to 10.
  • the rolling end temperature is preferably 900 to 1100 ° C. Further, the total rolling reduction in hot rolling is preferably 85.0 to 99.8%.
  • the hot-rolled steel sheet is wound.
  • the austenite phase is transformed into a ferrite phase, and the metal structure of the hot-rolled steel sheet is made mainly of the ferrite phase.
  • the winding temperature is set to 600 ° C. or higher.
  • the winding temperature is preferably 650 ° C. or higher, more preferably 700 ° C.
  • the upper limit of the winding temperature is not particularly limited, but is preferably 850 ° C. or lower.
  • the take-up temperature exceeds 850 ° C., the take-up temperature becomes a two-phase temperature range of an austenite phase and a ferrite phase. Therefore, the stability of the austenite phase becomes high, and the transformation from the austenite phase to the ferrite phase is delayed.
  • the austenite phase may be transformed into a hard martensite phase after air cooling (of the wound steel sheet) and before annealing the hot-rolled sheet.
  • the hot-rolled steel sheet may be significantly hardened or have a poor shape, which is not preferable.
  • Hot-rolled sheet annealing step Then, as a third step, the hot-rolled steel sheet obtained as described above is annealed by hot-rolled sheet to obtain a hot-rolled and annealed steel sheet.
  • the holding temperature is 750 to 900 ° C.
  • the holding time is 10 minutes or more.
  • Holding temperature for hot-rolled sheet annealing 750-900 ° C
  • the hot-rolled plate annealing is performed for the purpose of suppressing cracking (hereinafter, also referred to as machining cracking) during machining into a predetermined shape of a cutting tool or the like. Then, in this hot-rolled sheet annealing, the rolled structure (metal structure composed of strained crystal grains) formed by hot rolling is replaced with ferrite phase crystal grains containing almost no strain by recrystallization.
  • hot-rolled sheet annealing temperature when the holding temperature of hot-rolled sheet annealing (hereinafter, also referred to as hot-rolled sheet annealing temperature) becomes less than 750 ° C., the rolled structure formed during hot rolling remains. As a result, the ductility of the hot-rolled annealed steel sheet is reduced, and processing cracks are likely to occur. Further, when the hot-rolled plate annealing temperature exceeds 900 ° C., the crystal grains become coarse and the toughness decreases. As a result, processing cracks are likely to occur. Therefore, the hot-rolled sheet annealing temperature is in the range of 750 to 900 ° C. The hot-rolled plate annealing temperature is preferably 800 ° C. or higher.
  • the hot-rolled sheet annealing temperature is preferably 875 ° C. or lower, more preferably 850 ° C. or lower.
  • the hot-rolled sheet annealing temperature may be constant during holding, and may not always be constant during holding as long as it is within the above temperature range. The same applies to the cold-rolled sheet annealing temperature, quenching temperature, and tempering temperature described below.
  • Holding time of hot-rolled sheet annealing If the holding time of hot-rolled sheet annealing is less than 10 minutes, the material in the steel sheet cannot be sufficiently uniformized. Therefore, the holding time for hot-rolled sheet annealing is set to 10 minutes or more.
  • the holding time for hot-rolled sheet annealing is preferably 3 hours or longer, more preferably 6 hours or longer. If the holding time of hot-rolled sheet annealing exceeds 96 hours, the oxide scale may become thick and subsequent descaling treatment may become difficult. Therefore, the holding time for hot-rolled sheet annealing is preferably 96 hours or less.
  • the holding time for hot-rolled sheet annealing is preferably 24 hours or less, more preferably 12 hours or less.
  • cold rolling may be optionally performed as the fourth step, and the cold-rolled plate may be annealed as the fifth step.
  • the hot-rolled annealed steel sheet obtained after annealing the hot-rolled sheet is cold-rolled to obtain a cold-rolled steel sheet.
  • the cold rolling method is not particularly limited, and for example, a tandem mill or a cluster mill can be used.
  • the reduction ratio in cold rolling is also not particularly limited, but the reduction ratio in cold rolling is preferably 50% or more from the viewpoint of formability after annealing of a cold-rolled sheet and shape correction of a steel sheet. Further, from the viewpoint of avoiding an excessive rolling load, the rolling reduction in cold rolling is preferably 95% or less.
  • the cold-rolled steel sheet obtained after cold rolling has a holding temperature of 700 to 850 ° C. and a holding time of 5 seconds or more.
  • the rolled sheet is annealed to obtain a cold rolled annealed steel sheet.
  • the main purpose of cold rolled sheet annealing is to remove the rolled structure formed by cold rolling by recrystallization.
  • the holding temperature of the cold-rolled plate annealing (hereinafter, also referred to as the cold-rolled plate annealing temperature) is less than 700 ° C.
  • the rolled structure formed by the cold rolling remains and is obtained after the cold-rolled plate annealing.
  • the workability of cold-rolled annealed steel sheets is reduced.
  • the holding temperature in the cold-rolled sheet annealing exceeds 850 ° C.
  • an austenite phase is formed, and the austenite phase is transformed into a martensite phase during cooling after holding. Therefore, the cold-rolled annealed steel sheet obtained after the annealed cold-rolled sheet is hardened and the ductility is lowered, and as a result, processing cracks are caused.
  • the cold-rolled plate annealing temperature is in the range of 700 to 850 ° C.
  • the cold rolled sheet annealing temperature is preferably 720 ° C. or higher.
  • the cold rolled sheet annealing temperature is preferably 830 ° C. or lower.
  • the holding time of the cold-rolled sheet annealing is set to 5 seconds or more.
  • the holding time of the cold rolled sheet annealing is preferably 15 seconds or more.
  • the holding time for cold-rolled sheet annealing is preferably 24 hours or less.
  • the holding time for cold rolled sheet annealing is more preferably 15 minutes or less.
  • Quenching treatment step The hot-rolled annealed steel sheet, cold-rolled annealed steel sheet or cold-rolled annealed steel sheet obtained as described above is processed into, for example, a predetermined shape, and then, as a sixth step, a holding temperature. : 950 to 1200 ° C., holding time: 5 seconds to 30 minutes, average cooling rate after holding: 1 ° C./sec or more may be subjected to quenching treatment to obtain a hardened steel plate.
  • the holding temperature of the quenching treatment (hereinafter, also referred to as the quenching temperature) is less than 950 ° C., the austenite phase is not sufficiently formed during heating and holding in the quenching treatment, and sufficient quenching does not occur. If the quenching temperature exceeds 1200 ° C., a delta ferrite phase may be formed in the metal structure during heating and holding in the quenching treatment, and quenching may not be sufficiently performed. In addition, the crystal grains may be remarkably coarsened, causing shrinkage or processing cracks during cooling. Therefore, the quenching temperature is set in the range of 950 to 1200 ° C.
  • the quenching temperature is preferably 1000 ° C. or higher.
  • the quenching temperature is preferably 1150 ° C. or lower.
  • the holding time of the quenching treatment is less than 5 seconds, the austenite phase is not sufficiently formed during heating and holding, and sufficient quenching is not performed.
  • the holding time in the quenching treatment exceeds 30 minutes, coarsening of crystal grains may occur and processing cracks may occur. Therefore, the holding time of the quenching process is in the range of 5 seconds to 30 minutes.
  • the holding time in the quenching treatment is preferably 15 seconds or more.
  • the holding time in the quenching treatment is preferably 300 seconds or less, more preferably 120 seconds or less.
  • the average cooling rate after holding in the quenching treatment is set to 1 ° C./sec or more.
  • the average cooling rate after holding in the quenching treatment is preferably 5 ° C./sec or higher, more preferably 10 ° C./sec or higher.
  • the upper limit of the average cooling rate after holding in the quenching treatment is not particularly limited, but excessive quenching may cause deterioration of the steel sheet shape or quench cracking. Therefore, the average cooling rate after holding in the quenching treatment is preferably 1000 ° C./sec or less.
  • the cooling method is not particularly limited, and various methods such as air cooling, gas injection cooling, mist water cooling, roll cooling, water immersion, and mold cooling can be used.
  • the above-mentioned tempered steel sheet is further tempered as a seventh step with a holding temperature of 100 to 800 ° C. and a holding time of 5 minutes or more. It may be treated to be a tempered steel sheet.
  • the tempering temperature is in the range of 100 to 800 ° C.
  • the tempering temperature is preferably 200 ° C. or higher, more preferably 400 ° C. or higher.
  • the tempering temperature is preferably 750 ° C. or lower, more preferably 700 ° C. or lower.
  • the tempering time is set to 5 minutes or more.
  • the tempering time is preferably 10 minutes or longer, more preferably 15 minutes or longer.
  • the hardness tends to decrease as the tempering time becomes longer, but when the tempering time exceeds 60 minutes, the hardness becomes almost constant. Therefore, the tempering time is preferably 60 minutes or less.
  • the tempering time is more preferably 50 minutes or less, still more preferably 40 minutes or less.
  • pickling treatment, shot blasting, surface grinding, etc. are optionally performed after, for example, a hot rolling step, a hot rolling plate annealing step, a cold rolling step, a cold rolling plate annealing step, a quenching step, and a tempering step. You may. Further, depending on the application, temper rolling may be performed after the hot rolling step, the hot rolling plate annealing step, the cold rolling plate annealing step, the quenching treatment step, the tempering treatment step, and the like.
  • a knife such as a kitchen knife, scissors, a medical scalpel, a cutlery such as a knife or fork for a table, a spoon, and a precision tool such as tweezers can be obtained.
  • the end temperature after the 9th pass in hot rolling is also omitted in Table 2.
  • some hot-rolled annealed steel sheets were further subjected to cold rolling and / or cold-rolled sheet annealing under the conditions shown in Table 3 to obtain cold-rolled steel sheets and / or cold-rolled annealed steel sheets.
  • the metal structures of the hot-rolled and cold-rolled steel sheets, the cold-rolled steel sheets, and the cold-rolled and hardened steel sheets thus obtained were observed by the above-mentioned methods, and the metal structures were identified. The results are shown in Table 4. However, No. In No. 35, since cracks occurred during winding of the hot-rolled steel sheet, the metallographic structure was not identified and the subsequent evaluation was not performed.
  • the hot-rolled annealed steel sheet, the cold-rolled annealed steel sheet and the cold-rolled annealed steel sheet obtained as described above were punched in a rolling direction: 300 mm ⁇ width direction: 50 mm. Then, the processed steel sheet was quenched by air cooling under the conditions of quenching temperature: 1050 ° C., holding time: 15 minutes, and average cooling rate from the quenching temperature after holding to 400 ° C.: 5 ° C./s. ... In addition, No. 1A and 3A-1 and 3A-2 are No. 1 after quenching treatment. The steel sheets 1 and 3 are further tempered under the conditions shown in Table 3 (tempering steel sheet).
  • the hardness and surface quality were evaluated as follows. The hardness was evaluated using a hardened steel sheet. However, the tempered No. In 1A, 3A-1 and 3A-2, the hardness of the tempered steel sheet was also evaluated. In addition, the surface quality was evaluated by the finally obtained steel sheet, that is, No. In Nos. 1 to 37, the hardened steel sheet was used as No. In 1A, 3A-1 and 3A-2, tempered steel sheets were used.
  • FIG. 1 For reference, No. 1 of the invention example in which good surface quality was obtained. An optical microscope microstructure photograph in a cross section parallel to the rolling direction of No. 1 is shown in FIG. In addition, No. 1 of Comparative Example in which good surface quality could not be obtained. An optical microscope microstructure photograph of a cross section parallel to the rolling direction of 30 is shown in FIG.
  • the stainless steel plate of the present invention has high hardness and good surface quality, it can be used for knives such as kitchen knives, scissors, medical scalpels, cutlery such as table knives, forks and spoons, and precision tools such as tweezers. It can be suitably used as a material.

Abstract

The present invention is a stainless steel sheet which has a predetermined component composition, wherein the total volume percentage of Cr-based carbides having a particle diameter of 2.0 μm or more is 10% or less.

Description

ステンレス鋼板およびその製造方法、刃物、ならびに、カトラリーStainless steel sheet and its manufacturing method, cutlery, and cutlery
　本発明は、包丁やハサミ、医療用メス等の刃物、食卓用のナイフやフォーク、スプーン等のカトラリー、および、ピンセット等の精密工具に用いて好適な高い硬度と良好な表面品質とを有するステンレス鋼板に関するものである。 The present invention is a stainless steel having high hardness and good surface quality suitable for use in knives such as kitchen knives and scissors, medical scalpels, cutlery such as table knives and forks, spoons, and precision tools such as tweezers. It is about steel plates.
　包丁やハサミ、医療用メス等の刃物、および、ピンセット等の精密工具の素材には、ステンレス鋼板が使用される場合がある。
　例えば、包丁の場合、ステンレス鋼板をプレス加工等により所定形状にブランキングまたは鍛造加工する。ついで、所定形状に加工したステンレス鋼板に、焼入れ処理、または、焼入れ処理および焼戻し処理を施して硬質化させる。そして、硬質化させたステンレス鋼板に、刃付け研磨（刃先となる部分を研磨により薄肉化する処理）等を施し、最終製品（包丁）とする。 Stainless steel plates may be used as materials for knives, scissors, scalpels and other cutting tools, and precision tools such as tweezers.
For example, in the case of a kitchen knife, a stainless steel plate is blanked or forged into a predetermined shape by press working or the like. Then, the stainless steel sheet processed into a predetermined shape is subjected to quenching treatment, quenching treatment and tempering treatment to harden it. Then, the hardened stainless steel sheet is subjected to blade polishing (a process of thinning the portion to be the blade edge by polishing) or the like to obtain a final product (kitchen knife).
　上記した刃物および精密工具等の用途に用いられるステンレス鋼としては、例えば、１３ｍａｓｓ％Ｃｒ－０．３ｍａｓｓ％Ｃ鋼（ＪＩＳ　Ｇ　４３０４およびＪＩＳ　Ｇ　４３０５で規定されるＳＵＳ４２０Ｊ２）等が挙げられる。 Examples of stainless steel used for the above-mentioned applications such as cutting tools and precision tools include 13 mass% Cr-0.3 mass% C steel (SUS420J2 defined by JIS G 4304 and JIS G 4305).
　ところで、刃物および精密工具では、刃先の磨耗による切れ味の低下や錆の発生を極力抑制して、砥ぎ等のメンテナンス頻度を低減することも要求される。
　近年、この要求が特に高まっており、十分な耐食性を確保したうえで、切れ味が高く、さらに刃先の磨耗による切れ味の低下を長期にわたり抑制し得る、高硬度の高級刃物への市場ニーズが高まっている。 By the way, in cutting tools and precision tools, it is also required to reduce the frequency of maintenance such as grinding by suppressing the decrease in sharpness and the occurrence of rust due to the wear of the cutting edge as much as possible.
In recent years, this demand has been particularly increasing, and there is an increasing market need for high-hardness high-grade blades that can ensure sufficient corrosion resistance, have high sharpness, and can suppress deterioration of sharpness due to wear of the cutting edge for a long period of time. There is.
　このような高硬度の高級刃物に使用されるステンレス鋼として、例えば、欧州統一規格：ＥＮ１．４１１６に準拠する１４ｍａｓｓ％Ｃｒ－０．５ｍａｓｓ％Ｃ鋼が挙げられる。この欧州統一規格：ＥＮ１．４１１６に準拠する１４ｍａｓｓ％Ｃｒ－０．５ｍａｓｓ％Ｃ鋼は、１３ｍａｓｓ％Ｃｒ－０．３ｍａｓｓ％Ｃ鋼に比べてＣ量を増加させて硬度を高めた鋼である。
　また、特許文献１には、
「Ｃ ：０．８８ｍａｓｓ％以上１．２ｍａｓｓ％以下、Ｃｒ：１２．５ｍａｓｓ％以上１６．５０ｍａｓｓ％以下、Ｓｉ：０．０５ｍａｓｓ％以上０．２０ｍａｓｓ％以下、Ｎ ：０．００１ｍａｓｓ％以上０．０２ｍａｓｓ％以下、Ｍｎ：１．０ｍａｓｓ％以下、Ｃｕ：１．０ｍａｓｓ％以下、Ｐ ：０．０３ｍａｓｓ％以下、Ｓ ：０．０１０ｍａｓｓ％以下、および、Ｎｉ：１．０ｍａｓｓ％以下、を含み、残部がＦｅおよび不可避的不純物からなることを特徴とするステンレス帯鋼。」
が開示されている。 Examples of stainless steel used for such high-hardness high-grade blades include 14 mass% Cr-0.5 mass% C steel conforming to the European unified standard: EN1.4116. The 14 mass% Cr-0.5 mass% C steel conforming to this European unified standard: EN1.4116 is a steel in which the amount of C is increased and the hardness is increased as compared with the 13 mass% Cr-0.3 mass% C steel.
Further, in Patent Document 1,
"C: 0.88 mass% or more and 1.2 mass% or less, Cr: 12.5 mass% or more and 16.50 mass% or less, Si: 0.05 mass% or more and 0.20 mass% or less, N: 0.001 mass% or more and 0.02 mass" % Or less, Mn: 1.0 mass% or less, Cu: 1.0 mass% or less, P: 0.03 mass% or less, S: 0.010 mass% or less, and Ni: 1.0 mass% or less, and the balance is A stainless steel strip characterized by consisting of Fe and unavoidable impurities. "
Is disclosed.
特許第５０１０８１９号Patent No. 5010819
　しかし、欧州統一規格：ＥＮ１．４１１６に準拠する１４ｍａｓｓ％Ｃｒ－０．５ｍａｓｓ％Ｃ鋼や特許文献１に開示されるステンレス帯鋼から得た鋼板に、研磨や刃付け加工等を施すと、圧延方向に沿った筋模様が発生して外観の美麗性を大きく損う場合がある。 However, when a steel sheet obtained from 14 mass% Cr-0.5 mass% C steel conforming to the European unified standard: EN1.4116 or a stainless steel strip disclosed in Patent Document 1 is polished or bladed, it is rolled. A streak pattern along the direction may occur, which may greatly impair the beauty of the appearance.
　このような筋模様が発生した場合、研磨工程を追加する等して、筋模様を除去する必要がある。しかし、研磨工程の追加は製造コストの増加を招く。また、筋模様が著しい場合には、筋模様が除去しきれなかったり、筋模様を除去するために必要な研磨量が多くなって所定の形状が得られなくなる場合がある。その結果、歩留まりおよび生産性の大幅な低下を招く。 When such a streak pattern occurs, it is necessary to remove the streak pattern by adding a polishing process or the like. However, the addition of a polishing process leads to an increase in manufacturing cost. Further, when the streak pattern is remarkable, the streak pattern may not be completely removed, or the amount of polishing required for removing the streak pattern may increase and a predetermined shape may not be obtained. The result is a significant reduction in yield and productivity.
　そのため、製品として使用する際に高い硬度を有し、かつ、製品に加工する際の筋模様の発生を抑制した良好な表面品質を有する、ステンレス鋼板の開発が求められているのが現状である。 Therefore, the current situation is that there is a demand for the development of stainless steel sheets that have high hardness when used as products and have good surface quality that suppresses the occurrence of streaks when processed into products. ..
　本発明は、上記の現状に鑑み開発されたものであって、製品として使用する際に高い硬度（以下、単に高い硬度ともいう）を有し、かつ、製品に加工する際の筋模様の発生を抑制した良好な表面品質（以下、単に良好な表面品質ともいう）を有する、ステンレス鋼板を提供することを目的とする。
　また、本発明は、上記のステンレス鋼板の製造方法を提供することを目的とする。
　さらに、本発明は、上記のステンレス鋼板を用いてなる刃物およびカトラリーを提供することを目的とする。
　なお、上述したように、本発明のステンレス鋼板は、刃物およびカトラリーといった製品として使用する際に高い硬度が得られるものを対象としている。すなわち、本発明のステンレス鋼板には、硬質化させた後（焼き入れ処理後）の鋼板だけでなく、硬質化させる前（焼き入れ処理前）の製品素材となる鋼板が含まれる。 The present invention has been developed in view of the above-mentioned current situation, has high hardness when used as a product (hereinafter, also simply referred to as high hardness), and generates streaks when processed into a product. It is an object of the present invention to provide a stainless steel sheet having a good surface quality (hereinafter, also simply referred to as a good surface quality) in which the above is suppressed.
Another object of the present invention is to provide the above-mentioned method for manufacturing a stainless steel sheet.
Furthermore, an object of the present invention is to provide a cutlery and a cutlery made of the above-mentioned stainless steel plate.
As described above, the stainless steel sheet of the present invention is intended for those having high hardness when used as products such as cutlery and cutlery. That is, the stainless steel sheet of the present invention includes not only a steel sheet after being hardened (after quenching) but also a steel sheet as a product material before being hardened (before quenching).
　さて、発明者らは、上記の目的を達成すべく、鋭意検討を重ねた。
　まず、発明者らは、欧州統一規格：ＥＮ１．４１１６に準拠する１４ｍａｓｓ％Ｃｒ－０．５ｍａｓｓ％Ｃ鋼に、研磨や刃付け加工（以下、単に研磨ともいう）等を施した場合に、筋模様が発生する原因について、検討した。
　具体的には、
・欧州統一規格：ＥＮ１．４１１６に準拠する１４ｍａｓｓ％Ｃｒ－０．５ｍａｓｓ％Ｃ鋼の成分組成を有する鋼板（以下、単に鋼板ａともいう）、および
・ＪＩＳ　Ｇ　４３０４およびＪＩＳ　Ｇ　４３０５で規定されるＳＵＳ４２０Ｊ２に相当する１３ｍａｓｓ％Ｃｒ－０．３ｍａｓｓ％Ｃ鋼の成分組成を有する鋼板（以下、単に鋼板ｂともいう）
を、それぞれ従来公知の方法により同じ条件で製造し、製造した鋼板に、同じ条件で研磨を施した。
　その結果、鋼板ｂでは研磨を施しても筋模様が発生しなかった。一方、鋼板ａでは、研磨を施すと、筋模様が発生した。 Now, the inventors have made extensive studies to achieve the above-mentioned purpose.
First, the inventors found that when 14 mass% Cr-0.5 mass% C steel conforming to the European unified standard: EN1.4116 was polished or sharpened (hereinafter, also simply referred to as polishing), streaks were formed. The cause of the pattern was examined.
In particular,
-European unified standard: Steel sheet having a component composition of 14 mass% Cr-0.5 mass% C steel conforming to EN1.4116 (hereinafter, also simply referred to as steel sheet a), and ... JIS G 4304 and JIS G 4305. A steel sheet having a component composition of 13 mass% Cr-0.3 mass% C steel corresponding to SUS420J2 (hereinafter, also simply referred to as steel sheet b).
Was produced under the same conditions by a conventionally known method, and the produced steel sheet was polished under the same conditions.
As a result, the steel plate b did not have a streak pattern even after polishing. On the other hand, in the steel plate a, when polishing was performed, a streak pattern was generated.
　上記の結果から、発明者らは、次のように考えるに至った。
　すなわち、鋼板ａと鋼板ｂとでは、成分組成の違いにより、同じ製造条件で製造しても、析出物の析出状態が大きく異なるものとなる。そして、この析出物の析出状態の違いが原因で、鋼板ａでは筋模様が発生する。 From the above results, the inventors came to think as follows.
That is, even if the steel sheet a and the steel sheet b are manufactured under the same manufacturing conditions due to the difference in the component composition, the precipitation state of the precipitates is significantly different. Then, due to the difference in the precipitation state of the precipitate, a streak pattern is generated on the steel plate a.
　この考えに基づき、発明者らは、鋼板ａおよび鋼板ｂの金属組織を観察し、両者を詳細に対比した。
　その結果、筋模様が発生した鋼板ａでは、図２のように、金属組織中に粗大なＣｒ系炭化物が圧延方向に連なって存在しており、これが原因となって、筋模様が発生することを知見した。
　すなわち、Ｃｒ系炭化物は、ステンレス鋼板の母材（焼入れ前後ともに）よりも硬質である。そのため、金属組織に粗大なＣｒ系炭化物が存在していると、当該Ｃｒ系炭化物が存在している部位では、他の部位に比べて、研磨量が少なくなる。その結果、研磨後に、局所的に凸部が生じ、これらが筋模様として顕在化する。
　特に、鋼板ａ（欧州統一規格：ＥＮ１．４１１６に準拠する１４ｍａｓｓ％Ｃｒ－０．５ｍａｓｓ％Ｃ鋼）の成分組成では、より高い硬度を得るために、鋼板ｂ（１３ｍａｓｓ％Ｃｒ－０．３ｍａｓｓ％Ｃ鋼）に比べて多量のＣやＣｒが含まれている。そのため、鋼板ｂでは従来公知の方法により製造しても粗大なＣｒ系炭化物が多量には生成しないが、同じ条件で製造した鋼板ａでは、粗大なＣｒ系炭化物が多量に生成して筋模様が発生する。 Based on this idea, the inventors observed the metallographic structures of the steel plate a and the steel plate b and compared them in detail.
As a result, in the steel sheet a in which the streaks are generated, as shown in FIG. 2, coarse Cr-based carbides are continuously present in the metal structure in the rolling direction, which causes the streaks to be generated. Was found.
That is, the Cr-based carbide is harder than the base material of the stainless steel sheet (both before and after quenching). Therefore, when coarse Cr-based carbides are present in the metal structure, the amount of polishing is smaller in the portion where the Cr-based carbides are present than in other portions. As a result, after polishing, convex portions are locally generated, and these are manifested as streaks.
In particular, in the composition of the steel sheet a (14 mass% Cr-0.5 mass% C steel conforming to the European unified standard: EN1.4116), in order to obtain higher hardness, the steel sheet b (13 mass% Cr-0.3 mass%) is used. It contains a large amount of C and Cr as compared with C steel). Therefore, the steel sheet b does not generate a large amount of coarse Cr-based carbides even if it is manufactured by a conventionally known method, but the steel sheet a manufactured under the same conditions produces a large amount of coarse Cr-based carbides and has a streak pattern. appear.
　そして、発明者らは、上記の知見を基に、さらに検討を重ね、以下の知見を得た。
　すなわち、粒径：２．０μｍ以上のＣｒ系炭化物が研磨時の筋模様の発生に深く影響を与えている。そして、このような粗大なＣｒ系炭化物の生成を極力抑制する、特には、粒径：２．０μｍ以上のＣｒ系炭化物の体積率を１０％以下に抑制することによって、研磨時の筋模様の発生が大幅に抑制される。 Then, the inventors further studied based on the above findings and obtained the following findings.
That is, Cr-based carbides having a particle size of 2.0 μm or more have a profound effect on the generation of streaks during polishing. Then, by suppressing the formation of such coarse Cr-based carbides as much as possible, in particular, by suppressing the volume fraction of Cr-based carbides having a particle size of 2.0 μm or more to 10% or less, the streak pattern during polishing is suppressed. Occurrence is greatly suppressed.
　また、発明者らは、さらに検討を重ね、以下の知見を得た。
　すなわち、上記の粗大なＣｒ系炭化物は、鋳造時にスラブ断面の柱状晶と等軸晶の境界近傍において鋳造方向に沿って生成したものである。また、鋳造時に生成した粗大なＣｒ系炭窒物は、従来公知の一般的な製造条件では、鋳造工程以降の熱間圧延、熱延板焼鈍、冷間圧延および冷延板焼鈍工程を経た後にも、依然として圧延方向（鋳造方向と同一方向である）に残存する。 In addition, the inventors further studied and obtained the following findings.
That is, the above-mentioned coarse Cr-based carbide is produced along the casting direction in the vicinity of the boundary between the columnar crystal and the equiaxed crystal in the slab cross section at the time of casting. Further, the coarse Cr-based carbon dioxide produced during casting is subjected to hot rolling, hot rolling plate annealing, cold rolling and cold rolling plate annealing steps after the casting step under general manufacturing conditions known conventionally. Still remains in the rolling direction (the same direction as the casting direction).
　そこで、発明者らは、上記の知見を基に、高い硬度を得ながら、粗大なＣｒ系炭化物の生成を防止する方法について検討を重ねた。
　その結果、以下の知見を得た。
（１）成分組成を適正に調整する、特には、Ｃ含有量およびＣｒ含有量をそれぞれ０．４５～０．６０質量％、および、１３．０％以上１６．０％未満の範囲に調整し、
（２）そのうえで、鋼スラブの加熱、熱間圧延および熱延板焼鈍条件を適正に制御する、
　具体的には、
　（ａ）鋼スラブを１２００～１３５０℃で３０分以上保持し、かつ、
　（ｂ）熱間圧延における圧延パスのうち、終了温度：１０５０℃以上で、かつ、圧下率：２０％以上の圧延パス数を３パス以上とし、
　（ｃ）また、熱延鋼板の巻取り温度を６００℃以上とする、
ことが重要である。これにより、Ｃ含有量およびＣｒ含有量を一定量以上含有させる場合にも、粗大なＣｒ系炭化物の生成を抑制し、研磨時の筋模様の発生を有効に防止できる。 Therefore, based on the above findings, the inventors have repeatedly studied a method for preventing the formation of coarse Cr-based carbides while obtaining high hardness.
As a result, the following findings were obtained.
(1) Appropriately adjust the component composition, in particular, adjust the C content and Cr content to 0.45 to 0.60% by mass and 13.0% or more and less than 16.0%, respectively. ,
(2) Then, the heating, hot rolling and hot rolling sheet annealing conditions of the steel slab are appropriately controlled.
In particular,
(A) Hold the steel slab at 1200 to 1350 ° C. for 30 minutes or more, and
(B) Of the rolling passes in hot rolling, the number of rolling passes having an end temperature of 1050 ° C. or higher and a rolling reduction ratio of 20% or higher shall be 3 or more.
(C) Further, the winding temperature of the hot-rolled steel sheet is set to 600 ° C. or higher.
This is very important. As a result, even when the C content and the Cr content are contained in a certain amount or more, the formation of coarse Cr-based carbides can be suppressed, and the generation of streaks during polishing can be effectively prevented.
　なお、上記のように製造条件を制御することにより、粗大なＣｒ系炭化物の生成が抑制される理由について、発明者らは次のように考えている。
　すなわち、上記（２）（ａ）のように、鋼スラブを１２００～１３５０℃で３０分以上保持することにより、鋳造工程で生成した粗大なＣｒ系炭化物のオーステナイト相への固溶（Ｃｒ系炭化物がＣｒ原子、Ｃ原子等に分解してオーステナイト相中に原子状態で取り込まれること）が促進される。
　また、この状態で、上記（２）（ｂ）のように、熱間圧延における圧延パスを、高い温度でかつ、高い圧下率で行うことにより、Ｃｒ系炭化物のオーステナイト相への固溶がさらに促進される。加えて、圧延ひずみが鋼スラブの板厚中央部まで効果的に付与される。これにより、鋼スラブの柱状晶と等軸晶の境界部近傍において鋳造方向に沿って生成した粗大なＣｒ系炭化物が解消される。また、元素の転位上拡散（格子欠陥である転位を介した原子移動）が促進される。これにより、Ｃｒ系炭化物のオーステナイト相への固溶が一層促進される。さらに、オーステナイト相の動的再結晶および／または静的再結晶を促進することによって、オーステナイト相の結晶粒が微細化される。これにより、上記（２）（ｃ）における熱延鋼板の巻取りの際に、オーステナイト相の粒界から析出するＣｒ系炭化物の析出サイトが増加して、再析出するＣｒ系炭化物も微細化される。なお、再結晶とは、ひずみを有する結晶粒内または結晶粒界から、ひずみをほとんど含まない結晶粒が生成する現象である。
　上記の相乗効果により、Ｃ含有量およびＣｒ含有量を一定量含有させる場合にも、粗大なＣｒ系炭化物の生成を抑制して、研磨時の筋模様の発生を防止することが可能となる。
　本発明は、上記の知見に基づき、さらに検討を加えて完成されたものである。 The inventors consider the reason why the formation of coarse Cr-based carbides is suppressed by controlling the production conditions as described above.
That is, as described in (2) and (a) above, by holding the steel slab at 1200 to 1350 ° C. for 30 minutes or more, the coarse Cr-based carbide produced in the casting step is solid-solved in the austenite phase (Cr-based carbide). Is decomposed into Cr atoms, C atoms, etc. and incorporated into the austenite phase in the atomic state) is promoted.
Further, in this state, as described in (2) and (b) above, by performing the rolling pass in hot rolling at a high temperature and a high rolling reduction, the solid solution of the Cr-based carbide in the austenite phase is further increased. It is promoted. In addition, rolling strain is effectively applied to the central portion of the steel slab thickness. As a result, coarse Cr-based carbides formed along the casting direction in the vicinity of the boundary between the columnar crystal and the equiaxed crystal of the steel slab are eliminated. In addition, diffusion on dislocations of elements (atomic movement via dislocations, which are lattice defects) is promoted. This further promotes the solid solution of Cr-based carbides into the austenite phase. Furthermore, by promoting dynamic recrystallization and / or static recrystallization of the austenite phase, the crystal grains of the austenite phase are refined. As a result, when the hot-rolled steel sheet is wound in the above (2) and (c), the precipitation sites of Cr-based carbides precipitated from the grain boundaries of the austenite phase increase, and the Cr-based carbides reprecipitated are also refined. NS. Note that recrystallization is a phenomenon in which crystal grains containing almost no strain are generated from within the crystal grains having strain or from the grain boundaries.
Due to the above synergistic effect, even when the C content and the Cr content are contained in a certain amount, it is possible to suppress the formation of coarse Cr-based carbides and prevent the generation of streaks during polishing.
The present invention has been completed with further studies based on the above findings.
　すなわち、本発明の要旨構成は次のとおりである。
１．質量％で、
　Ｃ：０．４５～０．６０％、
　Ｓｉ：０．０５～１．００％、
　Ｍｎ：０．０５～１．００％、
　Ｐ：０．０５％以下、
　Ｓ：０．０２０％以下、
　Ｃｒ：１３．０％以上１６．０％未満、
　Ｎｉ：０．１０～１．００％および
　Ｎ：０．０１０～０．２００％
を含有し、残部がＦｅおよび不可避的不純物からなる成分組成を有し、
　粒径：２．０μｍ以上のＣｒ系炭化物の合計の体積率が１０％以下である、ステンレス鋼板。 That is, the gist structure of the present invention is as follows.
1. 1. By mass%
C: 0.45 to 0.60%,
Si: 0.05 to 1.00%,
Mn: 0.05 to 1.00%,
P: 0.05% or less,
S: 0.020% or less,
Cr: 13.0% or more and less than 16.0%,
Ni: 0.10 to 1.00% and N: 0.010 to 0.200%
Has a component composition in which the balance is composed of Fe and unavoidable impurities.
Particle size: A stainless steel sheet having a total volume fraction of Cr-based carbides of 2.0 μm or more of 10% or less.
２．前記成分組成が、さらに、質量％で、
　Ｍｏ：０．０５～１．００％、
　Ｃｕ：０．０５～１．００％および
　Ｃｏ：０．０５～０．５０％
のうちから選ばれる１種または２種以上を含有する、前記１に記載のステンレス鋼板。 2. The component composition is further increased by mass%.
Mo: 0.05-1.00%,
Cu: 0.05 to 1.00% and Co: 0.05 to 0.50%
The stainless steel sheet according to 1 above, which contains one or more selected from the above.
３．前記成分組成が、さらに、質量％で、
　Ａｌ：０．００１～０．１００％、
　Ｔｉ：０．０１～０．１０％、
　Ｎｂ：０．０１～０．１０％、
　Ｖ：０．０５～０．５０％、
　Ｚｒ：０．０１～０．１０％、
　Ｍｇ：０．０００２～０．００５０％、
　Ｂ：０．０００２～０．００５０％、
　Ｃａ：０．０００３～０．００３０％および
　ＲＥＭ：０．０１～０．１０％
のうちから選ばれる１種または２種以上を含有する、前記１または２に記載のステンレス鋼板。 3. 3. The component composition is further increased by mass%.
Al: 0.001 to 0.100%,
Ti: 0.01 to 0.10%,
Nb: 0.01 to 0.10%,
V: 0.05 to 0.50%,
Zr: 0.01-0.10%,
Mg: 0.0002 to 0.0050%,
B: 0.0002 to 0.0050%,
Ca: 0.0003 to 0.0030% and REM: 0.01 to 0.10%
The stainless steel sheet according to 1 or 2 above, which contains one or more selected from the above.
４．前記１～３のいずれかに記載のステンレス鋼板を製造するための方法であって、
　前記１～３のいずれかに記載の成分組成を有する鋼スラブを、１２００～１３５０℃で３０分以上保持する、第１の工程と、
　前記鋼スラブに、熱間圧延を施して熱延鋼板とし、該熱延鋼板を巻取る、第２の工程と、
　前記熱延鋼板に熱延板焼鈍を施し、熱延焼鈍鋼板とする、第３の工程と、をそなえ、
　前記第２の工程の熱間圧延における圧延パスのうち、終了温度：１０５０℃以上で、かつ、圧下率：２０％以上の圧延パス数が３パス以上であり、また、前記熱延鋼板の巻取り温度が６００℃以上であり、
　前記第３の工程の熱延板焼鈍における保持温度が７５０～９００℃、保持時間が１０分以上である、
　ステンレス鋼板の製造方法。 4. A method for manufacturing the stainless steel sheet according to any one of 1 to 3 above.
The first step of holding the steel slab having the component composition according to any one of 1 to 3 at 1200 to 1350 ° C. for 30 minutes or more.
The second step of hot-rolling the steel slab to obtain a hot-rolled steel sheet and winding the hot-rolled steel sheet.
A third step of subjecting the hot-rolled steel sheet to hot-rolled annealed steel sheet to obtain a hot-rolled and annealed steel sheet is provided.
Among the rolling passes in the hot rolling of the second step, the number of rolling passes having an end temperature of 1050 ° C. or higher and a reduction rate of 20% or higher is 3 or more, and the hot-rolled steel sheet is wound. The taking temperature is 600 ° C or higher,
The holding temperature in the hot-rolled sheet annealing of the third step is 750 to 900 ° C., and the holding time is 10 minutes or more.
Manufacturing method of stainless steel plate.
５．前記熱延焼鈍鋼板に、冷間圧延を施して冷延鋼板とする、第４の工程をそなえる、前記４に記載のステンレス鋼板の製造方法。 5. The method for producing a stainless steel sheet according to the above 4, further comprising a fourth step of cold-rolling the hot-rolled annealed steel sheet to obtain a cold-rolled steel sheet.
６．前記冷延鋼板に、冷延板焼鈍を施して冷延焼鈍鋼板とする、第５の工程をそなえ、
　前記冷延板焼鈍における保持温度が７００～８５０℃、保持時間が５秒以上である、前記５に記載のステンレス鋼板の製造方法。 6. The cold-rolled steel sheet is subjected to cold-rolled sheet annealing to obtain a cold-rolled annealed steel sheet, which is provided with a fifth step.
The method for producing a stainless steel sheet according to 5, wherein the holding temperature in the cold-rolled sheet annealing is 700 to 850 ° C., and the holding time is 5 seconds or more.
７．前記熱延焼鈍鋼板、前記冷延鋼板、または、前記冷延焼鈍鋼板に、焼入れ処理を施す、第６の工程をそなえ、
　前記焼入れ処理における保持温度が９５０～１２００℃、保持時間が５秒～３０分、保持後の平均冷却速度が１℃／秒以上である、前記４～６のいずれかに記載のステンレス鋼板の製造方法。 7. A sixth step of quenching the hot-rolled annealed steel sheet, the cold-rolled steel sheet, or the cold-rolled annealed steel sheet is provided.
The production of the stainless steel sheet according to any one of 4 to 6, wherein the holding temperature in the quenching treatment is 950 to 1200 ° C., the holding time is 5 seconds to 30 minutes, and the average cooling rate after holding is 1 ° C./sec or more. Method.
８．前記焼入れ処理を施した鋼板に、焼戻し処理を施す、第７の工程をそなえ、
　前記焼戻し処理における保持温度が１００～８００℃、保持時間が５分以上である、前記７に記載のステンレス鋼板の製造方法。 8. A seventh step of tempering the hardened steel sheet is provided.
The method for producing a stainless steel sheet according to 7. above, wherein the holding temperature in the tempering treatment is 100 to 800 ° C. and the holding time is 5 minutes or more.
９．前記１～３のいずれかに記載のステンレス鋼板を用いてなる、刃物。 9. A cutting tool made of the stainless steel plate according to any one of 1 to 3 above.
１０．前記１～３のいずれかに記載のステンレス鋼板を用いてなる、カトラリー。 10. A cutlery made of the stainless steel plate according to any one of 1 to 3 above.
　本発明によれば、高い硬度を有し、かつ、良好な表面品質を有するステンレス鋼板を、得ることができる。 According to the present invention, it is possible to obtain a stainless steel sheet having high hardness and good surface quality.
発明例であるＮｏ．１の光学顕微鏡組織写真である。No. which is an example of the invention. It is an optical microscope tissue photograph of 1. 比較例であるＮｏ．３０の光学顕微鏡組織写真である。No. which is a comparative example. 30 optical microscope tissue photographs. 表面品質の評価において、試験片に切削加工を施したときの状態を示す模式図である。It is a schematic diagram which shows the state when the test piece was cut in the evaluation of the surface quality.
　本発明を、以下の実施形態に基づき説明する。
　まず、本発明の一実施形態に係るステンレス鋼板の成分組成について、説明する。なお、成分組成における単位はいずれも「質量％」であるが、以下、特に断らない限り、単に「％」で示す。 The present invention will be described based on the following embodiments.
First, the component composition of the stainless steel sheet according to the embodiment of the present invention will be described. The unit in the component composition is "mass%", but hereinafter, unless otherwise specified, it is simply indicated by "%".
Ｃ：０．４５～０．６０％
　Ｃは、焼入れ処理により得られるマルテンサイト相を硬質化させる効果がある。ここで、Ｃ含有量が０．４５％未満では、焼入れ処理後の硬度が不足し、高級刃物に求められる切れ味が十分に得られない。一方、Ｃ含有量が０．６０％を超えると、製造条件を適正に制御しても、粗大な炭化物の発生を十分には抑制できず、良好な表面品質が得られない。また、焼入れ処理時に焼割れが生じやすくなり、刃物を安定的に製造することが困難となる。
　そのため、Ｃ含有量は０．４５～０．６０％の範囲とする。Ｃ含有量は、好ましくは０．５５％以下、より好ましくは０．５０％以下である。 C: 0.45 to 0.60%
C has the effect of hardening the martensite phase obtained by the quenching treatment. Here, if the C content is less than 0.45%, the hardness after the quenching treatment is insufficient, and the sharpness required for a high-grade blade cannot be sufficiently obtained. On the other hand, when the C content exceeds 0.60%, even if the production conditions are appropriately controlled, the generation of coarse carbides cannot be sufficiently suppressed, and good surface quality cannot be obtained. In addition, quench cracking is likely to occur during the quenching process, making it difficult to stably manufacture the cutting tool.
Therefore, the C content is set in the range of 0.45 to 0.60%. The C content is preferably 0.55% or less, more preferably 0.50% or less.
Ｓｉ：０．０５～１．００％
　Ｓｉは、鋼の溶製時に脱酸剤として作用する。このような効果を得るため、Ｓｉ含有量は０．０５％以上とする。しかし、Ｓｉ含有量が１．００％を超えると、焼入れ処理前に鋼板が過度に硬質化して、刃物等の所定形状に成形する際の加工性が十分に得られなくなる。
　そのため、Ｓｉ含有量は０．０５～１．００％の範囲とする。Ｓｉ含有量は、好ましくは０．２０％以上である。また、Ｓｉ含有量は、好ましくは０．６０％以下である。 Si: 0.05 to 1.00%
Si acts as an antacid during the melting of steel. In order to obtain such an effect, the Si content is set to 0.05% or more. However, if the Si content exceeds 1.00%, the steel sheet becomes excessively hard before the quenching treatment, and it becomes difficult to obtain sufficient workability when forming a predetermined shape such as a cutting tool.
Therefore, the Si content is in the range of 0.05 to 1.00%. The Si content is preferably 0.20% or more. The Si content is preferably 0.60% or less.
Ｍｎ：０．０５～１．００％
　Ｍｎは、オーステナイト相の生成を促進するとともに、焼入れ性を向上させる効果がある。このような効果を得るため、Ｍｎ含有量は０．０５％以上とする。しかし、Ｍｎ含有量が１．００％を超えると、耐食性の低下を招く。
　そのため、Ｍｎ含有量は０．０５～１．００％の範囲とする。Ｍｎ含有量は、好ましくは０．４０％以上である。また、Ｍｎ含有量は、好ましくは０．８０％以下である。 Mn: 0.05 to 1.00%
Mn has the effect of promoting the formation of the austenite phase and improving the hardenability. In order to obtain such an effect, the Mn content is set to 0.05% or more. However, if the Mn content exceeds 1.00%, the corrosion resistance is lowered.
Therefore, the Mn content is set in the range of 0.05 to 1.00%. The Mn content is preferably 0.40% or more. The Mn content is preferably 0.80% or less.
Ｐ：０．０５％以下
　Ｐは、粒界偏析による粒界破壊を助長する元素である。そのため、Ｐは、可能な限り低減することが望ましい。
　よって、Ｐ含有量は０．０５％以下とする。Ｐ含有量は、好ましくは０．０４％以下、より好ましくは０．０３％以下である。
　なお、Ｐ含有量の下限は特に限定されない。ただし、過度の脱Ｐはコストの増加を招くので、Ｐ含有量は０．００５％以上が好ましい。 P: 0.05% or less P is an element that promotes grain boundary fracture due to grain boundary segregation. Therefore, it is desirable to reduce P as much as possible.
Therefore, the P content is set to 0.05% or less. The P content is preferably 0.04% or less, more preferably 0.03% or less.
The lower limit of the P content is not particularly limited. However, since excessive de-P causes an increase in cost, the P content is preferably 0.005% or more.
Ｓ：０．０２０％以下
　Ｓは、ＭｎＳ等の硫化物系介在物として鋼中に存在して、延性や耐食性等を低下させる元素である。そのため、Ｓは、可能な限り低減することが望ましい。
　よって、Ｓ含有量は０．０２０％以下とする。Ｓ含有量は、好ましくは０．０１５％以下である。
　なお、Ｓ含有量の下限は特に限定されない。ただし、過度の脱Ｓはコストの増加を招くので、Ｓ含有量は０．０００５％以上が好ましい。 S: 0.020% or less S is an element that exists in steel as a sulfide-based inclusion such as MnS and lowers ductility, corrosion resistance, and the like. Therefore, it is desirable to reduce S as much as possible.
Therefore, the S content is set to 0.020% or less. The S content is preferably 0.015% or less.
The lower limit of the S content is not particularly limited. However, since excessive de-S causes an increase in cost, the S content is preferably 0.0005% or more.
Ｃｒ：１３．０％以上１６．０％未満
　Ｃｒは、耐食性を向上させる効果がある。このような効果を得るため、Ｃｒ含有量は１３．０％以上とする。しかし、Ｃｒ含有量が１６．０％以上になると、焼入れ処理の加熱・保持時に生成するオーステナイト量が減少する。そのため、焼入れ処理後に得られるマルテンサイト相が減少し、十分な硬度が得られない。よって、Ｃｒ含有量は１３．０％以上１６．０％未満の範囲とする。Ｃｒ含有量は、好ましくは１４．０％以上である。また、Ｃｒ含有量は好ましくは１５．５％以下、より好ましくは１５．０％以下である。 Cr: 13.0% or more and less than 16.0% Cr has an effect of improving corrosion resistance. In order to obtain such an effect, the Cr content is set to 13.0% or more. However, when the Cr content is 16.0% or more, the amount of austenite produced during heating and holding of the quenching treatment decreases. Therefore, the martensite phase obtained after the quenching treatment is reduced, and sufficient hardness cannot be obtained. Therefore, the Cr content is in the range of 13.0% or more and less than 16.0%. The Cr content is preferably 14.0% or more. The Cr content is preferably 15.5% or less, more preferably 15.0% or less.
Ｎｉ：０．１０～１．００％
　Ｎｉは、耐食性を向上させるとともに、焼入れ後の靭性を向上させる効果がある。このような効果を得るため、Ｎｉ含有量は０．１０％以上とする。しかし、Ｎｉ含有量が１．００％を超えると、その効果が飽和する。また、固溶Ｎｉ量の増加によって、焼入れ処理前に鋼板が過度に硬質化して、刃物等の所定形状に成形する際の加工性が十分に得られなくなる。
　そのため、Ｎｉ含有量は０．１０～１．００％の範囲とする。Ｎｉ含有量は、好ましくは０．１５％以上、より好ましくは０．２０％以上である。また、Ｎｉ含有量は、好ましくは０．８０％以下、より好ましくは０．６０％以下である。 Ni: 0.10 to 1.00%
Ni has the effect of improving corrosion resistance and toughness after quenching. In order to obtain such an effect, the Ni content is set to 0.10% or more. However, when the Ni content exceeds 1.00%, the effect is saturated. Further, due to the increase in the amount of solid-dissolved Ni, the steel sheet becomes excessively hard before the quenching treatment, and it becomes difficult to obtain sufficient workability when forming a predetermined shape of a cutting tool or the like.
Therefore, the Ni content is set in the range of 0.10 to 1.00%. The Ni content is preferably 0.15% or more, more preferably 0.20% or more. The Ni content is preferably 0.80% or less, more preferably 0.60% or less.
Ｎ：０．０１０～０．２００％
　Ｎは、Ｃと同様に、焼入れ処理により得られるマルテンサイト相を硬質化させる効果がある。また、Ｎは、焼入れ処理後の耐食性を向上させる効果もある。このような効果を得るため、Ｎ含有量は０．０１０％以上とする。しかし、Ｎ含有量が０．２００％を超えると、鋳造時に気泡が発生し、表面欠陥の発生を誘引する。
　そのため、Ｎ含有量は０．０１０～０．２００％の範囲とする。Ｎ含有量は好ましくは０．０１５％以上、より好ましくは０．０２０％以上である。また、Ｎ含有量は好ましくは０．１５０％以下、より好ましくは０．１００％以下である。 N: 0.010 to 0.200%
Like C, N has the effect of hardening the martensite phase obtained by quenching. In addition, N also has the effect of improving the corrosion resistance after the quenching treatment. In order to obtain such an effect, the N content is set to 0.010% or more. However, if the N content exceeds 0.200%, bubbles are generated during casting, which induces the occurrence of surface defects.
Therefore, the N content is set in the range of 0.010 to 0.200%. The N content is preferably 0.015% or more, more preferably 0.020% or more. The N content is preferably 0.150% or less, more preferably 0.100% or less.
　以上、本発明の一実施形態に係るステンレス鋼板の基本成分組成について説明したが、さらに、
　Ｍｏ：０．０５～１．００％、Ｃｕ：０．０５～１．００％およびＣｏ：０．０５～０．５０％のうちから選ばれる１種または２種以上、
　ならびに／または、
　Ａｌ：０．００１～０．１００％、Ｔｉ：０．０１～０．１０％、Ｎｂ：０．０１～０．１０％、Ｖ：０．０５～０．５０％、Ｚｒ：０．０１～０．１０％、Ｍｇ：０．０００２～０．００５０％、Ｂ：０．０００２～０．００５０％、Ｃａ：０．０００３～０．００３０％およびＲＥＭ：０．０１～０．１０％のうちから選ばれる１種または２種以上、
を含有させることができる。 The basic composition of the stainless steel sheet according to the embodiment of the present invention has been described above.
One or more selected from Mo: 0.05 to 1.00%, Cu: 0.05 to 1.00% and Co: 0.05 to 0.50%,
And / or
Al: 0.001 to 0.100%, Ti: 0.01 to 0.10%, Nb: 0.01 to 0.10%, V: 0.05 to 0.50%, Zr: 0.01 to Of 0.10%, Mg: 0.0002 to 0.0050%, B: 0.0002 to 0.0050%, Ca: 0.0003 to 0.0030% and REM: 0.01 to 0.10% One or more selected from,
Can be contained.
Ｍｏ：０．０５～１．００％
　Ｍｏは、耐食性を向上させる効果がある。このような効果を得るため、Ｍｏ含有量は０．０５％以上とすることが好ましい。しかし、Ｍｏ含有量が１．００％を超えると、焼入れ処理の加熱・保持時に生成するオーステナイト量が減少し、焼入れ処理後に十分な硬度が得られなくなる。
　そのため、Ｍｏを含有させる場合、Ｍｏ含有量は０．０５～１．００％の範囲とすることが好ましい。Ｍｏ含有量は、より好ましくは０．１０％以上、さらに好ましくは０．５０％以上である。また、Ｍｏ含有量は、より好ましくは０．８０％以下、さらに好ましくは０．６５％以下である。 Mo: 0.05-1.00%
Mo has the effect of improving corrosion resistance. In order to obtain such an effect, the Mo content is preferably 0.05% or more. However, if the Mo content exceeds 1.00%, the amount of austenite produced during heating and holding of the quenching treatment decreases, and sufficient hardness cannot be obtained after the quenching treatment.
Therefore, when Mo is contained, the Mo content is preferably in the range of 0.05 to 1.00%. The Mo content is more preferably 0.10% or more, still more preferably 0.50% or more. The Mo content is more preferably 0.80% or less, still more preferably 0.65% or less.
Ｃｕ：０．０５～１．００％
　Ｃｕは、焼入れ処理後の鋼板において、焼戻し軟化抵抗を向上させる効果がある。このような効果を得るため、Ｃｕ含有量は０．０５％以上とすることが好ましい。しかし、Ｃｕ含有量が１．００％を超えると、耐食性の低下を招く。
　そのため、Ｃｕを含有させる場合、Ｃｕ含有量は０．０５～１．００％の範囲とすることが好ましい。Ｃｕ含有量は、より好ましくは０．１０％以上である。また、Ｃｕ含有量は、より好ましくは０．５０％以下、さらに好ましくは０．２０％以下である。 Cu: 0.05-1.00%
Cu has the effect of improving the temper softening resistance of the hardened steel sheet. In order to obtain such an effect, the Cu content is preferably 0.05% or more. However, if the Cu content exceeds 1.00%, the corrosion resistance is lowered.
Therefore, when Cu is contained, the Cu content is preferably in the range of 0.05 to 1.00%. The Cu content is more preferably 0.10% or more. The Cu content is more preferably 0.50% or less, still more preferably 0.20% or less.
Ｃｏ：０．０５～０．５０％
　Ｃｏは、靭性を向上させる効果がある。このような効果を得るため、Ｃｏ含有量は０．０５％以上とすることが好ましい。しかし、Ｃｏ含有量が０．５０％を超えると、焼入れ処理前に、鋼板を刃物等の所定形状に成形する際の加工性が十分に得られなくなる。
　そのため、Ｃｏを含有させる場合、Ｃｏ含有量は０．０５～０．５０％の範囲とすることが好ましい。Ｃｏ含有量は、より好ましくは０．１０％以上である。また、Ｃｏ含有量は、より好ましくは０．２０％以下である。 Co: 0.05 to 0.50%
Co has the effect of improving toughness. In order to obtain such an effect, the Co content is preferably 0.05% or more. However, if the Co content exceeds 0.50%, the workability when forming the steel sheet into a predetermined shape such as a cutting tool cannot be sufficiently obtained before the quenching treatment.
Therefore, when Co is contained, the Co content is preferably in the range of 0.05 to 0.50%. The Co content is more preferably 0.10% or more. The Co content is more preferably 0.20% or less.
Ａｌ：０．００１～０．１００％
　Ａｌは、Ｓｉと同様に、脱酸剤として作用する。このような効果を得るため、Ａｌ含有量は０．００１％以上とすることが好ましい。しかし、Ａｌ含有量が０．１００％を超えると、焼入れ性が低下する。
　そのため、Ａｌを含有させる場合、Ａｌ含有量は０．００１～０．１００％の範囲とすることが好ましい。Ａｌ含有量は、より好ましくは０．０５０％以下、さらに好ましくは０．０１０％以下である。 Al: 0.001 to 0.100%
Al acts as an antacid in the same manner as Si. In order to obtain such an effect, the Al content is preferably 0.001% or more. However, if the Al content exceeds 0.100%, the hardenability is lowered.
Therefore, when Al is contained, the Al content is preferably in the range of 0.001 to 0.100%. The Al content is more preferably 0.050% or less, still more preferably 0.010% or less.
Ｔｉ：０．０１～０．１０％
　Ｔｉは、Ｃｒと同様、ＣおよびＮとの親和力が高く、鋼中において炭化物を形成する元素である。また、Ｔｉは、焼戻し軟化抵抗を向上させる効果がある。そのため、焼戻しを行った際の軟質化を抑制しつつ、靭性を向上させることが可能となる。このような効果を得るため、Ｔｉ含有量は０．０１％以上とすることが好ましい。しかし、Ｔｉ含有量が０．１０％を超えると、その効果は飽和する。また、却って靭性が低下する。
　そのため、Ｔｉを含有させる場合、Ｔｉ含有量は０．０１～０．１０％の範囲とすることが好ましい。Ｔｉ含有量は、より好ましくは０．０２％以上である。また、Ｔｉ含有量は、より好ましくは０．０５％以下である。 Ti: 0.01 to 0.10%
Like Cr, Ti has a high affinity for C and N and is an element that forms carbides in steel. In addition, Ti has the effect of improving tempering and softening resistance. Therefore, it is possible to improve the toughness while suppressing the softening at the time of tempering. In order to obtain such an effect, the Ti content is preferably 0.01% or more. However, when the Ti content exceeds 0.10%, the effect is saturated. Moreover, the toughness is rather lowered.
Therefore, when Ti is contained, the Ti content is preferably in the range of 0.01 to 0.10%. The Ti content is more preferably 0.02% or more. The Ti content is more preferably 0.05% or less.
Ｎｂ：０．０１～０．１０％
　Ｎｂは、Ｔｉと同様、ＣおよびＮとの親和力が高く、鋼中において炭化物を形成する元素である。また、Ｎｂは、焼戻し軟化抵抗を向上させる効果がある。そのため、焼戻しを行った際の軟質化を抑制しつつ、靭性を向上させることが可能となる。このような効果を得るため、Ｎｂ含有量は０．０１％以上とすることが好ましい。しかし、Ｎｂ含有量が０．１０％を超えると、その効果は飽和する。また、金属間化合物の析出に起因した靭性の低下が生じる場合がある。
　そのため、Ｎｂを含有させる場合、Ｎｂ含有量は０．０１～０．１０％の範囲とすることが好ましい。Ｎｂ含有量は、より好ましくは０．０２％以上である。また、Ｎｂ含有量は、より好ましくは０．０５％以下である。 Nb: 0.01 to 0.10%
Like Ti, Nb is an element that has a high affinity for C and N and forms carbides in steel. In addition, Nb has the effect of improving tempering softening resistance. Therefore, it is possible to improve the toughness while suppressing the softening at the time of tempering. In order to obtain such an effect, the Nb content is preferably 0.01% or more. However, when the Nb content exceeds 0.10%, the effect is saturated. In addition, the toughness may decrease due to the precipitation of intermetallic compounds.
Therefore, when Nb is contained, the Nb content is preferably in the range of 0.01 to 0.10%. The Nb content is more preferably 0.02% or more. The Nb content is more preferably 0.05% or less.
Ｖ：０．０５～０．５０％
　Ｖは、ＴｉやＮｂと同様、ＣおよびＮとの親和力が高く、鋼中において炭化物を形成する元素である。また、Ｖは、焼戻し軟化抵抗を向上させる効果がある。そのため、焼戻しを行った際の軟質化を抑制しつつ、靭性を向上させることが可能となる。このような効果を得るため、Ｖ含有量は０．０５％以上とすることが好ましい。しかし、Ｖ含有量が０．５０％を超えると、その効果は飽和する。また、金属間化合物の析出に起因した靭性の低下が生じる場合がある。
　そのため、Ｖを含有させる場合、Ｖ含有量は０．０５～０．５０％の範囲とすることが好ましい。Ｖ含有量は、より好ましくは０．１０％以上である。また、Ｖ含有量は、より好ましくは０．３０％以下、さらに好ましくは０．２０％以下である。 V: 0.05 to 0.50%
Like Ti and Nb, V has a high affinity for C and N and is an element that forms carbides in steel. Further, V has an effect of improving tempering softening resistance. Therefore, it is possible to improve the toughness while suppressing the softening at the time of tempering. In order to obtain such an effect, the V content is preferably 0.05% or more. However, when the V content exceeds 0.50%, the effect is saturated. In addition, the toughness may decrease due to the precipitation of intermetallic compounds.
Therefore, when V is contained, the V content is preferably in the range of 0.05 to 0.50%. The V content is more preferably 0.10% or more. The V content is more preferably 0.30% or less, still more preferably 0.20% or less.
Ｚｒ：０．０１～０．１０％
　Ｚｒは、ＴｉやＮｂと同様、ＣおよびＮとの親和力が高く、鋼中において炭化物を形成する元素である。また、Ｚｒは、焼戻し軟化抵抗を向上させる効果がある。そのため、焼戻しを行った際の軟質化を抑制しつつ、靭性を向上させることが可能となる。このような効果を得るため、Ｚｒ含有量は０．０１％以上とすることが好ましい。しかし、Ｚｒ含有量が０．１０％を超えると、その効果は飽和する。また、金属間化合物の析出に起因した靭性の低下が生じる場合がある。
　そのため、Ｚｒを含有させる場合、Ｚｒ含有量は０．０１～０．１０％の範囲とすることが好ましい。Ｚｒ含有量は、より好ましくは０．０２％以上である。また、Ｚｒ含有量は、より好ましくは０．０５％以下である。 Zr: 0.01-0.10%
Like Ti and Nb, Zr has a high affinity for C and N and is an element that forms carbides in steel. In addition, Zr has the effect of improving tempering and softening resistance. Therefore, it is possible to improve the toughness while suppressing the softening at the time of tempering. In order to obtain such an effect, the Zr content is preferably 0.01% or more. However, when the Zr content exceeds 0.10%, the effect is saturated. In addition, the toughness may decrease due to the precipitation of intermetallic compounds.
Therefore, when Zr is contained, the Zr content is preferably in the range of 0.01 to 0.10%. The Zr content is more preferably 0.02% or more. The Zr content is more preferably 0.05% or less.
Ｍｇ：０．０００２～０．００５０％
　Ｍｇは、スラブの等軸晶率を向上させ、加工性や靭性を向上させる効果がある。このような効果を得るため、Ｍｇ含有量は０．０００２％以上とすることが好ましい。しかし、Ｍｇ含有量が０．００５０％を超えると、鋼板の表面性状が悪化する場合がある。
　そのため、Ｍｇを含有させる場合、Ｍｇ含有量は０．０００２～０．００５０％の範囲とすることが好ましい。Ｍｇ含有量は、より好ましくは０．００１０％以上である。また、Ｍｇ含有量は、より好ましくは０．００２０％以下である。 Mg: 0.0002 to 0.0050%
Mg has the effect of improving the equiaxed crystal ratio of the slab and improving workability and toughness. In order to obtain such an effect, the Mg content is preferably 0.0002% or more. However, if the Mg content exceeds 0.0050%, the surface properties of the steel sheet may deteriorate.
Therefore, when Mg is contained, the Mg content is preferably in the range of 0.0002 to 0.0050%. The Mg content is more preferably 0.0010% or more. The Mg content is more preferably 0.0020% or less.
Ｂ：０．０００２～０．００５０％
　Ｂは、鋳造および熱間圧延時の熱間加工性の向上させる効果がある。また、Ｂは、フェライト相およびオーステナイト相の粒界に偏析して粒界強度を上昇させる。これによって、鋳造および熱間圧延時の割れの発生を抑制する。このような効果を得るため、Ｂ含有量は０．０００２％以上とすることが好ましい。しかし、Ｂ含有量が０．００５０％を超えると、焼入れ処理前に、鋼板を刃物等の所定形状に成形する際の加工性が十分に得られなくなる。また、靭性の低下を招く。
　そのため、Ｂを含有させる場合、Ｂ含有量は０．０００２～０．００５０％の範囲とすることが好ましい。Ｂ含有量は、より好ましくは０．０００５％以上である。また、Ｂ含有量は、より好ましくは０．００３０％以下、さらに好ましくは０．００２０％以下である。 B: 0.0002 to 0.0050%
B has the effect of improving hot workability during casting and hot rolling. Further, B segregates at the grain boundaries of the ferrite phase and the austenite phase to increase the grain boundary strength. This suppresses the occurrence of cracks during casting and hot rolling. In order to obtain such an effect, the B content is preferably 0.0002% or more. However, if the B content exceeds 0.0050%, the workability when forming the steel sheet into a predetermined shape such as a cutting tool cannot be sufficiently obtained before the quenching treatment. It also causes a decrease in toughness.
Therefore, when B is contained, the B content is preferably in the range of 0.0002 to 0.0050%. The B content is more preferably 0.0005% or more. The B content is more preferably 0.0030% or less, still more preferably 0.0020% or less.
Ｃａ：０．０００３～０．００３０％
　Ｃａは、製錬ならびに連続鋳造時に生成する介在物を微細化する効果があり、特に連続鋳造におけるノズルの閉塞を防止するのに有効である。このような効果を得るため、Ｃａ含有量は０．０００３％以上とすることが好ましい。しかし、Ｃａ含有量が０．００３０％を超えると、ＣａＳの生成により耐食性が低下する場合がある。
　そのため、Ｃａを含有する場合、Ｃａ含有量は０．０００３～０．００３０％の範囲とすることが好ましい。Ｃａ含有量は、より好ましくは０．０００５％以上、さらに好ましくは０．０００７％以上である。また、Ｃａ含有量は、より好ましくは０．００２０％以下、さらに好ましくは０．００１５％以下である。 Ca: 0.0003 to 0.0030%
Ca has the effect of miniaturizing inclusions generated during smelting and continuous casting, and is particularly effective in preventing nozzle blockage during continuous casting. In order to obtain such an effect, the Ca content is preferably 0.0003% or more. However, if the Ca content exceeds 0.0030%, the corrosion resistance may decrease due to the formation of CaS.
Therefore, when Ca is contained, the Ca content is preferably in the range of 0.0003 to 0.0030%. The Ca content is more preferably 0.0005% or more, still more preferably 0.0007% or more. The Ca content is more preferably 0.0020% or less, still more preferably 0.0015% or less.
ＲＥＭ：０．０１～０．１０％
　ＲＥＭ（Ｒａｒｅ　Ｅａｒｔｈ　Ｍｅｔａｌｓ：希土類金属）は、熱間延性を向上させる効果がある。また、ＲＥＭは、熱間圧延時の鋼板端面部の割れや肌荒れを抑制する効果もある。このような効果を得るため、ＲＥＭ含有量は０．０１％以上とすることが好ましい。しかし、ＲＥＭ含有量が０．１０％を超えると、その効果は飽和する。また、ＲＥＭは、高価な元素でもある。
　そのため、ＲＥＭを含有する場合、ＲＥＭ含有量は０．０１～０．１０％の範囲とすることが好ましい。ＲＥＭ含有量は、より好ましくは０．０５％以下である。 REM: 0.01 to 0.10%
REM (Rare Earth Metals) has the effect of improving hot ductility. In addition, REM also has an effect of suppressing cracking and rough skin of the end face portion of the steel sheet during hot rolling. In order to obtain such an effect, the REM content is preferably 0.01% or more. However, when the REM content exceeds 0.10%, the effect is saturated. REM is also an expensive element.
Therefore, when REM is contained, the REM content is preferably in the range of 0.01 to 0.10%. The REM content is more preferably 0.05% or less.
　上記以外の成分の残部は、Ｆｅおよび不可避的不純物である。 The rest of the components other than the above are Fe and unavoidable impurities.
　次に、本発明の一実施形態に係るステンレス鋼板の金属組織について、説明する。
　本発明の一実施形態に係るステンレス鋼板の金属組織は、焼入れ処理の前後で、主体となる組織が変化する。
　例えば、本発明の一実施形態に係るステンレス鋼板を製品に加工する場合、まず、鋼板が硬質化していない段階で、鋼板にプレス加工等により所定形状にブランキングまたは鍛造加工する。ついで、所定形状に加工した鋼板に、焼入れ処理、または、焼入れおよび焼戻し処理を施して硬質化させる。すなわち、焼入れ処理の前後で、主体となる組織を変化させる、具体的には、フェライト相からマルテンサイト相に変化させる。
　ただし、粒径：２．０μｍ以上のＣｒ系炭化物は、焼入れ処理の前後でもあまり変わらず、ほぼ維持される。
　よって、本発明の一実施形態に係るステンレス鋼板の金属組織では、焼入れ処理前後を問わず、粒径：２．０μｍ以上のＣｒ系炭化物の体積率を１０％以下とすることが極めて重要となる。 Next, the metal structure of the stainless steel sheet according to the embodiment of the present invention will be described.
The metal structure of the stainless steel sheet according to the embodiment of the present invention changes its main structure before and after the quenching treatment.
For example, when processing a stainless steel sheet according to an embodiment of the present invention into a product, first, the steel sheet is blanked or forged into a predetermined shape by press working or the like at a stage where the steel sheet is not hardened. Then, the steel sheet processed into a predetermined shape is hardened by quenching or quenching and tempering. That is, before and after the quenching treatment, the main structure is changed, specifically, the ferrite phase is changed to the martensite phase.
However, Cr-based carbides having a particle size of 2.0 μm or more do not change much before and after the quenching treatment and are almost maintained.
Therefore, in the metal structure of the stainless steel sheet according to the embodiment of the present invention, it is extremely important that the volume ratio of Cr-based carbide having a particle size of 2.0 μm or more is 10% or less regardless of before and after the quenching treatment. ..
粒径：２．０μｍ以上のＣｒ系炭化物の体積率：１０％以下
　Ｃｒ系炭化物は、ステンレス鋼板の母材（焼入れ前後ともに）よりも硬質である。そのため、金属組織に粗大なＣｒ系炭化物、特に、粒径：２．０μｍ以上のＣｒ系炭化物が多量に存在している状態で研磨や刃付け加工等を施すと、当該Ｃｒ系炭化物が存在している部位では、他の部位に比べて、研磨量が少なくなる。その結果、研磨後に、局所的に凸部が生じ、これらが筋模様として顕在化する。
　そのため、粒径：２．０μｍ以上のＣｒ系炭化物の体積率は、１０％以下とする。粒径：２．０μｍ以上のＣｒ系炭化物の体積率は、好ましくは５％以下、より好ましくは２％以下である。なお、粒径：２．０μｍ以上のＣｒ系炭化物の体積率は０％であってもよい。
　なお、粒径２．０μｍ未満のＣｒ系炭化物については、研磨時に肉眼で識別できるほどの凹凸を生じさせず、筋模様の発生には関与しない。そのため、粒径２．０μｍ未満のＣｒ系炭化物の体積率は、特に限定されない。 Particle size: 2.0 μm or more Volume fraction of Cr-based carbide: 10% or less Cr-based carbide is harder than the base material of stainless steel sheet (both before and after quenching). Therefore, if polishing or cutting is performed in a state where a large amount of coarse Cr-based carbides, particularly Cr-based carbides having a particle size of 2.0 μm or more, are present in the metal structure, the Cr-based carbides are present. The amount of polishing is smaller in the area where the material is used than in other areas. As a result, after polishing, convex portions are locally generated, and these are manifested as streaks.
Therefore, the volume fraction of Cr-based carbide having a particle size of 2.0 μm or more is set to 10% or less. The volume ratio of Cr-based carbide having a particle size of 2.0 μm or more is preferably 5% or less, more preferably 2% or less. The volume fraction of Cr-based carbide having a particle size of 2.0 μm or more may be 0%.
For Cr-based carbides having a particle size of less than 2.0 μm, irregularities that can be discerned with the naked eye are not generated during polishing and are not involved in the generation of streaks. Therefore, the volume fraction of Cr-based carbide having a particle size of less than 2.0 μm is not particularly limited.
　また、ここでいうＣｒ系炭化物は、主にＣｒ23Ｃ6である。また、Ｃｒ炭化物における一部のＣｒが、ＦｅやＭｎ、Ｔｉ、Ｎｂ、Ｖ、Ｚｒ等の元素に置換されたものや、一部のＣがＮに置換されたものも、ここでいうＣｒ系炭化物に含むものとする。 The Cr-based carbide referred to here is mainly Cr 23 C 6 . Further, a part of Cr in the Cr carbide is replaced with an element such as Fe, Mn, Ti, Nb, V, Zr, and a part of C is replaced with N. It shall be contained in carbide.
　また、本発明の一実施形態に係るステンレス鋼板におけるＣｒ系炭化物以外の組織は、フェライト相とマルテンサイト相の合計の体積率が９５％以上、より好ましくは９８％以上である金属組織となる。フェライト相とマルテンサイト相の合計の体積率は１００％であってもよい。フェライト相、マルテンサイト相および上記したＣｒ系炭化物以外の残部組織としては、残留オーステナイト相やその他の析出物（粒径：２．０μｍ未満のＣｒ系炭化物も含む）、介在物（例えば、ＡｌやＳｉ等の酸化物およびＭｎ等の硫化物など）が挙げられる。残部組織の体積率は、好ましくは５％以下、より好ましくは２％以下である。残部組織の体積率は０％であってもよい。
　なお、本発明の一実施形態に係るステンレス鋼板には、焼入れ処理前後の両方の鋼板が含まれ、例えば、熱延鋼板、熱延焼鈍鋼板、冷延鋼板および冷延焼鈍鋼板、ならびに、これらの鋼板に、焼入れ処理および／または焼戻し処理を施した鋼板（後述する焼入れ処理鋼板および焼戻し処理鋼板）等が含まれる。 Further, the structure other than the Cr-based carbide in the stainless steel sheet according to the embodiment of the present invention is a metal structure in which the total volume ratio of the ferrite phase and the martensite phase is 95% or more, more preferably 98% or more. The total volume fraction of the ferrite phase and the martensite phase may be 100%. Remaining structures other than the ferrite phase, martensite phase and the above-mentioned Cr-based carbides include residual austenite phase and other precipitates (including Cr-based carbides having a particle size of less than 2.0 μm), inclusions (for example, Al and the like. Oxides such as Si and sulfides such as Mn) can be mentioned. The volume fraction of the residual tissue is preferably 5% or less, more preferably 2% or less. The volume fraction of the residual tissue may be 0%.
The stainless steel sheet according to the embodiment of the present invention includes both steel sheets before and after the quenching treatment, for example, a hot-rolled steel sheet, a hot-rolled tempered steel sheet, a cold-rolled steel sheet, a cold-rolled fired steel sheet, and these. The steel sheet includes a steel sheet that has undergone a quenching treatment and / or a tempering treatment (a hardened steel sheet and a tempered steel sheet described later) and the like.
　なお、熱延鋼板、熱延焼鈍鋼板、冷延鋼板および冷延焼鈍鋼板の段階では、Ｃｒ系炭化物以外の組織は、フェライト相主体の組織となる。
　具体的には、フェライト相が体積率で８０％以上、好ましくは９０％以上、より好ましくは９５％以上、さらに好ましくは９８％以上である金属組織となる。フェライト相の体積率が１００％であってもよい。フェライト相および上記したＣｒ系炭化物以外の残部組織としては、マルテンサイト相や残留オーステナイト相、その他の析出物（粒径：２．０μｍ未満のＣｒ系炭化物も含む）、介在物（例えば、ＡｌやＳｉ等の酸化物およびＭｎ等の硫化物など）が挙げられる。残部組織の体積率は、好ましくは２０％以下、より好ましくは１０％以下、さらに好ましくは５％以下、よりさらに好ましくは２％以下である。残部組織の体積率は０％であってもよい。
　なお、熱延鋼板には、熱延ままの鋼板に加え、熱延ままの鋼板に酸洗等の酸化スケールの除去処理を施して得た鋼板が含まれる。また、熱延焼鈍鋼板には、熱延鋼板に熱延板焼鈍を施して得た鋼板に加え、該熱延板焼鈍を施して得た鋼板にさらに酸洗等の酸化スケールの除去処理を施して得た鋼板が含まれる。冷延鋼板には、冷延ままの鋼板に加え、冷延ままの鋼板に酸洗等の酸化スケールの除去処理を施して得た鋼板が含まれる。 At the stage of the hot-rolled steel sheet, the hot-rolled annealed steel sheet, the cold-rolled steel sheet, and the cold-rolled annealed steel sheet, the structure other than the Cr-based carbide is mainly a ferrite phase structure.
Specifically, the ferrite phase has a metal structure having a volume fraction of 80% or more, preferably 90% or more, more preferably 95% or more, and further preferably 98% or more. The volume fraction of the ferrite phase may be 100%. As the residual structure other than the ferrite phase and the above-mentioned Cr-based carbides, martensite phase, retained austenite phase, other precipitates (including Cr-based carbides having a particle size of less than 2.0 μm), inclusions (for example, Al and Oxides such as Si and sulfides such as Mn) can be mentioned. The volume fraction of the residual tissue is preferably 20% or less, more preferably 10% or less, still more preferably 5% or less, still more preferably 2% or less. The volume fraction of the residual tissue may be 0%.
The hot-rolled steel sheet includes, in addition to the hot-rolled steel sheet, a steel sheet obtained by subjecting the hot-rolled steel sheet to an oxidation scale removal treatment such as pickling. Further, for the hot-rolled hardened steel sheet, in addition to the steel sheet obtained by hot-rolling the hot-rolled steel sheet, the steel sheet obtained by hot-rolling the hardened steel sheet is further subjected to an oxidation scale removal treatment such as pickling. The steel sheet obtained from the above is included. The cold-rolled steel sheet includes, in addition to the cold-rolled steel sheet, a steel sheet obtained by subjecting the cold-rolled steel sheet to an oxidation scale removal treatment such as pickling.
　さらに、熱延鋼板、熱延焼鈍鋼板、冷延鋼板および冷延焼鈍鋼板に焼入れ処理を施した鋼板（以下、焼入れ処理鋼板ともいう）では、Ｃｒ系炭化物以外の組織は、マルテンサイト相主体の組織となる。
　具体的には、マルテンサイト相が体積率で８０％以上、好ましくは９０％以上、より好ましくは９５％以上、さらに好ましくは９８％以上である金属組織となる。マルテンサイト相の体積率が１００％であってもよい。マルテンサイト相および上記したＣｒ系炭化物以外の残部組織としては、フェライト相や残留オーステナイト相、その他の析出物（粒径：２．０μｍ未満のＣｒ系炭化物も含む）、介在物（例えば、ＡｌやＳｉ等の酸化物およびＭｎ等の硫化物など）が挙げられる。残部組織の体積率は、好ましくは２０％以下、より好ましくは１０％以下、さらに好ましくは５％以下、よりさらに好ましくは２％以下である。残部組織の体積率は０％であってもよい。
　なお、焼入れ処理により硬質化されるので、焼入れ処理鋼板では、ロックウェル硬度がＨＲＣ５５以上となる。 Further, in a hot-rolled steel sheet, a hot-rolled tempered steel sheet, a cold-rolled steel sheet, and a steel sheet obtained by quenching a cold-rolled hardened steel sheet (hereinafter, also referred to as a hardened steel sheet), the structure other than Cr-based carbides is mainly composed of martensite phase. Become an organization.
Specifically, the martensite phase has a metal structure having a volume fraction of 80% or more, preferably 90% or more, more preferably 95% or more, and further preferably 98% or more. The volume fraction of the martensite phase may be 100%. The residual structure other than the martensite phase and the above-mentioned Cr-based carbides includes a ferrite phase, a retained austenite phase, other precipitates (including Cr-based carbides having a particle size of less than 2.0 μm), inclusions (for example, Al and the like. Oxides such as Si and sulfides such as Mn) can be mentioned. The volume fraction of the residual tissue is preferably 20% or less, more preferably 10% or less, still more preferably 5% or less, still more preferably 2% or less. The volume fraction of the residual tissue may be 0%.
Since it is hardened by the quenching treatment, the rockwell hardness of the hardened steel sheet is HRC55 or more.
　加えて、焼入れ処理鋼板に焼戻し処理を施した鋼板（以下、焼戻し処理鋼板ともいう）では、Ｃｒ系炭化物以外の組織は、焼入れ処理後に比べて転位密度ならびに固溶Ｃ、Ｎが減少したマルテンサイト相（焼戻しマルテンサイト相と呼称される場合がある）主体の組織となり、焼戻し処理前のマルテンサイト分率がほぼ維持される。
　具体的には、マルテンサイト相が体積率で８０％以上、好ましくは９０％以上、より好ましくは９５％以上、さらに好ましくは９８％以上である金属組織となる。また、フェライト相が体積率で２０％以下、好ましくは１０％以下、より好ましくは５％以下、さらに好ましくは２％以下である金属組織となる。フェライト相、マルテンサイト相および上記したＣｒ系炭化物以外の残部組織としては、残留オーステナイト相やその他の析出物（粒径：２．０μｍ未満のＣｒ系炭化物も含む）、介在物（例えば、ＡｌやＳｉ等の酸化物およびＭｎ等の硫化物など）が挙げられる。残部組織の体積率は、好ましくは５％以下、より好ましくは２％以下である。
　ここで、焼戻し処理は、焼入れ処理により硬質化した鋼板の硬度および耐久性を調整するために行われるものであり、焼戻し処理前の焼入れ処理鋼板に比べると、焼戻し処理鋼板では硬度が低下する。具体的には、焼戻し処理鋼板では、ロックウェル硬度がＨＲＣ４０～５０になる。 In addition, in a steel plate obtained by tempering a hardened steel plate (hereinafter, also referred to as a tempered steel plate), martensite having a reduced dislocation density and solid-melt C and N in structures other than Cr-based carbides as compared with that after the quenching treatment. The structure is mainly phase (sometimes called tempered martensite phase), and the martensite fraction before tempering is almost maintained.
Specifically, the martensite phase has a metal structure having a volume fraction of 80% or more, preferably 90% or more, more preferably 95% or more, and further preferably 98% or more. Further, the ferrite phase has a metal structure having a volume fraction of 20% or less, preferably 10% or less, more preferably 5% or less, still more preferably 2% or less. Remaining structures other than the ferrite phase, martensite phase and the above-mentioned Cr-based carbides include residual austenite phase and other precipitates (including Cr-based carbides having a particle size of less than 2.0 μm), inclusions (for example, Al and the like. Oxides such as Si and sulfides such as Mn) can be mentioned. The volume fraction of the residual tissue is preferably 5% or less, more preferably 2% or less.
Here, the tempering treatment is performed to adjust the hardness and durability of the steel sheet hardened by the quenching treatment, and the hardness of the tempered steel sheet is lower than that of the tempered steel sheet before the tempering treatment. Specifically, in the tempered steel sheet, the Rockwell hardness is HRC40 to 50.
　また、粒径：２．０μｍ以上のＣｒ系炭化物の体積率は以下のようにして測定する。
　すなわち、供試材となる鋼板の板幅中央部から組織観察用の試験片を採取する。ついで、試験片の圧延方向断面を鏡面研磨後、ピクリン酸塩酸水溶液を用いてエッチングを行い、倍率：５００倍の光学顕微鏡写真を１０視野撮影する。得られた組織写真中のＣｒ系炭化物の面積を画像解析により測定し、円相当直径が２．０μｍ以上になるＣｒ系炭化物を特定する。そして、特定した円相当直径が２．０μｍ以上になるＣｒ系炭化物の合計の面積率を算出し、その算出した値を、粒径：２．０μｍ以上のＣｒ系炭化物の体積率とする。
　ここで、上記の画像解析では、組織写真のデジタルデータについて、画像解析装置を用いて、コントラスト差により母相（フェライト相またはマルテンサイト相）の粒界と析出物の境界を自動検出させる（粒界および境界は線状の黒いコントラストを呈し、結晶粒は比較的明るいコントラストを呈する）。次に、母相と析出物の境界線に囲われた領域を、析出物とし、各析出物の領域の面積を自動測定する。その後、後述する方法によりＣｒ系炭化物と同定した析出物について、面積が３．１４μｍ2以上（すなわち、円相当直径が２.０μｍ以上）となるもののみを特定する。そして、特定した析出物の合計の面積を算出する。
　そして、（円相当直径：２.０μｍ以上の析出物（Ｃｒ系炭化物）の合計の面積）÷（組織写真の全面積）×１００［％］を求め、求めた値を粒径：２．０μｍ以上のＣｒ系炭化物の体積率とする。 The volume fraction of Cr-based carbide having a particle size of 2.0 μm or more is measured as follows.
That is, a test piece for observing the structure is collected from the central portion of the width of the steel plate used as the test material. Then, the cross section of the test piece in the rolling direction is mirror-polished and then etched with an aqueous solution of hydrochloric acid picrinate, and an optical micrograph with a magnification of 500 times is taken in 10 fields. The area of Cr-based carbides in the obtained microstructure photograph is measured by image analysis, and Cr-based carbides having a circle-equivalent diameter of 2.0 μm or more are identified. Then, the total area ratio of the Cr-based carbides having the specified circle-equivalent diameter of 2.0 μm or more is calculated, and the calculated value is used as the volume fraction of the Cr-based carbides having a particle size of 2.0 μm or more.
Here, in the above image analysis, the boundary between the grain boundary of the matrix phase (ferrite phase or martensite phase) and the boundary of the precipitate is automatically detected by the contrast difference using the image analysis device for the digital data of the tissue photograph (grains). The boundaries and boundaries have a linear black contrast, and the grains have a relatively bright contrast). Next, the region surrounded by the boundary line between the matrix and the precipitate is designated as a precipitate, and the area of each precipitate region is automatically measured. Then, among the precipitates identified as Cr-based carbides by the method described later, only those having an area of 3.14 μm 2 or more (that is, a circle-equivalent diameter of 2.0 μm or more) are specified. Then, the total area of the specified precipitates is calculated.
Then, (total area of precipitates (Cr-based carbides) having a diameter equivalent to a circle: 2.0 μm or more) ÷ (total area of microstructure photograph) × 100 [%] was calculated, and the obtained value was calculated as particle size: 2.0 μm. Let it be the volume fraction of the above Cr-based carbides.
　また、上記の組織写真中の析出物がＣｒ系炭化物であることの同定は、以下のようにして行う。
　すなわち、上記の組織写真を撮影した同じ視野において、ＳＥＭ－ＥＤＳ（Ｓｃａｎｎｉｎｇ　Ｅｌｅｃｔｒｏｎ　Ｍｉｃｒｏｓｃｏｐｅ－Ｅｎｅｒｇｙ　Ｄｉｓｐｅｒｓｉｖｅ　Ｘ－ｒａｙ　Ｓｐｅｃｔｒｏｓｃｏｐｙ）を用いた点分析を行い、観察される析出物の主成分を測定する。
　具体的には、析出物におけるＣｒおよびＦｅの合計の含有量が６０質量％以上であり、かつ、析出物におけるＦｅおよびＣｒの合計の含有量に対する、析出物におけるＣｒ含有量の比（［Ｃｒ含有量（質量％）］／（［Ｆｅ含有量（質量％）］＋［Ｃｒ含有量（質量％）］）が０．４以上となる場合、当該析出物をＣｒ系炭化物と同定する。 Further, the identification that the precipitate in the above-mentioned microstructure photograph is a Cr-based carbide is performed as follows.
That is, point analysis using SEM-EDS (Scanning Electron Microscope-Energy Dispersive X-ray Spectroscopy) is performed in the same field of view in which the above-mentioned tissue photograph is taken, and the main component of the observed precipitate is measured.
Specifically, the total content of Cr and Fe in the precipitate is 60% by mass or more, and the ratio of the Cr content in the precipitate to the total content of Fe and Cr in the precipitate ([Cr. When the content (% by mass)] / ([Fe content (% by mass)] + [Cr content (% by mass)]) is 0.4 or more, the precipitate is identified as a Cr-based carbide.
　加えて、フェライト相およびマルテンサイト相の体積率は、以下のようにして求める。
　すなわち、上記の組織写真において、組織形状とエッチング強度からマルテンサイト相とフェライト相とを区別する（なお、マルテンサイト相はフェライト相よりより深くエッチングされる。そのため、マルテンサイト相は、フェライト相よりコントラストが暗い。）。ついで、画像処理により、視野ごとにフェライト相およびマルテンサイト相の体積率をそれぞれ算出する。ついで、視野ごとに得られたフェライト相およびマルテンサイト相の体積率の算術平均値を算出し、その値をフェライト相およびマルテンサイト相の体積率とする。 In addition, the volume fractions of the ferrite phase and the martensite phase are determined as follows.
That is, in the above microstructure photograph, the martensite phase and the ferrite phase are distinguished from each other based on the structure shape and the etching strength (note that the martensite phase is etched deeper than the ferrite phase. Therefore, the martensite phase is more than the ferrite phase. The contrast is dark.). Then, the volume fractions of the ferrite phase and the martensite phase are calculated for each field of view by image processing. Then, the arithmetic mean value of the volume ratios of the ferrite phase and the martensite phase obtained for each field of view is calculated, and the value is used as the volume ratio of the ferrite phase and the martensite phase.
　なお、本発明の一実施形態に係るステンレス鋼板の厚みは特に限定されるものではないが、包丁やカミソリ、医療用刃物等へ適用する観点から、０．１～５．０ｍｍとすることが好適である。本発明の一実施形態に係るステンレス鋼板の厚みは、より好ましくは０．５ｍｍ以上、さらに好ましくは１．０ｍｍ以上である。また、本発明の一実施形態に係るステンレス鋼板の厚みは、より好ましくは４．０ｍｍ以下、さらに好ましくは２．５ｍｍ以下である。 The thickness of the stainless steel plate according to the embodiment of the present invention is not particularly limited, but is preferably 0.1 to 5.0 mm from the viewpoint of application to kitchen knives, razors, medical blades, and the like. Is. The thickness of the stainless steel sheet according to the embodiment of the present invention is more preferably 0.5 mm or more, still more preferably 1.0 mm or more. The thickness of the stainless steel sheet according to the embodiment of the present invention is more preferably 4.0 mm or less, still more preferably 2.5 mm or less.
　次に、本発明の一実施形態に係るステンレス鋼板の製造方法について説明する。
　すなわち、転炉または電気炉等の溶解炉で溶鋼を溶製する。ついで、該溶鋼に、取鍋精錬または真空精錬による二次精錬を施して、上記の成分組成に調整する。ついで、該溶鋼を、連続鋳造法または造塊－分塊圧延法等により、鋼素材（鋼スラブ）とする。 Next, a method for manufacturing a stainless steel sheet according to an embodiment of the present invention will be described.
That is, molten steel is melted in a melting furnace such as a converter or an electric furnace. Then, the molten steel is subjected to secondary refining by ladle refining or vacuum refining to adjust to the above-mentioned composition. Then, the molten steel is made into a steel material (steel slab) by a continuous casting method, an ingot-bulk rolling method, or the like.
・第１の工程（鋼スラブ加熱工程）
　そして、第１の工程として、上記の鋼スラブを１２００～１３５０℃で３０分以上保持する。 ・ First process (steel slab heating process)
Then, as the first step, the steel slab is held at 1200 to 1350 ° C. for 30 minutes or more.
鋼スラブを１２００～１３５０℃で３０分以上保持
　熱間圧延前に行う鋼スラブの加熱では、鋳造時に鋼スラブ断面の柱状晶と等軸晶との境界部近傍の鋳造方向に沿って生成した粗大なＣｒ系炭化物を、極力、オーステナイト相へ固溶させる必要がある。
　ここで、鋼スラブの保持温度（以下、スラブ加熱温度ともいう）が、１２００℃未満では、オーステナイト相へのＣｒ系炭化物の固溶が十分に促進されない。そのため、粗大なＣｒ系炭化物の生成が十分に抑制されず、良好な表面品質が得られない。一方、スラブ加熱温度が１３５０℃を超えると、鋼スラブの金属組織がオーステナイト相とデルタフェライト相の二相組織、または、デルタフェライトの単相組織となって、オーステナイト相へのＣｒ系炭化物の固溶が十分に促進されない。そのため、粗大なＣｒ系炭化物の生成が十分に抑制されず、良好な表面品質が得られない。
　よって、スラブ加熱温度は１２００～１３５０℃の範囲とする。スラブ加熱温度は、好ましくは１３００℃以下、より好ましくは１２５０℃以下である。 Hold the steel slab at 1200 to 1350 ° C. for 30 minutes or more In the heating of the steel slab performed before hot rolling, the coarseness generated along the casting direction near the boundary between the columnar crystal and the equiaxed crystal of the steel slab cross section during casting. Cr-based carbides need to be dissolved in the austenite phase as much as possible.
Here, if the holding temperature of the steel slab (hereinafter, also referred to as the slab heating temperature) is less than 1200 ° C., the solid solution of Cr-based carbides in the austenite phase is not sufficiently promoted. Therefore, the formation of coarse Cr-based carbides is not sufficiently suppressed, and good surface quality cannot be obtained. On the other hand, when the slab heating temperature exceeds 1350 ° C., the metal structure of the steel slab becomes a two-phase structure of an austenite phase and a delta ferrite phase or a single phase structure of delta ferrite, and the Cr-based carbide solid solution to the austenite phase. Melting is not sufficiently promoted. Therefore, the formation of coarse Cr-based carbides is not sufficiently suppressed, and good surface quality cannot be obtained.
Therefore, the slab heating temperature is in the range of 1200 to 1350 ° C. The slab heating temperature is preferably 1300 ° C. or lower, more preferably 1250 ° C. or lower.
　また、１２００～１３５０℃での保持時間が３０分未満の場合、やはりオーステナイト相へのＣｒ系炭化物の固溶が不十分となる。そのため、粗大なＣｒ系炭化物の生成が十分に抑制されず、良好な表面品質が得られない。
　よって、１２００～１３５０℃での保持時間は３０分以上とする。
　なお、当該保持時間が２４時間を超えると、鋼スラブの加熱中に生成する酸化スケールが厚くなって、表面欠陥が発生しやすくなる。また、生産性も低下する。そのため、当該保持時間は２４時間以下とすることが好ましい。当該保持時間は、より好ましくは１２時間以下、さらに好ましくは３時間以下である。 Further, when the holding time at 1200 to 1350 ° C. is less than 30 minutes, the solid solution of the Cr-based carbide in the austenite phase is also insufficient. Therefore, the formation of coarse Cr-based carbides is not sufficiently suppressed, and good surface quality cannot be obtained.
Therefore, the holding time at 1200 to 1350 ° C. is set to 30 minutes or more.
If the holding time exceeds 24 hours, the oxide scale generated during heating of the steel slab becomes thick, and surface defects are likely to occur. In addition, productivity is also reduced. Therefore, the holding time is preferably 24 hours or less. The holding time is more preferably 12 hours or less, still more preferably 3 hours or less.
・第２の工程：熱間圧延工程
　ついで、第２の工程として、該鋼スラブに、熱間圧延を施して熱延鋼板とし、該熱延鋼板を巻取る。
　この際、熱間圧延における圧延パスのうち、終了温度：１０５０℃以上で、かつ、圧下率：２０％以上となる圧延パス数を３パス以上とし、また、熱延鋼板の巻取り温度を６００℃以上とすることが重要である。 -Second step: Hot rolling step Then, as a second step, the steel slab is hot-rolled to obtain a hot-rolled steel sheet, and the hot-rolled steel sheet is wound up.
At this time, among the rolling passes in hot rolling, the number of rolling passes having an end temperature of 1050 ° C. or higher and a reduction rate of 20% or higher is 3 or more, and the winding temperature of the hot-rolled steel sheet is 600. It is important to keep the temperature above ℃.
熱間圧延における圧延パスのうち、終了温度：１０５０℃以上で、かつ、圧下率：２０％以上となる圧延パス数：３パス以上
　熱間圧延では、Ｃｒ系炭化物のオーステナイト相への固溶をさらに促進して、鋼スラブ加熱後に残留する粗大なＣｒ系炭化物を解消する。また、オーステナイト相の動的再結晶および／または静的再結晶の促進により、オーステナイト相の結晶粒が微細化される。これにより、その後の熱延鋼板の巻取りの際に、オーステナイト相の粒界から析出するＣｒ系炭化物の析出サイトが増加して、再析出するＣｒ系炭化物も微細化される。
　特に、１０５０℃以上の温度で圧延を施すことにより、オーステナイト相の動的再結晶および／または静的再結晶が有効に促進される。また、圧延パスごとの圧下率を２０％以上とすることにより、圧延ひずみが鋼スラブの板厚中央部まで効果的に付与される。これにより、鋼スラブの柱状晶と等軸晶の境界部近傍において鋳造方向に沿って生成した粗大なＣｒ系炭化物がより有効に解消される。
　そのため、熱間圧延における圧延パスのうち、終了温度：１０５０℃以上で、かつ、圧下率：２０％以上となる圧延パス数（以下、所定条件を満足する圧延パスともいう）を３パス以上とする必要がある。
　なお、所定条件を満足する圧延パス数の上限は特に限定されるものではないが、過度に増加すると圧延温度の維持に多大な入熱が必要となって製造コストの増加を招く、よって、所定条件を満足する圧延パス数は、１０パス以下とすることが好ましい。 Of the rolling passes in hot rolling, the number of rolling passes at which the end temperature is 1050 ° C or higher and the rolling reduction is 20% or higher: 3 passes or higher. Further accelerating, the coarse Cr-based carbides remaining after heating the steel slab are eliminated. Further, by promoting dynamic recrystallization and / or static recrystallization of the austenite phase, the crystal grains of the austenite phase are refined. As a result, the precipitation sites of Cr-based carbides precipitated from the grain boundaries of the austenite phase increase during the subsequent winding of the hot-rolled steel sheet, and the Cr-based carbides reprecipitated are also refined.
In particular, rolling at a temperature of 1050 ° C. or higher effectively promotes dynamic recrystallization and / or static recrystallization of the austenite phase. Further, by setting the rolling reduction ratio for each rolling pass to 20% or more, rolling strain is effectively applied to the central portion of the plate thickness of the steel slab. As a result, coarse Cr-based carbides formed along the casting direction in the vicinity of the boundary between the columnar crystal and the equiaxed crystal of the steel slab are more effectively eliminated.
Therefore, among the rolling passes in hot rolling, the number of rolling passes at which the end temperature is 1050 ° C. or higher and the rolling reduction is 20% or higher (hereinafter, also referred to as rolling passes satisfying predetermined conditions) is 3 or more. There is a need to.
The upper limit of the number of rolling passes that satisfy the predetermined conditions is not particularly limited, but if it is excessively increased, a large amount of heat is required to maintain the rolling temperature, which leads to an increase in manufacturing cost. The number of rolling passes that satisfy the conditions is preferably 10 or less.
　また、熱間圧延における圧延パスごとの圧下率の上限は特に限定されるものではないが、圧延パスごとの圧下率が過度に大きくなると、圧延荷重が増加して圧延が困難になる。そのため、圧延パスごとの圧下率は６０％以下とすることが好ましい。
　ここで、圧延パスごとの圧下率とは、（［当該圧延パス開始時の被圧延材の板厚（ｍｍ）］－［当該圧延パス終了時の被圧延材の板厚（ｍｍ）］）／［当該圧延パス開始時の被圧延材の板厚（ｍｍ）］×１００として求めたものである。 Further, the upper limit of the rolling reduction ratio for each rolling pass in hot rolling is not particularly limited, but if the rolling reduction ratio for each rolling pass becomes excessively large, the rolling load increases and rolling becomes difficult. Therefore, the rolling reduction ratio for each rolling pass is preferably 60% or less.
Here, the rolling reduction ratio for each rolling pass is ([plate thickness of the material to be rolled (mm) at the start of the rolling pass]-[plate thickness of the material to be rolled (mm) at the end of the rolling pass]) / It was obtained as [plate thickness (mm) of the material to be rolled at the start of the rolling pass] × 100.
　なお、熱間圧延の圧延パス数（総数）は、８～２０パスとすることが好適である。また、熱間圧延は、粗圧延と仕上げ圧延により構成されることが一般的である。この場合、粗圧延の圧延パス数は３～１０パス、仕上げ圧延の圧延パス数は５～１０パスとすることが好適である。また、圧延終了温度は９００～１１００℃とすることが好適である。さらに、熱間圧延における総圧下率は、８５．０～９９．８％とすることが好適である。 The number of rolling passes (total number) for hot rolling is preferably 8 to 20 passes. Further, hot rolling is generally composed of rough rolling and finish rolling. In this case, it is preferable that the number of rolling passes for rough rolling is 3 to 10 and the number of rolling passes for finish rolling is 5 to 10. The rolling end temperature is preferably 900 to 1100 ° C. Further, the total rolling reduction in hot rolling is preferably 85.0 to 99.8%.
巻取り温度：６００℃以上
　熱間圧延の仕上げ圧延後に、熱延鋼板を巻取る。この際、オーステナイト相をフェライト相に変態させ、熱延鋼板の金属組織をフェライト相主体の組織とする。巻取り温度が６００℃未満の場合、オーステナイト相がマルテンサイト相に変態し、鋼板の硬質化を招く。また、鋼板の平坦度が悪化し、以降の工程の実施が困難となる場合がある。さらに、鋼板に焼割れが生じる場合がある。
　そのため、巻取り温度は６００℃以上とする。巻取り温度は、好ましくは６５０℃以上、より好ましくは７００℃以上、さらに好ましくは７５０℃以上である。巻取り温度の上限は特に限定されるものではないが、８５０℃以下とすることが好ましい。巻取り温度が８５０℃超になると、巻取り温度がオーステナイト相とフェライト相との二相温度域となる。そのため、オーステナイト相の安定性が高くなって、オーステナイト相からフェライト相への変態の遅滞が生じる。これにより、（巻取った鋼板の）大気放冷後でかつ熱延板焼鈍前に、オーステナイト相が硬質なマルテンサイト相へと変態する場合がある。その結果、熱延鋼板の著しい硬質化や形状不良が発生する場合があるため、好ましくない。 Winding temperature: 600 ° C or higher After the finish rolling of hot rolling, the hot-rolled steel sheet is wound. At this time, the austenite phase is transformed into a ferrite phase, and the metal structure of the hot-rolled steel sheet is made mainly of the ferrite phase. When the winding temperature is less than 600 ° C., the austenite phase is transformed into the martensite phase, which causes the steel sheet to become hard. In addition, the flatness of the steel sheet may deteriorate, making it difficult to carry out the subsequent steps. Further, the steel sheet may be cracked.
Therefore, the winding temperature is set to 600 ° C. or higher. The winding temperature is preferably 650 ° C. or higher, more preferably 700 ° C. or higher, and even more preferably 750 ° C. or higher. The upper limit of the winding temperature is not particularly limited, but is preferably 850 ° C. or lower. When the take-up temperature exceeds 850 ° C., the take-up temperature becomes a two-phase temperature range of an austenite phase and a ferrite phase. Therefore, the stability of the austenite phase becomes high, and the transformation from the austenite phase to the ferrite phase is delayed. As a result, the austenite phase may be transformed into a hard martensite phase after air cooling (of the wound steel sheet) and before annealing the hot-rolled sheet. As a result, the hot-rolled steel sheet may be significantly hardened or have a poor shape, which is not preferable.
・第３の工程：熱延板焼鈍工程
　ついで、第３の工程として、上記のようにして得た熱延鋼板に熱延板焼鈍を施し、熱延焼鈍鋼板とする。
　この熱延板焼鈍では、保持温度を７５０～９００℃、保持時間を１０分以上とする。 -Third step: Hot-rolled sheet annealing step Then, as a third step, the hot-rolled steel sheet obtained as described above is annealed by hot-rolled sheet to obtain a hot-rolled and annealed steel sheet.
In this hot-rolled sheet annealing, the holding temperature is 750 to 900 ° C., and the holding time is 10 minutes or more.
熱延板焼鈍の保持温度：７５０～９００℃
　熱延板焼鈍は、刃物等の所定形状への加工時の割れ（以下、加工割れともいう）を抑制することを目的として行われる。そして、この熱延板焼鈍では、再結晶により、熱間圧延により形成された圧延加工組織（ひずみを受けた結晶粒からなる金属組織）を、ひずみをほとんど含まないフェライト相の結晶粒に置き換える。 Holding temperature for hot-rolled sheet annealing: 750-900 ° C
The hot-rolled plate annealing is performed for the purpose of suppressing cracking (hereinafter, also referred to as machining cracking) during machining into a predetermined shape of a cutting tool or the like. Then, in this hot-rolled sheet annealing, the rolled structure (metal structure composed of strained crystal grains) formed by hot rolling is replaced with ferrite phase crystal grains containing almost no strain by recrystallization.
　ただし、熱延板焼鈍の保持温度（以下、熱延板焼鈍温度ともいう）が７５０℃未満になると、熱間圧延時に形成された圧延加工組織が残存するようになる。これにより、熱延焼鈍鋼板の延性が低下し、加工割れが生じやすくなる。また、熱延板焼鈍温度が９００℃を超えると、結晶粒が粗大化して、靭性が低下する。これにより、加工割れが生じやすくなる。
　そのため、熱延板焼鈍温度は７５０～９００℃の範囲とする。熱延板焼鈍温度は、好ましくは８００℃以上である。また、熱延板焼鈍温度は、好ましくは８７５℃以下、より好ましくは８５０℃以下である。
　なお、熱延板焼鈍温度は、保持中、一定であってもよく、また、上記の温度範囲内にあれば、保持中、常に一定としなくてもよい。以下で説明する冷延板焼鈍温度や焼入れ温度、焼戻し温度についても同様である。 However, when the holding temperature of hot-rolled sheet annealing (hereinafter, also referred to as hot-rolled sheet annealing temperature) becomes less than 750 ° C., the rolled structure formed during hot rolling remains. As a result, the ductility of the hot-rolled annealed steel sheet is reduced, and processing cracks are likely to occur. Further, when the hot-rolled plate annealing temperature exceeds 900 ° C., the crystal grains become coarse and the toughness decreases. As a result, processing cracks are likely to occur.
Therefore, the hot-rolled sheet annealing temperature is in the range of 750 to 900 ° C. The hot-rolled plate annealing temperature is preferably 800 ° C. or higher. The hot-rolled sheet annealing temperature is preferably 875 ° C. or lower, more preferably 850 ° C. or lower.
The hot-rolled sheet annealing temperature may be constant during holding, and may not always be constant during holding as long as it is within the above temperature range. The same applies to the cold-rolled sheet annealing temperature, quenching temperature, and tempering temperature described below.
熱延板焼鈍の保持時間：１０分以上
　熱延板焼鈍の保持時間が１０分未満の場合、鋼板内の材質を十分に均一化することができない。そのため、熱延板焼鈍の保持時間は１０分以上とする。熱延板焼鈍の保持時間は、好ましくは３時間以上、より好ましくは６時間以上である。なお、熱延板焼鈍の保持時間が９６時間超の場合、酸化スケールが厚くなって、その後の脱スケール処理が困難となる場合がある。そのため、熱延板焼鈍の保持時間は９６時間以下が好ましい。また、熱延板焼鈍の保持時間は、好ましくは２４時間以下、より好ましくは１２時間以下である。 Holding time of hot-rolled sheet annealing: If the holding time of hot-rolled sheet annealing is less than 10 minutes, the material in the steel sheet cannot be sufficiently uniformized. Therefore, the holding time for hot-rolled sheet annealing is set to 10 minutes or more. The holding time for hot-rolled sheet annealing is preferably 3 hours or longer, more preferably 6 hours or longer. If the holding time of hot-rolled sheet annealing exceeds 96 hours, the oxide scale may become thick and subsequent descaling treatment may become difficult. Therefore, the holding time for hot-rolled sheet annealing is preferably 96 hours or less. The holding time for hot-rolled sheet annealing is preferably 24 hours or less, more preferably 12 hours or less.
　また、熱延板焼鈍後、任意に、第４の工程として冷間圧延、さらに、第５の工程として、冷延板焼鈍を行ってもよい。 Further, after the hot-rolled plate is annealed, cold rolling may be optionally performed as the fourth step, and the cold-rolled plate may be annealed as the fifth step.
第４の工程：冷間圧延工程
　第４の工程では、熱延板焼鈍後に得られる熱延焼鈍鋼板に、冷間圧延を施して冷延鋼板とする。
　冷間圧延の手法に特に限定はなく、例えば、タンデムミルやクラスターミルを用いることができる。また、冷間圧延における圧下率についても特に限定はないが、冷延板焼鈍後の成形性や鋼板の形状矯正の観点から、冷間圧延における圧下率は５０％以上とすることが好ましい。また、過度の圧延荷重を避ける観点から、冷間圧延における圧下率は９５％以下とすることが好ましい。 Fourth Step: Cold Rolling Step In the fourth step, the hot-rolled annealed steel sheet obtained after annealing the hot-rolled sheet is cold-rolled to obtain a cold-rolled steel sheet.
The cold rolling method is not particularly limited, and for example, a tandem mill or a cluster mill can be used. The reduction ratio in cold rolling is also not particularly limited, but the reduction ratio in cold rolling is preferably 50% or more from the viewpoint of formability after annealing of a cold-rolled sheet and shape correction of a steel sheet. Further, from the viewpoint of avoiding an excessive rolling load, the rolling reduction in cold rolling is preferably 95% or less.
第５の工程：冷延板焼鈍工程
　第５の工程（冷延板焼鈍工程）では、冷間圧延後に得られる冷延鋼板に、保持温度：７００～８５０℃、保持時間：５秒以上の冷延板焼鈍を施して冷延焼鈍鋼板とする。
　冷延板焼鈍は、冷間圧延によって形成された圧延加工組織を再結晶により除去することを主目的として行う。 Fifth step: Cold-rolled sheet annealing step In the fifth step (cold-rolled sheet annealing step), the cold-rolled steel sheet obtained after cold rolling has a holding temperature of 700 to 850 ° C. and a holding time of 5 seconds or more. The rolled sheet is annealed to obtain a cold rolled annealed steel sheet.
The main purpose of cold rolled sheet annealing is to remove the rolled structure formed by cold rolling by recrystallization.
　ここで、冷延板焼鈍の保持温度（以下、冷延板焼鈍温度ともいう）が７００℃未満の場合、冷間圧延によって形成された圧延加工組織が残存して、冷延板焼鈍後に得られる冷延焼鈍鋼板の加工性が低下する。一方、冷延板焼鈍の保持における保持温度が８５０℃を超えると、オーステナイト相が生成し、保持後の冷却の際に、オーステナイト相がマルテンサイト相へと変態する。そのため、冷延板焼鈍後に得られる冷延焼鈍鋼板の硬質化および延性の低下を招き、結果的に、加工割れを招く。
　そのため、冷延板焼鈍を行う場合、冷延板焼鈍温度は７００～８５０℃の範囲とする。冷延板焼鈍温度は、好ましくは７２０℃以上である。また、冷延板焼鈍温度は、好ましくは８３０℃以下である。 Here, when the holding temperature of the cold-rolled plate annealing (hereinafter, also referred to as the cold-rolled plate annealing temperature) is less than 700 ° C., the rolled structure formed by the cold rolling remains and is obtained after the cold-rolled plate annealing. The workability of cold-rolled annealed steel sheets is reduced. On the other hand, when the holding temperature in the cold-rolled sheet annealing exceeds 850 ° C., an austenite phase is formed, and the austenite phase is transformed into a martensite phase during cooling after holding. Therefore, the cold-rolled annealed steel sheet obtained after the annealed cold-rolled sheet is hardened and the ductility is lowered, and as a result, processing cracks are caused.
Therefore, when the cold-rolled plate is annealed, the cold-rolled plate annealing temperature is in the range of 700 to 850 ° C. The cold rolled sheet annealing temperature is preferably 720 ° C. or higher. The cold rolled sheet annealing temperature is preferably 830 ° C. or lower.
　また、冷延板焼鈍の保持時間が５秒未満の場合、冷間圧延によって形成された圧延加工組織が残存し、冷延板焼鈍後に得られる冷延焼鈍鋼板の加工性が低下する。そのため、冷延板焼鈍を行う場合、冷延板焼鈍の保持時間は５秒以上とする。冷延板焼鈍の保持時間は、好ましくは１５秒以上である。
　一方、冷延板焼鈍の保持時間が２４時間を超えると、結晶粒が粗大化して、加工割れを招く場合がある。そのため、冷延板焼鈍の保持時間は、好ましくは２４時間以下とすることが好ましい。冷延板焼鈍の保持時間は、より好ましくは１５分以下である。 Further, when the holding time of the cold-rolled sheet annealing is less than 5 seconds, the rolled structure formed by the cold-rolling remains, and the workability of the cold-rolled annealed steel sheet obtained after the cold-rolled sheet annealing is lowered. Therefore, when the cold-rolled plate is annealed, the holding time of the cold-rolled plate annealing is set to 5 seconds or more. The holding time of the cold rolled sheet annealing is preferably 15 seconds or more.
On the other hand, if the holding time of the cold rolled sheet annealing exceeds 24 hours, the crystal grains may become coarse and cause processing cracks. Therefore, the holding time for cold-rolled sheet annealing is preferably 24 hours or less. The holding time for cold rolled sheet annealing is more preferably 15 minutes or less.
第６の工程：焼入れ処理工程
　上記のようにして得た熱延焼鈍鋼板、冷延鋼板または冷延焼鈍鋼板を、例えば、所定の形状に加工したのち、さらに、第６の工程として、保持温度：９５０～１２００℃、保持時間：５秒～３０分、保持後の平均冷却速度：１℃／秒以上の焼入れ処理を施して、焼入れ処理鋼板としてもよい。 Sixth step: Quenching treatment step The hot-rolled annealed steel sheet, cold-rolled annealed steel sheet or cold-rolled annealed steel sheet obtained as described above is processed into, for example, a predetermined shape, and then, as a sixth step, a holding temperature. : 950 to 1200 ° C., holding time: 5 seconds to 30 minutes, average cooling rate after holding: 1 ° C./sec or more may be subjected to quenching treatment to obtain a hardened steel plate.
　焼入れ処理の保持温度（以下、焼入れ温度ともいう）が９５０℃未満では、焼入れ処理における加熱および保持時にオーステナイト相が十分に生成せず、十分な焼きが入らない。焼入れ温度が１２００℃超では、焼入れ処理における加熱および保持時に金属組織中にデルタフェライト相が生成して焼きが十分に入らない場合がある。また、結晶粒が著しく粗大化して、冷却時の焼割れや加工割れが生じる場合がある。
　そのため、焼入れ温度は９５０～１２００℃の範囲とする。焼入れ温度は、好ましくは１０００℃以上である。また、焼入れ温度は、好ましくは１１５０℃以下である。 If the holding temperature of the quenching treatment (hereinafter, also referred to as the quenching temperature) is less than 950 ° C., the austenite phase is not sufficiently formed during heating and holding in the quenching treatment, and sufficient quenching does not occur. If the quenching temperature exceeds 1200 ° C., a delta ferrite phase may be formed in the metal structure during heating and holding in the quenching treatment, and quenching may not be sufficiently performed. In addition, the crystal grains may be remarkably coarsened, causing shrinkage or processing cracks during cooling.
Therefore, the quenching temperature is set in the range of 950 to 1200 ° C. The quenching temperature is preferably 1000 ° C. or higher. The quenching temperature is preferably 1150 ° C. or lower.
　また、焼入れ処理の保持時間が５秒未満では、加熱および保持時にオーステナイト相が十分に生成せず、十分な焼きが入らない。一方、焼入れ処理における保持時間が３０分を超えると、結晶粒の粗大化が生じて、加工割れが生じる場合がある。
　そのため、焼入れ処理の保持時間は５秒～３０分の範囲とする。焼入れ処理における保持時間は、好ましくは１５秒以上である。また、焼入れ処理における保持時間は、好ましくは３００秒以下、より好ましくは１２０秒以下である。 Further, if the holding time of the quenching treatment is less than 5 seconds, the austenite phase is not sufficiently formed during heating and holding, and sufficient quenching is not performed. On the other hand, if the holding time in the quenching treatment exceeds 30 minutes, coarsening of crystal grains may occur and processing cracks may occur.
Therefore, the holding time of the quenching process is in the range of 5 seconds to 30 minutes. The holding time in the quenching treatment is preferably 15 seconds or more. The holding time in the quenching treatment is preferably 300 seconds or less, more preferably 120 seconds or less.
　さらに、焼入れ処理における保持後に冷却する。この冷却の際の平均冷却速度、具体的には、焼入れ温度から４００℃までの平均冷却速度が１℃／秒未満の場合、加熱時に生成したオーステナイト相がマルテンサイト相ではなく、フェライト相へと変態するため、十分な焼きが入らない。
　そのため、焼入れ処理における保持後の平均冷却速度は１℃／秒以上とする。焼入れ処理における保持後の平均冷却速度は、好ましくは５℃／秒以上、より好ましくは１０℃／秒以上である。焼入れ処理における保持後の平均冷却速度の上限は特に限定されるものではないが、過度な急冷を行うと、鋼板形状の悪化や焼割れが生じる場合がある。そのため、焼入れ処理における保持後の平均冷却速度は、１０００℃／秒以下とすることが好ましい。 Further, it is cooled after being held in the quenching process. When the average cooling rate during this cooling, specifically, the average cooling rate from the quenching temperature to 400 ° C. is less than 1 ° C./sec, the austenite phase generated during heating becomes a ferrite phase instead of the martensite phase. Because it transforms, it does not get enough quenching.
Therefore, the average cooling rate after holding in the quenching treatment is set to 1 ° C./sec or more. The average cooling rate after holding in the quenching treatment is preferably 5 ° C./sec or higher, more preferably 10 ° C./sec or higher. The upper limit of the average cooling rate after holding in the quenching treatment is not particularly limited, but excessive quenching may cause deterioration of the steel sheet shape or quench cracking. Therefore, the average cooling rate after holding in the quenching treatment is preferably 1000 ° C./sec or less.
　なお、冷却の手法に特に限定はなく、空冷、ガス噴射冷却、ミスト水冷却、ロール冷却、水浸漬、金型冷却等の種々の方法を用いることができる。 The cooling method is not particularly limited, and various methods such as air cooling, gas injection cooling, mist water cooling, roll cooling, water immersion, and mold cooling can be used.
第７の工程：焼戻し処理工程
　ついで、硬度および耐久性を調整するため、上記の焼入れ処理鋼板に、第７の工程として、さらに、保持温度：１００～８００℃、保持時間：５分以上の焼戻し処理を施して、焼戻し処理鋼板としてもよい。 Seventh step: Tempering process Next, in order to adjust the hardness and durability, the above-mentioned tempered steel sheet is further tempered as a seventh step with a holding temperature of 100 to 800 ° C. and a holding time of 5 minutes or more. It may be treated to be a tempered steel sheet.
　焼戻し処理の保持温度（以下、焼戻し温度ともいう）が１００℃未満の場合、マルテンサイト相中の転位の回復が著しく遅くなる。そのため、焼戻し処理で目的とする軟質化効果を十分に得ることが困難となる。一方、焼戻し温度が８００℃を超えると、マルテンサイト相が再度オーステナイト相へと変態し、保持後の冷却時に、再度マルテンサイト相へと変態して硬質化する。そのため、焼戻し処理で目的とする軟質化効果を十分に得ることが困難となる。
　よって、焼戻し温度は、１００～８００℃の範囲とする。焼戻し温度は、好ましくは２００℃以上、より好ましくは４００℃以上である。また、焼戻し温度は、好ましくは７５０℃以下、より好ましくは７００℃以下である。 When the holding temperature of the tempering treatment (hereinafter, also referred to as tempering temperature) is less than 100 ° C., the recovery of dislocations in the martensite phase is significantly delayed. Therefore, it becomes difficult to sufficiently obtain the desired softening effect in the tempering treatment. On the other hand, when the tempering temperature exceeds 800 ° C., the martensite phase is transformed into the austenite phase again, and at the time of cooling after holding, it is transformed into the martensite phase again and hardened. Therefore, it becomes difficult to sufficiently obtain the desired softening effect in the tempering treatment.
Therefore, the tempering temperature is in the range of 100 to 800 ° C. The tempering temperature is preferably 200 ° C. or higher, more preferably 400 ° C. or higher. The tempering temperature is preferably 750 ° C. or lower, more preferably 700 ° C. or lower.
　また、焼戻し処理の保持時間（以下、焼戻し時間ともいう）が５分未満になると、マルテンサイト相中の転位の回復が不十分となる。そのため、焼戻し処理で目的とする軟質化効果を十分に得ることが困難となる。よって、焼戻し時間は５分以上とする。焼戻し時間は、好ましくは１０分以上、より好ましくは１５分以上である。
　なお、硬度は焼戻し時間が長くなるほど、低下する傾向にあるが、焼戻し時間が６０分を超えると、硬度はほぼ一定となる。そのため、焼戻し時間は６０分以下とすることが好ましい。焼戻し時間は、より好ましくは５０分以下、さらに好ましくは４０分以下である。 Further, when the holding time of the tempering treatment (hereinafter, also referred to as tempering time) is less than 5 minutes, the recovery of dislocations in the martensite phase becomes insufficient. Therefore, it becomes difficult to sufficiently obtain the desired softening effect in the tempering treatment. Therefore, the tempering time is set to 5 minutes or more. The tempering time is preferably 10 minutes or longer, more preferably 15 minutes or longer.
The hardness tends to decrease as the tempering time becomes longer, but when the tempering time exceeds 60 minutes, the hardness becomes almost constant. Therefore, the tempering time is preferably 60 minutes or less. The tempering time is more preferably 50 minutes or less, still more preferably 40 minutes or less.
　なお、上記以外の条件については、常法に従えばよい。
　また、任意に酸洗処理やショットブラストや表面研削等を、例えば、熱間圧延工程、熱延板焼鈍工程、冷間圧延工程、冷延板焼鈍工程、焼入れ工程および焼戻し工程の後等に行ってもよい。さらに、用途によっては、熱間圧延工程、熱延板焼鈍工程、冷延板焼鈍工程、焼入れ処理工程、および焼戻し処理工程の後等に、調質圧延を施してもよい。
　そして、上記のようにして得た鋼板を用いて、包丁やハサミ、医療用メス等の刃物、食卓用のナイフやフォーク、スプーン等のカトラリー、ならびに、ピンセット等の精密工具を得ることができる。 For conditions other than the above, the conventional method may be followed.
In addition, pickling treatment, shot blasting, surface grinding, etc. are optionally performed after, for example, a hot rolling step, a hot rolling plate annealing step, a cold rolling step, a cold rolling plate annealing step, a quenching step, and a tempering step. You may. Further, depending on the application, temper rolling may be performed after the hot rolling step, the hot rolling plate annealing step, the cold rolling plate annealing step, the quenching treatment step, the tempering treatment step, and the like.
Then, using the steel plate obtained as described above, a knife such as a kitchen knife, scissors, a medical scalpel, a cutlery such as a knife or fork for a table, a spoon, and a precision tool such as tweezers can be obtained.
　表１に示した成分組成を有する鋼（残部はＦｅおよび不可避的不純物）を、容量：１５０ｔｏｎの転炉による精錬と強攪拌・真空酸素脱炭処理（ＳＳ－ＶＯＤ）による精錬で溶製し、連続鋳造により、幅：１０００ｍｍ、厚さ：２００ｍｍの鋼スラブとした。
　該鋼スラブを、表２に記載の条件で保持後、表２および３に記載の条件で熱間圧延および熱延板焼鈍を施し、熱延焼鈍鋼板とした。なお、熱間圧延の（合計の）パス数は、いずれも１４パスとした。また、熱間圧延における１～５パス目の終了温度は、６パス目の終了温度よりも高い温度であるため、表２では記載を省略した。加えて、熱間圧延における９パス目以降の終了温度も、表２では記載を省略した。
　ついで、一部の熱延焼鈍鋼板について、さらに、表３に記載の条件で、冷間圧延および／または冷延板焼鈍を施し、冷延鋼板および／または冷延焼鈍鋼板を得た。
　かくして得られた熱延焼鈍鋼板、冷延鋼板および冷延焼鈍鋼板について、上述の方法により金属組織の観察を行い、金属組織を同定した。結果を表４に示す。ただし、Ｎｏ．３５では、熱延鋼板の巻取りの際に、割れが生じたため、金属組織の同定、および、これ以降の評価は行わなかった。 Steel having the composition shown in Table 1 (the balance is Fe and unavoidable impurities) is smelted by refining with a converter having a capacity of 150 tons and refining with strong stirring and vacuum oxygen decarburization (SS-VOD). A steel slab having a width of 1000 mm and a thickness of 200 mm was obtained by continuous casting.
The steel slab was held under the conditions shown in Table 2 and then hot-rolled and hot-rolled and annealed under the conditions shown in Tables 2 and 3 to obtain a hot-rolled annealed steel sheet. The number of (total) passes for hot rolling was 14 in each case. Further, since the end temperature of the 1st to 5th passes in hot rolling is higher than the end temperature of the 6th pass, the description is omitted in Table 2. In addition, the end temperature after the 9th pass in hot rolling is also omitted in Table 2.
Then, some hot-rolled annealed steel sheets were further subjected to cold rolling and / or cold-rolled sheet annealing under the conditions shown in Table 3 to obtain cold-rolled steel sheets and / or cold-rolled annealed steel sheets.
The metal structures of the hot-rolled and cold-rolled steel sheets, the cold-rolled steel sheets, and the cold-rolled and hardened steel sheets thus obtained were observed by the above-mentioned methods, and the metal structures were identified. The results are shown in Table 4. However, No. In No. 35, since cracks occurred during winding of the hot-rolled steel sheet, the metallographic structure was not identified and the subsequent evaluation was not performed.
　また、上記のようにして得た熱延焼鈍鋼板、冷延鋼板および冷延焼鈍鋼板を、圧延方向：３００ｍｍ×幅方向：５０ｍｍに打ち抜き加工した。そして、加工後の鋼板に、焼入れ温度：１０５０℃、保持時間：１５分間、保持後の焼入れ温度から４００℃までの平均冷却速度：５℃／ｓの条件で、空冷により、焼入れ処理を施した。
　なお、Ｎｏ．１Ａおよび３Ａ－１、３Ａ－２は、焼入れ処理後のＮｏ．１および３の鋼板に、さらに、表３に記載の条件で焼戻し処理を行ったもの（焼戻し処理鋼板）である。 Further, the hot-rolled annealed steel sheet, the cold-rolled annealed steel sheet and the cold-rolled annealed steel sheet obtained as described above were punched in a rolling direction: 300 mm × width direction: 50 mm. Then, the processed steel sheet was quenched by air cooling under the conditions of quenching temperature: 1050 ° C., holding time: 15 minutes, and average cooling rate from the quenching temperature after holding to 400 ° C.: 5 ° C./s. ..
In addition, No. 1A and 3A-1 and 3A-2 are No. 1 after quenching treatment. The steel sheets 1 and 3 are further tempered under the conditions shown in Table 3 (tempering steel sheet).
　かくして得られた焼入れ処理鋼板および焼戻し処理鋼板について、上述の方法により金属組織の観察を行い、金属組織を同定した。結果を表４に併記する。 The metallographic structure of the thus obtained hardened steel sheet and tempered steel sheet was observed by the above-mentioned method, and the metallographic structure was identified. The results are also shown in Table 4.
　また、以下の要領で、硬度および表面品質の評価を行った。
　なお、硬度の評価は、焼入れ処理鋼板を用いて行った。ただし、焼き戻し処理を行ったＮｏ．１Ａおよび３Ａ－１、３Ａ－２では、焼戻し処理後の鋼板についても、硬度の評価を行った。
　また、表面品質の評価は、最終的に得られた鋼板、つまり、Ｎｏ．１～３７では焼入れ処理鋼板を、Ｎｏ．１Ａおよび３Ａ－１、３Ａ－２では焼戻し処理鋼板を用いて行った。 In addition, the hardness and surface quality were evaluated as follows.
The hardness was evaluated using a hardened steel sheet. However, the tempered No. In 1A, 3A-1 and 3A-2, the hardness of the tempered steel sheet was also evaluated.
In addition, the surface quality was evaluated by the finally obtained steel sheet, that is, No. In Nos. 1 to 37, the hardened steel sheet was used as No. In 1A, 3A-1 and 3A-2, tempered steel sheets were used.
＜硬度の評価＞
　上記のようにして得た鋼板の圧延面において、ＪＩＳ　Ｚ　２２４５（２０１６年）に準拠したロックウェル硬さ試験を任意の５点で行い、当該５点でのロックウェル硬度の平均値を求めた。なお、鋼板の圧延面は、試験前に、♯４００の耐水エメリー研磨紙で表面研磨した。そして、以下の基準により、硬度の評価を行った。評価結果を表４に併記する。
・焼戻し処理を行わない場合
　〇（合格）：ロックウェル硬度の平均値がＨＲＣ５５以上
　×（不合格）：ロックウェル硬度の平均値がＨＲＣ５５未満
・焼戻し処理を行う場合
　〇（合格）：焼戻し処理前のロックウェル硬度の平均値がＨＲＣ５５以上でかつ、焼戻し処理後のロックウェル硬度の平均値がＨＲＣ４０以上
　×（不合格）：焼戻し処理前のロックウェル硬度の平均値がＨＲＣ５５未満、または、焼戻し処理後のロックウェル硬度の平均値がＨＲＣ４０未満 <Evaluation of hardness>
On the rolled surface of the steel sheet obtained as described above, a Rockwell hardness test conforming to JIS Z 2245 (2016) was performed at any 5 points, and the average value of the Rockwell hardness at the 5 points was obtained. .. The rolled surface of the steel sheet was surface-polished with # 400 water-resistant emery polishing paper before the test. Then, the hardness was evaluated according to the following criteria. The evaluation results are also shown in Table 4.
・ When tempering is not performed 〇 (Pass): Average value of Rockwell hardness is HRC55 or more × (Failure): Average value of Rockwell hardness is less than HRC55 ・ When tempering is performed 〇 (Pass): Before tempering The average value of Rockwell hardness before tempering is HRC55 or more and the average value of Rockwell hardness after tempering is HRC40 or more × (Failure): The average value of Rockwell hardness before tempering is less than HRC55 or tempering The average value of the later Rockwell hardness is less than HRC40
＜表面品質の評価＞
　上記のようにして得た鋼板から、圧延方向：１００ｍｍ×幅方向：５０ｍｍの試験片を１０枚採取した。ついで、図３に示すように、各試験片に対して、圧延方向および幅方向に平行な端面の１つを、幅方向に対して３．５°の角度で切削加工を施した。ついで、切削面を、＃４００→＃６００→＃８００→＃１２００→＃２０００の耐水エメリーペーパーの順で、湿式クロス研磨（次番手の研磨を先番手の研磨方向と直角方向に行う研磨）することにより、刃付け研磨面を設けた。
　そして、当該刃付け研磨面を目視により観察し、以下の基準で、表面品質の評価を行った。評価結果を表４に併記する。
　〇（合格）：１０枚全ての試験片で、刃付け研磨面に長さ：２．０ｍｍ以上の筋模様が観察されない。
　×（不合格）：１０枚の試験片の少なくとも１つで、刃付け研磨面に長さ：２．０ｍｍ以上の筋模様が観察される。 <Evaluation of surface quality>
From the steel sheet obtained as described above, 10 test pieces having a rolling direction of 100 mm and a width direction of 50 mm were collected. Then, as shown in FIG. 3, one of the end faces parallel to the rolling direction and the width direction was cut on each test piece at an angle of 3.5 ° with respect to the width direction. Next, the cutting surface is subjected to wet cloth polishing (polishing in the direction perpendicular to the polishing direction of the first count) in the order of # 400 → # 600 → # 800 → # 1200 → # 2000 water resistant emery paper. As a result, a polished surface with a blade was provided.
Then, the polished surface with a blade was visually observed, and the surface quality was evaluated according to the following criteria. The evaluation results are also shown in Table 4.
〇 (Pass): No streaks with a length of 2.0 mm or more are observed on the polished surface of the blade on all 10 test pieces.
X (Failure): With at least one of the 10 test pieces, a streak pattern having a length of 2.0 mm or more is observed on the polished surface of the blade.
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-I000003
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-I000003
Figure JPOXMLDOC01-appb-T000004
Figure JPOXMLDOC01-appb-T000004
Figure JPOXMLDOC01-appb-T000005
Figure JPOXMLDOC01-appb-T000005
　表４に示すように、発明例ではいずれも、高い硬度を有し、かつ、良好な表面品質が得られていた。
　一方、比較例であるＮｏ．３０、３３および３４では、熱間圧延における所定条件を満足する圧延パス数が３パス未満であったため、粒径：２．０μｍ以上のＣｒ系炭化物の合計の体積率が１０％超となった。そのため、良好な表面品質が得られなかった。
　Ｎｏ．３１では、スラブ加熱温度が適正範囲を超えたため、粒径：２．０μｍ以上のＣｒ系炭化物の合計の体積率が１０％超となった。そのため、良好な表面品質が得られなかった。
　Ｎｏ．３２では、スラブ加熱温度が適正範囲に満たなかったため、粒径：２．０μｍ以上のＣｒ系炭化物の合計の体積率が１０％超となった。そのため、良好な表面品質が得られなかった。
　Ｎｏ．３５では、熱間圧延の巻取り温度が適正範囲に満たなかったため、熱延鋼板に割れが発生した。
　Ｎｏ．３６および３７では、Ｃ含有量が適正範囲に満たなかったため、焼入れ処理後の硬度が適正範囲に満たなかった。なお、Ｎｏ．３６では、Ｃ含有量が適正範囲に満たなかったため、熱間圧延における所定条件を満足する圧延パス数が３パス未満であるものの、粒径：２．０μｍ以上のＣｒ系炭化物の合計の体積率は１０％以下となった。 As shown in Table 4, all of the examples of the invention had high hardness and good surface quality.
On the other hand, No. In Nos. 30, 33 and 34, the number of rolling passes satisfying the predetermined conditions in hot rolling was less than 3, so that the total volume fraction of Cr-based carbides having a particle size of 2.0 μm or more was more than 10%. .. Therefore, good surface quality could not be obtained.
No. In No. 31, since the slab heating temperature exceeded the appropriate range, the total volume fraction of Cr-based carbides having a particle size of 2.0 μm or more was more than 10%. Therefore, good surface quality could not be obtained.
No. In No. 32, since the slab heating temperature was not within the appropriate range, the total volume fraction of Cr-based carbides having a particle size of 2.0 μm or more was more than 10%. Therefore, good surface quality could not be obtained.
No. In No. 35, the take-up temperature of hot rolling was not within the appropriate range, so that the hot-rolled steel sheet was cracked.
No. In 36 and 37, the C content was not in the appropriate range, so the hardness after the quenching treatment was not in the appropriate range. In addition, No. In No. 36, since the C content was not within the appropriate range, the number of rolling passes satisfying the predetermined conditions in hot rolling was less than 3, but the total volume fraction of Cr-based carbides having a particle size of 2.0 μm or more. Was less than 10%.
　なお、参考のため、良好な表面品質が得られていた発明例のＮｏ．１の圧延方向に平行な断面における光学顕微鏡組織写真を図１に示す。また、良好な表面品質が得られなかった比較例のＮｏ．３０の圧延方向に平行な断面の光学顕微鏡組織写真を図２に示す。 For reference, No. 1 of the invention example in which good surface quality was obtained. An optical microscope microstructure photograph in a cross section parallel to the rolling direction of No. 1 is shown in FIG. In addition, No. 1 of Comparative Example in which good surface quality could not be obtained. An optical microscope microstructure photograph of a cross section parallel to the rolling direction of 30 is shown in FIG.
　本発明のステンレス鋼板は、高い硬度と良好な表面品質とを有するので、包丁やハサミ、医療用メス等の刃物、食卓用のナイフやフォーク、スプーン等のカトラリー、ならびに、ピンセット等の精密工具の材料に、好適に用いることができる。 Since the stainless steel plate of the present invention has high hardness and good surface quality, it can be used for knives such as kitchen knives, scissors, medical scalpels, cutlery such as table knives, forks and spoons, and precision tools such as tweezers. It can be suitably used as a material.

Claims (10)

  1. 　質量％で、
    　Ｃ：０．４５～０．６０％、
    　Ｓｉ：０．０５～１．００％、
    　Ｍｎ：０．０５～１．００％、
    　Ｐ：０．０５％以下、
    　Ｓ：０．０２０％以下、
    　Ｃｒ：１３．０％以上１６．０％未満、
    　Ｎｉ：０．１０～１．００％および
    　Ｎ：０．０１０～０．２００％
    を含有し、残部がＦｅおよび不可避的不純物からなる成分組成を有し、
    　粒径：２．０μｍ以上のＣｒ系炭化物の合計の体積率が１０％以下である、ステンレス鋼板。     By mass%
    C: 0.45 to 0.60%,
    Si: 0.05 to 1.00%,
    Mn: 0.05 to 1.00%,
    P: 0.05% or less,
    S: 0.020% or less,
    Cr: 13.0% or more and less than 16.0%,
    Ni: 0.10 to 1.00% and N: 0.010 to 0.200%
    Has a component composition in which the balance is composed of Fe and unavoidable impurities.
    Particle size: A stainless steel sheet having a total volume fraction of Cr-based carbides of 2.0 μm or more of 10% or less.
  2. 　前記成分組成が、さらに、質量％で、
    　Ｍｏ：０．０５～１．００％、
    　Ｃｕ：０．０５～１．００％および
    　Ｃｏ：０．０５～０．５０％
    のうちから選ばれる１種または２種以上を含有する、請求項１に記載のステンレス鋼板。     The component composition is further increased by mass%.
    Mo: 0.05-1.00%,
    Cu: 0.05 to 1.00% and Co: 0.05 to 0.50%
    The stainless steel sheet according to claim 1, which contains one or more selected from the above.
  3. 　前記成分組成が、さらに、質量％で、
    　Ａｌ：０．００１～０．１００％、
    　Ｔｉ：０．０１～０．１０％、
    　Ｎｂ：０．０１～０．１０％、
    　Ｖ：０．０５～０．５０％、
    　Ｚｒ：０．０１～０．１０％、
    　Ｍｇ：０．０００２～０．００５０％、
    　Ｂ：０．０００２～０．００５０％、
    　Ｃａ：０．０００３～０．００３０％および
    　ＲＥＭ：０．０１～０．１０％
    のうちから選ばれる１種または２種以上を含有する、請求項１または２に記載のステンレス鋼板。     The component composition is further increased by mass%.
    Al: 0.001 to 0.100%,
    Ti: 0.01 to 0.10%,
    Nb: 0.01 to 0.10%,
    V: 0.05 to 0.50%,
    Zr: 0.01-0.10%,
    Mg: 0.0002 to 0.0050%,
    B: 0.0002 to 0.0050%,
    Ca: 0.0003 to 0.0030% and REM: 0.01 to 0.10%
    The stainless steel sheet according to claim 1 or 2, which contains one or more selected from the above.
  4. 　請求項１～３のいずれかに記載のステンレス鋼板を製造するための方法であって、
    　請求項１～３のいずれかに記載の成分組成を有する鋼スラブを、１２００～１３５０℃で３０分以上保持する、第１の工程と、
    　前記鋼スラブに、熱間圧延を施して熱延鋼板とし、該熱延鋼板を巻取る、第２の工程と、
    　前記熱延鋼板に熱延板焼鈍を施し、熱延焼鈍鋼板とする、第３の工程と、をそなえ、
    　前記第２の工程の熱間圧延における圧延パスのうち、終了温度：１０５０℃以上で、かつ、圧下率：２０％以上の圧延パス数が３パス以上であり、また、前記熱延鋼板の巻取り温度が６００℃以上であり、
    　前記第３の工程の熱延板焼鈍における保持温度が７５０～９００℃、保持時間が１０分以上である、
    　ステンレス鋼板の製造方法。     A method for manufacturing the stainless steel sheet according to any one of claims 1 to 3.
    The first step of holding the steel slab having the component composition according to any one of claims 1 to 3 at 1200 to 1350 ° C. for 30 minutes or more.
    The second step of hot-rolling the steel slab to obtain a hot-rolled steel sheet and winding the hot-rolled steel sheet.
    A third step of subjecting the hot-rolled steel sheet to hot-rolled annealed steel sheet to obtain a hot-rolled and annealed steel sheet is provided.
    Among the rolling passes in the hot rolling of the second step, the number of rolling passes having an end temperature of 1050 ° C. or higher and a reduction rate of 20% or higher is 3 or more, and the hot-rolled steel sheet is wound. The taking temperature is 600 ° C or higher,
    The holding temperature in the hot-rolled sheet annealing of the third step is 750 to 900 ° C., and the holding time is 10 minutes or more.
    Manufacturing method of stainless steel plate.
  5. 　前記熱延焼鈍鋼板に、冷間圧延を施して冷延鋼板とする、第４の工程をそなえる、請求項４に記載のステンレス鋼板の製造方法。 The method for manufacturing a stainless steel sheet according to claim 4, further comprising a fourth step of cold-rolling the hot-rolled annealed steel sheet to obtain a cold-rolled steel sheet.
  6. 　前記冷延鋼板に、冷延板焼鈍を施して冷延焼鈍鋼板とする、第５の工程をそなえ、
    　前記冷延板焼鈍における保持温度が７００～８５０℃、保持時間が５秒以上である、請求項５に記載のステンレス鋼板の製造方法。     The cold-rolled steel sheet is subjected to cold-rolled sheet annealing to obtain a cold-rolled annealed steel sheet, which is provided with a fifth step.
    The method for producing a stainless steel sheet according to claim 5, wherein the holding temperature in the cold rolled sheet annealing is 700 to 850 ° C. and the holding time is 5 seconds or more.
  7. 　前記熱延焼鈍鋼板、前記冷延鋼板、または、前記冷延焼鈍鋼板に、焼入れ処理を施す、第６の工程をそなえ、
    　前記焼入れ処理における保持温度が９５０～１２００℃、保持時間が５秒～３０分、保持後の平均冷却速度が１℃／秒以上である、請求項４～６のいずれかに記載のステンレス鋼板の製造方法。     A sixth step of quenching the hot-rolled annealed steel sheet, the cold-rolled steel sheet, or the cold-rolled annealed steel sheet is provided.
    The stainless steel sheet according to any one of claims 4 to 6, wherein the holding temperature in the quenching treatment is 950 to 1200 ° C., the holding time is 5 seconds to 30 minutes, and the average cooling rate after holding is 1 ° C./sec or more. Production method.
  8. 　前記焼入れ処理を施した鋼板に、焼戻し処理を施す、第７の工程をそなえ、
    　前記焼戻し処理における保持温度が１００～８００℃、保持時間が５分以上である、請求項７に記載のステンレス鋼板の製造方法。     A seventh step of tempering the hardened steel sheet is provided.
    The method for producing a stainless steel sheet according to claim 7, wherein the holding temperature in the tempering treatment is 100 to 800 ° C. and the holding time is 5 minutes or more.
  9. 　請求項１～３のいずれかに記載のステンレス鋼板を用いてなる、刃物。 A knife made of the stainless steel plate according to any one of claims 1 to 3.
  10. 　請求項１～３のいずれかに記載のステンレス鋼板を用いてなる、カトラリー。 Cutlery using the stainless steel plate according to any one of claims 1 to 3.
PCT/JP2021/014829 2020-04-30 2021-04-07 Stainless steel sheet, method for producing same, edged tools and cutlery WO2021220754A1 (en)

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