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WO2017002147A1 - Ferritic stainless steel sheet and method for manufacturing same - Google Patents

Ferritic stainless steel sheet and method for manufacturing same Download PDF

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Publication number
WO2017002147A1
WO2017002147A1 PCT/JP2015/003339 JP2015003339W WO2017002147A1 WO 2017002147 A1 WO2017002147 A1 WO 2017002147A1 JP 2015003339 W JP2015003339 W JP 2015003339W WO 2017002147 A1 WO2017002147 A1 WO 2017002147A1
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WIPO (PCT)
Prior art keywords
phase
stainless steel
ferritic stainless
annealing
hot
Prior art date
Application number
PCT/JP2015/003339
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French (fr)
Japanese (ja)
Inventor
正崇 吉野
映斗 水谷
光幸 藤澤
力 上
Original Assignee
Ｊｆｅスチール株式会社
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by Ｊｆｅスチール株式会社 filed Critical Ｊｆｅスチール株式会社
Priority to KR1020177036666A priority Critical patent/KR102027769B1/en
Priority to JP2015552676A priority patent/JP5884211B1/en
Priority to CN201580081418.2A priority patent/CN107709592B/en
Priority to US15/737,932 priority patent/US20180171430A1/en
Priority to PCT/JP2015/003339 priority patent/WO2017002147A1/en
Priority to TW104122004A priority patent/TW201702406A/en
Publication of WO2017002147A1 publication Critical patent/WO2017002147A1/en

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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/004—Heat treatment of ferrous alloys containing Cr and Ni
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/004—Very low carbon steels, i.e. having a carbon content of less than 0,01%
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel with vanadium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • 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
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/52—Ferrous alloys, e.g. steel alloys containing chromium with nickel with cobalt
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel with boron
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite

Definitions

  • the present invention relates to a ferritic stainless steel sheet having sufficient corrosion resistance and excellent formability and ridging resistance, and a method for producing the same.
  • Ferritic stainless steel sheet is more economical than austenitic stainless steel containing a large amount of expensive Ni.
  • SUS430 stainless steel plate (16 to 18% by mass Cr) is particularly economical and is used in various applications such as building materials, transportation equipment, home appliances, kitchen appliances, and automobile parts. The range of application has been further expanded in recent years. In order to apply to these applications, not only corrosion resistance but also sufficient formability that can be processed into a predetermined shape is required.
  • SUS430 stainless steel sheet is often applied to applications that require a good appearance, and is required to have excellent ridging resistance.
  • Ridging is a surface irregularity generated due to distortion in molding.
  • a crystal grain group (colony) having a similar crystal orientation may be generated during casting and / or hot rolling.
  • colony crystal grain group
  • a large difference occurs in the strain amount between the colony part and other parts at the time of forming, and thus surface irregularities (ridging) occur after forming.
  • a polishing step is required to remove surface irregularities, and the manufacturing cost of the molded product increases.
  • Patent Document 1 by mass, C: 0.02 to 0.06%, Si: 1.0% or less, Mn: 1.0% or less, P: 0.05% or less, S: 0.01%
  • Al 0.005% or less
  • Ti 0.005% or less
  • Cr 11 to 30%
  • Ni 0.7% or less
  • An excellent ferritic stainless steel is disclosed. However, when the present inventors manufactured ferritic stainless steel by the method described in Patent Document 1, excellent elongation at break was obtained in the rolling direction of the steel sheet.
  • box annealing for example, annealing for 8 hours at 860 degreeC
  • box annealing has a problem of low productivity because it takes about one week when heating and cooling processes are included.
  • the manufacturing cost increases because a technique for reducing solid solution N by adding V, which is an expensive transition metal element, is used.
  • box annealing is performed in the single-phase temperature range of ferrite for hot-rolled sheet annealing, ferrite colonies remain with almost no destruction, and there is also a problem that ridging resistance is significantly lowered.
  • Patent Document 2 by mass, C: 0.01 to 0.10%, Si: 0.05 to 0.50%, Mn: 0.05 to 1.00%, Ni: 0.01 to 0.00. 50%, Cr: 10 to 20%, Mo: 0.005 to 0.50%, Cu: 0.01 to 0.50%, V: 0.001 to 0.50%, Ti: 0.001 to 0 .50%, Al: 0.01-0.20%, Nb: 0.001-0.50%, N: 0.005-0.050% and B: 0.00010-0.00500%
  • hot-rolled sheet annealing is performed in the ferrite single-phase temperature range using a box furnace or a continuous furnace of AP line (Annealing and Pickling line), followed by cold rolling and finish annealing.
  • Patent Document 2 As in Patent Document 1, when an attempt was made to produce a press member mainly composed of stretch forming, it could not be formed into a predetermined shape, and as expected from elongation at break. No stretch formability was obtained. Furthermore, in general, ferritic stainless steel as in Patent Document 2 generates a group of crystal grains (colony) having a similar crystal orientation during casting or hot rolling, and hot-rolled sheet annealing is performed in a ferrite single-phase temperature range. The ferrite phase colonies cannot be sufficiently destroyed. Therefore, there is a problem that the colony expands and remains in the rolling direction by cold rolling after hot-rolled sheet annealing, and significant ridging occurs after forming.
  • colony crystal grains
  • Patent Document 3 a ferritic stainless steel containing 0.15% or less of C and 13 to 25% of Cr, and a range of 930 to 990 ° C. in which austenite and a ferrite phase coexist on the hot-rolled sheet of this steel.
  • the structure becomes a two-phase structure of martensite phase and ferrite phase, and then cold rolling and cold-rolled sheet annealing are performed, which is excellent in ridging resistance and workability
  • a method for producing a ferritic stainless steel sheet is disclosed.
  • Patent Document 3 only elongation is mentioned as workability.
  • Japanese Patent No. 3584881 (Republished WO 00/60134) Japanese Patent No. 3581801 (Japanese Patent Laid-Open No. 2001-3134) Japanese Patent Publication No.47-1878
  • An object of the present invention is to solve such problems and to provide a ferritic stainless steel sheet having sufficient corrosion resistance, excellent formability and ridging resistance, and a method for producing the same.
  • sufficient corrosion resistance refers to a salt spray cycle test (salt spray (35 ° C., 5 ° C., JIS H 8502) applied to a steel plate whose surface is polished with # 600 emery paper and the end face is sealed. (Mass% NaCl, spraying 2 hr) ⁇ drying (60 ° C., relative humidity 40%, 4 hr) ⁇ wetting (50 ° C., relative humidity ⁇ 95%, 2
  • excellent moldability means having excellent stretch formability, elongation at break, and average r value.
  • Excellent stretch formability means that the minimum value of the maximum logarithmic strain of the forming limit determined based on the forming limit diagram (FLD) of steel is 0.15 or more.
  • the excellent elongation at break means that the elongation at break (El) in a tensile test according to JIS Z 2241 is 28% or more in a test piece perpendicular to the rolling direction.
  • An excellent average r value is an average rankford value (hereinafter referred to as an average r value) calculated by the following equation (1) when a strain of 15% is applied in a tensile test based on JIS Z 2241: It means that it is more than .75.
  • r L is an r value when a tensile test is performed in a direction parallel to the rolling direction
  • r D is an r value when a tensile test is performed in a direction of 45 ° with respect to the rolling direction
  • r C is a direction perpendicular to the rolling direction. The r value when a tensile test is performed.
  • excellent ridging resistance means that the ridging height measured by the following method is 2.5 ⁇ m or less.
  • a JIS No. 5 tensile test piece is taken in parallel with the rolling direction.
  • 20% tensile strain is applied.
  • the arithmetic average waviness (Wa) defined by JIS B 0601 (2001) is measured with a surface roughness meter in the direction perpendicular to the rolling direction on the polishing surface at the center of the parallel part of the test piece.
  • the measurement conditions are a measurement length of 16 mm, a high cut filter wavelength of 0.8 mm, and a low cut filter wavelength of 8 mm. This arithmetic mean swell is defined as the ridging height.
  • An appropriate component ferritic stainless steel sheet is annealed at a suitable temperature range of a ferrite phase and an austenite phase after hot rolling and before cold rolling (hereinafter referred to as hot rolled sheet annealing). Further, by annealing the cold-rolled steel sheet at a temperature that becomes a ferrite single-phase region (hereinafter referred to as cold-rolled sheet annealing), it is a ferrite single-phase structure, but an intragranular carbonitride. A mixed grain structure of ferrite grains having a large amount of ferrite and ferrite grains having a small amount of carbonitride in the grains. As a result, it has been found that a ferritic stainless steel sheet having sufficient corrosion resistance and excellent formability and ridging resistance can be obtained.
  • all% which shows the component of steel is the mass%.
  • a ferritic stainless steel sheet having sufficient corrosion resistance and excellent formability and ridging resistance can be obtained.
  • the ferritic stainless steel sheet of the present invention is intended to be used in various applications such as building material parts, home appliance parts, kitchen appliances, or automobile parts by press working. In order to apply to these uses, sufficient moldability is required.
  • the present inventors conducted an overhang forming test assuming a ventilation hood using various ferritic stainless steel sheets (including those corresponding to Patent Documents 1 to 3) having different components and production methods.
  • various ferritic stainless steel sheets including those corresponding to Patent Documents 1 to 3 having different components and production methods.
  • the superiority or inferiority of the stretch formability is not necessarily determined by the magnitude of the elongation at break. Therefore, when the steel sheet used in the bulge forming test was prepared and the bulge formability was evaluated in detail by creating an FLD (formation limit diagram), good formability was obtained by the bulge hood assumed above. It has become clear that a stretchable formability of 0.15 or more, preferably 0.18 or more is required as the minimum value of the maximum logarithmic strain at the forming limit based on FLD.
  • the present inventors investigated the cause of the case where the superiority or inferiority of the stretch formability of the ferritic stainless steel plate obtained by the conventional technique does not correspond to the magnitude of the elongation at break.
  • the structure after cold rolling annealing is a ferrite single-phase structure in which carbonitrides are dispersed in a large amount and uniformly, and this is the cause. It was.
  • voids are generated in the structure as the amount of strain increases, and when these voids are connected, they become cracks and eventually break.
  • the ferritic stainless steel sheet obtained by the conventional technique is a ferrite single-phase structure in which carbonitride is dispersed in a large amount and uniformly. A very large amount of voids are generated from the entire surface. That is, in the prior art, cracks due to the connection of voids are likely to occur. As a result, even in the case of uniaxial deformation such as a tensile test, rupture occurs because void connection occurs in all directions in stretch forming in which multiaxial stress and strain are applied, even though it shows high elongation at break. It was easy to find out that sufficient stretchability could not be obtained.
  • the present inventors performed hot rolling sheet annealing on a steel sheet having an appropriate component in a two-phase region of a ferrite phase and an austenite phase, and then cold-rolled the steel sheet in a conventional manner, and further performed cold rolling sheet annealing on a single ferrite sheet.
  • an austenite phase with an area ratio of 3 to 20% is generated by hot-rolled sheet annealing. Almost all of the austenite phase is transformed into a martensite phase in the cooling process after hot-rolled sheet annealing.
  • the martensite phase is decomposed into a ferrite phase and a carbonitride during cold-rolled sheet annealing.
  • the structure after the cold-rolled sheet annealing becomes ferrite grains formed by the decomposition of ferrite grains that were ferrite phases and martensite phases from the beginning. That is, there are a large amount of carbonitrides in the grain boundaries and in the grains of the ferrite phase formed by the decomposition of the martensite phase, and in the entire metal structure, ferrite with an extremely large amount of carbonitrides in and on the grain boundaries. It becomes a mixed grain structure composed of ferrite grains with few grains and carbonitrides.
  • the ferrite grains having a high carbonitride are relatively hard, and a hardness difference of a grain unit occurs in the metal structure. It was found that when such a steel sheet is stretched, voids are mainly generated from carbonitrides on the interface between ferrite grains with a large amount of carbonitride and few ferrite grains, and the amount of voids generated in other parts is small. .
  • the steel according to the present invention there is little void formation in a portion where ferrite grains having a large amount of carbonitride are continuously located, a portion where ferrite grains having a small amount of carbonitride are continuous, and a ferrite grain. Therefore, the distance between the voids is longer than that of the ferritic stainless steel plate obtained by the prior art, cracks due to void connection during overhang forming are less likely to occur, and the minimum value of the maximum logarithmic strain at the forming limit based on FLD is small. High stretch formability of 0.15 or more is manifested.
  • the ferrite grains with many carbonitrides generated by subsequent cold-rolled sheet annealing increase, and the interface area between the ferrite grains with many carbonitrides and the few ferrite grains that become the starting point of voids during processing increases.
  • the overhang moldability cannot be expressed. Therefore, the upper limit of each of the C content and the N content needs to be 0.025%.
  • a predetermined amount of austenite phase can be stably secured by annealing in the two-phase region of the ferrite phase and the austenite phase, particularly in the range of 900 to 1100 ° C. Good surface quality can be obtained without excessively coarsening the particle size.
  • the steel having the C content and the N content is subjected to hot-rolled sheet annealing at a two-phase region temperature of a ferrite phase and an austenite phase, which is one of the technical features of the present invention, to thereby obtain an elongation at break and an average r It has been found that beneficial effects are also obtained with respect to value and ridging resistance.
  • hot-rolled sheet annealing was performed at a ferrite single-phase temperature, but in the present invention, hot-rolled sheet annealing is performed at a high temperature that is a two-phase region of a ferrite phase and an austenite phase. Grain growth is further promoted and the crystal grain size is increased appropriately.
  • stimulation of the development of an annealing texture are acquired.
  • the elongation at break is also improved for the following reasons.
  • the amount of carbonitride produced after cold-rolled sheet annealing is reduced, and the generation of voids and the connection of voids during tensile deformation are suppressed. . This also improves the elongation at break.
  • the reason why a beneficial effect can be obtained with respect to ridging resistance is as follows.
  • the austenite phase is generated with a crystal orientation different from that of the ferrite phase before annealing.
  • the metal structure after hot-rolled sheet annealing becomes a two-phase structure of a martensite phase and a ferrite phase.
  • rolling strain is locally concentrated in the ferrite phase sandwiched between the martensite phases, and an orientation difference is formed in the ferrite phase.
  • recrystallization occurs preferentially at a site where the orientation difference is introduced in the subsequent cold-rolled sheet annealing.
  • the ferrite phase colonies are effectively destroyed, and excellent ridging resistance with a ridging height of 2.5 ⁇ m or less is obtained.
  • the steel components have a C content and an N content generated by the austenite phase.
  • the C content and the N content are reduced within a range in which a predetermined amount of austenite phase can be generated.
  • About steel which has such a component after performing hot-rolled sheet annealing at the two-phase region temperature of a ferrite phase and austenite, cold rolling and cold-rolled sheet annealing are performed. As a result, it is necessary to obtain a ferrite single-phase structure composed of ferrite grains containing a large amount of carbonitride and few ferrite grains.
  • C 0.005 to 0.025%
  • C promotes the formation of the austenite phase and has the effect of expanding the two-phase temperature range where the ferrite phase and the austenite phase appear during hot-rolled sheet annealing.
  • a content of 0.005% or more is necessary.
  • the amount of C exceeds 0.025%, the amount of austenite phase produced in the hot-rolled sheet annealing becomes excessive, and the amount of ferrite grains containing many carbonitrides becomes excessive after cold-rolled sheet annealing.
  • the distance between voids in the metal structure is reduced, and breakage due to void connection is likely to occur at the time of molding, and sufficient stretch formability cannot be obtained. Therefore, the C content is in the range of 0.005 to 0.025%. Preferably it is 0.010 to 0.020% of range.
  • Si 0.02 to 0.50% Si is an element that acts as a deoxidizer during steel melting. In order to acquire this effect, 0.02% or more needs to be contained. However, if the amount of Si exceeds 0.50%, the steel sheet becomes hard, the rolling load during hot rolling increases, and the ductility after finish annealing decreases. For this reason, the Si content is in the range of 0.02 to 0.50%. Preferably it is 0.10 to 0.35% of range. More preferably, it is in the range of 0.10 to 0.20%.
  • Mn 0.55 to 1.00% Mn, like C, promotes the formation of an austenite phase and has the effect of expanding the two-phase temperature range in which a ferrite phase and an austenite phase appear during hot-rolled sheet annealing. In order to acquire this effect, 0.55% or more needs to be contained. However, if the amount of Mn exceeds 1.00%, the amount of MnS produced increases and the corrosion resistance decreases. Therefore, the Mn content is set in the range of 0.55 to 1.00%. Preferably it is 0.60 to 0.90% of range. More preferably, it is in the range of 0.75 to 0.85%.
  • P 0.04% or less Since P is an element that promotes grain boundary fracture due to grain boundary segregation, the lower one is desirable, and the upper limit is made 0.04%. Preferably it is 0.03% or less. More preferably, it is 0.01% or less.
  • S 0.01% or less
  • S is an element that exists as sulfide inclusions such as MnS and lowers ductility, corrosion resistance, etc., and particularly when the content exceeds 0.01%, their adverse effects Is noticeable. For this reason, the S amount is desirably as low as possible.
  • the upper limit of the S amount is set to 0.01%. Preferably it is 0.007% or less. More preferably, it is 0.005% or less.
  • Al 0.001 to 0.10%
  • Al is an element that acts as a deoxidizing agent like Si. In order to acquire this effect, 0.001% or more needs to be contained. However, when the Al content exceeds 0.10%, Al-based inclusions such as Al 2 O 3 increase, and the surface properties tend to decrease. Therefore, the Al content is set in the range of 0.001 to 0.10%. Preferably it is 0.001 to 0.07% of range. More preferably, it is in the range of 0.001 to 0.05%.
  • Cr 15.5 to 18.0% Cr is an element having an effect of improving the corrosion resistance by forming a passive film on the surface of the steel sheet. In order to obtain this effect, the Cr amount needs to be 15.5% or more. However, if the Cr content exceeds 18.0%, the austenite phase is not sufficiently generated during hot-rolled sheet annealing, and desired material characteristics cannot be obtained. Therefore, the Cr content is in the range of 15.5 to 18.0%. Preferably it is 16.0 to 17.0% of range. More preferably, it is in the range of 16.0 to 16.5%.
  • Ni 0.1 to 1.0%
  • Ni is an element that improves corrosion resistance, and it is effective to contain it particularly when high corrosion resistance is required.
  • Ni also has the effect of promoting the formation of the austenite phase and expanding the two-phase temperature range in which the ferrite phase and austenite phase appear during hot-rolled sheet annealing. These effects become significant when the content is 0.1% or more. However, if the Ni content exceeds 1.0%, the formability deteriorates, which is not preferable. Therefore, when Ni is contained, the content is made 0.1 to 1.0%. Preferably it is 0.1 to 0.3% of range.
  • N 0.005 to 0.025%
  • N like C and Mn, promotes the formation of the austenite phase and has the effect of expanding the two-phase temperature range in which the ferrite phase and austenite phase appear during hot-rolled sheet annealing.
  • the N amount needs to be 0.005% or more.
  • the N content exceeds 0.025%, the ductility is remarkably lowered, and the amount of austenite phase generated in hot-rolled sheet annealing becomes excessive, and the amount of ferrite grains that are rich in carbonitride after cold-rolled sheet annealing is increased. Becomes excessive.
  • the N content is set in the range of 0.005 to 0.025%. Preferably it is 0.010 to 0.020% of range.
  • the balance is Fe and inevitable impurities.
  • Cu 0.1 to 1.0%
  • V 0.01 to 0.10%
  • Ti 0.001 to 0.05%
  • Nb 0.001 to 0.05%
  • Mo 0.1 to One or more selected from 0.5%
  • Co 0.01 to 0.2%
  • Cu 0.1 to 1.0%
  • Cu is an element that improves corrosion resistance, and it is effective to contain it particularly when high corrosion resistance is required.
  • Cu has an effect of promoting the generation of an austenite phase and expanding a two-phase temperature range in which a ferrite phase and an austenite phase appear during hot-rolled sheet annealing. These effects become significant when the content is 0.1% or more. However, if the Cu content exceeds 1.0%, formability may be deteriorated, which is not preferable. Therefore, when Cu is contained, the content is made 0.1 to 1.0%. Preferably it is 0.2 to 0.3% of range.
  • V 0.01-0.10% V combines with C and N in the steel to reduce solute C and solute N. This improves the average r value. In order to acquire this effect, it is necessary to contain V amount 0.01% or more. However, if the amount of V exceeds 0.10%, the workability is lowered and the manufacturing cost is increased. Therefore, when V is contained, the content is made 0.01 to 0.10%. Preferably it is 0.02 to 0.08% of range.
  • Ti and Nb are elements having a high affinity with C and N, like V, and precipitate as carbide or nitride during hot rolling, reducing the solid solution C and solid solution N in the matrix, and cold rolling. There is an effect of improving workability after sheet annealing. In order to obtain this effect, it is necessary to contain 0.001% or more of Ti and 0.001% or more of Nb. However, when the Ti content exceeds 0.05% or the Nb content exceeds 0.05%, good surface properties cannot be obtained due to excessive precipitation of TiN and NbC.
  • the range when Ti is contained, the range is 0.001 to 0.05%, and when Nb is contained, the range is 0.001 to 0.05%.
  • the amount of Ti is preferably in the range of 0.003 to 0.03%. More preferably, it is in the range of 0.005 to 0.015%.
  • the amount of Nb is preferably in the range of 0.003 to 0.03%. More preferably, it is in the range of 0.005 to 0.015%.
  • Mo 0.1 to 0.5%
  • Mo is an element that improves corrosion resistance, and it is effective to contain it particularly when high corrosion resistance is required. This effect becomes remarkable when the content is 0.1% or more. However, if the Mo content exceeds 0.5%, the austenite phase is not sufficiently generated during hot-rolled sheet annealing, and desired material characteristics cannot be obtained. Therefore, when it contains Mo, it is 0.1 to 0.5%. Preferably it is 0.2 to 0.3% of range.
  • Co 0.01 to 0.2% Co is an element that improves toughness. This effect is obtained when the content is 0.01% or more. On the other hand, if the content exceeds 0.2%, the moldability is lowered. Therefore, if Co is contained, the content is made 0.01 to 0.2%.
  • Mg 0.0002 to 0.0050%
  • Ca 0.0002 to 0.0020%
  • B 0.0002 to 0.0050%
  • REM 0.01 to 0.10%
  • Mg is an element that has an effect of improving hot workability. In order to acquire this effect, 0.0002% or more needs to be contained. However, when the amount of Mg exceeds 0.0050%, the surface quality deteriorates. Therefore, when Mg is contained, the content is made 0.0002 to 0.0050%. Preferably it is 0.0005 to 0.0035% of range. More preferably, it is in the range of 0.0005 to 0.0020%.
  • Ca 0.0002 to 0.0020%
  • Ca is an effective component for preventing nozzle clogging due to crystallization of inclusions that are likely to occur during continuous casting. In order to acquire the effect, 0.0002% or more needs to be contained. However, if the Ca content exceeds 0.0020%, CaS is generated and the corrosion resistance is lowered. Therefore, when Ca is contained, the content is made 0.0002 to 0.0020%. Preferably it is 0.0005 to 0.0015% of range. More preferably, it is in the range of 0.0005 to 0.0010%.
  • B 0.0002 to 0.0050%
  • B is an element effective for preventing embrittlement at low temperature secondary work. In order to acquire this effect, 0.0002% or more needs to be contained. However, when the amount of B exceeds 0.0050%, hot workability deteriorates. Therefore, when B is contained, the content is made 0.0002 to 0.0050%. Preferably it is 0.0005 to 0.0035% of range. More preferably, it is in the range of 0.0005 to 0.0020%.
  • REM 0.01-0.10% REM (Rare Earth Metals) is an element that improves oxidation resistance, and is particularly effective in suppressing the formation of an oxide film on the welded portion and improving the corrosion resistance of the welded portion.
  • a content of 0.01% or more is necessary. However, if the content exceeds 0.10%, productivity such as pickling at the time of cold rolling annealing is lowered.
  • REM is an expensive element, excessive inclusion is not preferable because it causes an increase in manufacturing cost. Therefore, when REM is contained, the content is made 0.01 to 0.10%. Preferably it is 0.01 to 0.05% of range.
  • the ferritic stainless steel sheet of the present invention is subjected to hot rolling on a steel slab having the above component composition, followed by hot rolling sheet annealing at a temperature range of 900 to 1100 ° C. for 5 seconds to 15 minutes, After cold rolling, it is obtained by annealing a cold-rolled sheet that is held at a temperature range of 800 to 900 ° C. for 5 seconds to 5 minutes.
  • the molten steel having the above component composition is melted by a known method such as a converter, electric furnace, vacuum melting furnace or the like, and is made into a steel material (slab) by a continuous casting method or an ingot-bundling method.
  • the slab is heated at 1100 to 1250 ° C. for 1 to 24 hours, or directly hot-rolled as cast without heating to form a hot-rolled sheet.
  • the winding temperature is preferably 500 ° C. or higher and 850 ° C. or lower. If it is less than 500 degreeC, a martensite phase will produce
  • Hot-rolled sheet annealing held for 5 seconds to 15 minutes in the temperature range of 900 to 1100 ° C. Thereafter, hot rolling for 5 seconds to 15 minutes in the temperature range of 900 to 1100 ° C. which is a two-phase temperature range of the ferrite phase and austenite phase Sheet annealing is performed.
  • Hot-rolled sheet annealing is an extremely important process for the present invention to obtain excellent formability and ridging resistance.
  • the hot-rolled sheet annealing temperature is less than 900 ° C., sufficient recrystallization does not occur and the ferrite single-phase region is obtained, so that the effects of the present invention that are manifested by annealing in the two-phase temperature region may not be obtained.
  • the annealing temperature exceeds 1100 ° C.
  • the amount of austenite phase produced is significantly reduced, and the predetermined ridging resistance may not be obtained.
  • the annealing time is less than 5 seconds, even if annealing is performed at a predetermined temperature, generation of austenite phase and recrystallization of the ferrite phase do not occur sufficiently, so that desired formability may not be obtained.
  • the annealing time exceeds 15 minutes, C concentration in the austenite phase is promoted and the martensite phase becomes excessively hard.
  • the hot-rolled sheet annealing is held in the temperature range of 900 to 1100 ° C. for 5 seconds to 15 minutes.
  • the temperature is maintained at 920 to 1080 ° C. for 15 seconds to 5 minutes. More preferably, the temperature is kept at 940 to 1040 ° C. for 30 seconds to 3 minutes.
  • cold rolling is preferably performed at a rolling reduction of 50% or more from the viewpoints of extensibility, bendability, press formability, and shape correction.
  • cold rolling and annealing may be repeated twice or more. Further, in order to improve the surface properties, grinding or polishing may be performed.
  • Cold-rolled sheet annealing is carried out for 5 seconds to 5 minutes in a temperature range of 800 to 900 ° C. Next, cold-rolled sheet annealing is performed.
  • Cold-rolled sheet annealing is an important process for making a two-phase structure of a ferrite phase and a martensite phase formed by hot-rolled sheet annealing into a ferrite single-phase structure. If the cold-rolled sheet annealing temperature is less than 800 ° C., sufficient recrystallization does not occur and a predetermined formability cannot be obtained. On the other hand, when the cold-rolled sheet annealing temperature exceeds 900 ° C., the steel component in which the temperature exceeding 900 ° C.
  • the ferrite phase becomes the two-phase temperature range of the ferrite phase and the austenite phase generates a martensite phase after the cold-rolled sheet annealing. Becomes hard, and the predetermined elongation at break and stretchability cannot be obtained. Moreover, even if it is a steel component in which the temperature exceeding 900 ° C. is the ferrite single phase temperature range, the glossiness of the steel sheet is lowered due to the remarkable coarsening of crystal grains, which is not preferable from the viewpoint of surface quality.
  • the annealing time is less than 5 seconds, even if annealing is performed at a predetermined temperature, the ferrite phase is not sufficiently recrystallized, and therefore, a predetermined formability cannot be obtained.
  • cold-rolled sheet annealing is held for 5 seconds to 5 minutes in a temperature range of 800 to 900 ° C.
  • the temperature is maintained at 850 ° C. to 900 ° C. for 15 seconds to 3 minutes.
  • BA annealing (bright annealing) may be performed.
  • Stainless steel having the chemical composition shown in Table 1 was melted in a 50 kg small vacuum melting furnace. These steel ingots were heated at 1150 ° C. for 1 hour and then hot rolled to form hot rolled sheets having a thickness of 3.5 mm. Subsequently, these hot-rolled sheets were subjected to hot-rolled sheet annealing under the conditions shown in Table 2, and then the surfaces were descaled by shot blasting and pickling. Further, after cold rolling to a plate thickness of 0.8 mm, cold rolled sheet annealing was performed under the conditions shown in Table 2. Furthermore, descaling treatment by pickling was performed to obtain a cold-rolled pickling annealed plate (ferritic stainless steel plate).
  • the cold roll pickling annealed plate (ferritic stainless steel plate) thus obtained was evaluated as follows.
  • Photograph the surface of the specimen after 8 cycles of salt spray cycle test measure the rusting area on the specimen surface by image analysis, and calculate the rusting rate (( Rust area / total area of test piece) ⁇ 100 [%]) was calculated.
  • a rusting rate of 10% or less was determined to pass with excellent corrosion resistance ()), more than 10% to 25% or less passed ( ⁇ ), and more than 25% to reject (x).
  • the Cr content falls below the scope of the present invention. With 38 (steel No. S30), although predetermined formability and ridging characteristics were obtained, the predetermined corrosion resistance was not obtained because the Cr content was insufficient.
  • No. C content is below the scope of the present invention.
  • No. 33 (steel No. S25), a predetermined elongation at break and an average r value were obtained, but since austenite generation ability was insufficient, an austenite phase was not formed in hot-rolled sheet annealing, and predetermined ridging resistance and overhanging properties were obtained. Formability could not be obtained.
  • the C content exceeds the range of the present invention.
  • Predetermined ridging resistance and stretch formability were obtained, but the steel sheet was hardened, so that the elongation was lowered and the predetermined breaking elongation was not obtained.
  • No. 27 (steel No. S27) was hardened by excessive Si content, and could not obtain a predetermined elongation at break.
  • No. N content is below the range of the present invention.
  • 40 steel No. S32
  • a predetermined elongation at break and an average r value were obtained, but since austenite forming ability was insufficient, an austenite phase was not formed in hot-rolled sheet annealing, and a predetermined ridging resistance and overhanging property were obtained. Formability could not be obtained.
  • the N content exceeds the range of the present invention.
  • No. 41 (steel No. S33), predetermined ridging resistance characteristics and stretch formability were obtained, but the predetermined breaking elongation was not obtained because the steel plate was hardened. Furthermore, sensitization caused by the precipitation of a large amount of Cr nitride in the structure occurred, and the predetermined corrosion resistance could not be obtained.
  • the ferritic stainless steel sheet obtained by the present invention is particularly suitable for applications requiring press-formed products mainly composed of overhang forming, such as kitchen appliances and tableware.

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Abstract

Provided are a ferritic stainless steel sheet having adequate corrosion resistance and excellent forming properties and ridging resistance, and a method for manufacturing the same. This ferritic stainless steel sheet contains, in terms of % by mass, 0.005-0.025% C, 0.02-0.50% Si, 0.55-1.00% Mn, 0.04% or less of P, 0.01% or less of S, 0.001-0.10% Al, 15.5-18.0% Cr, 0.1-1.0% Ni, and 0.005-0.025% N, the remainder comprising Fe and unavoidable impurities, the breaking elongation of the ferritic stainless steel sheet being 28% or greater, the average r value thereof being 0.75 or greater, and the minimum value of the maximum logarithmic strain of the forming limit based on a forming limit diagram (FLD) being 0.15 or greater.

Description

フェライト系ステンレス鋼板およびその製造方法Ferritic stainless steel sheet and manufacturing method thereof
　本発明は、十分な耐食性を有し、成形性および耐リジング性に優れたフェライト系ステンレス鋼板およびその製造方法に関するものである。 The present invention relates to a ferritic stainless steel sheet having sufficient corrosion resistance and excellent formability and ridging resistance, and a method for producing the same.
　フェライト系ステンレス鋼板は、高価なNiを多量に含むオーステナイト系ステンレス鋼より経済的である。フェライト系ステンレス鋼の中でも特にＳＵＳ４３０系ステンレス鋼板（1６～1８質量％Ｃｒ）は経済的なため、建材、輸送機器、家電製品、厨房器具、または、自動車部品などのさまざまな用途に使用されており、その適用範囲は近年さらに拡大しつつある。これらの用途に適用するためには、耐食性だけでなく、所定の形状に加工できる十分な成形性が求められる。 Ferritic stainless steel sheet is more economical than austenitic stainless steel containing a large amount of expensive Ni. Among ferritic stainless steels, SUS430 stainless steel plate (16 to 18% by mass Cr) is particularly economical and is used in various applications such as building materials, transportation equipment, home appliances, kitchen appliances, and automobile parts. The range of application has been further expanded in recent years. In order to apply to these applications, not only corrosion resistance but also sufficient formability that can be processed into a predetermined shape is required.
　一方、ＳＵＳ４３０ 系ステンレス鋼板では外観が良好であることが求められる用途へ適用される場合が多く、耐リジング特性に優れることも必要とされる。リジングとは成形加工のひずみに起因して発生する表面凹凸のことである。フェライト系ステンレス鋼板では鋳造および／または熱延時に類似した結晶方位を有する結晶粒群（コロニー）が生成する場合がある。コロニーが残存する鋼板では成形加工時にコロニー部とその他の部位でひずみ量に大きな差が生じるために、成形後に表面凹凸（リジング）が発生する。成形後に過度のリジングが発生した場合、表面凹凸を除去するために研磨工程が必要となり成形品の製造コストが上昇するという問題がある。 On the other hand, SUS430 stainless steel sheet is often applied to applications that require a good appearance, and is required to have excellent ridging resistance. Ridging is a surface irregularity generated due to distortion in molding. In a ferritic stainless steel sheet, a crystal grain group (colony) having a similar crystal orientation may be generated during casting and / or hot rolling. In a steel sheet in which colonies remain, a large difference occurs in the strain amount between the colony part and other parts at the time of forming, and thus surface irregularities (ridging) occur after forming. When excessive ridging occurs after molding, there is a problem that a polishing step is required to remove surface irregularities, and the manufacturing cost of the molded product increases.
　特許文献１では、質量％で、Ｃ:０．０２～０．０６％、Ｓｉ：１．０％以下、Ｍｎ:１．０％以下、Ｐ:０．０５％以下、Ｓ:０．０１％以下、Ａｌ:０．００５％以下、Ｔｉ:０．００５％以下、Ｃｒ:１１～３０％、Ｎｉ:０．７％以下を含み、かつ０．０６≦（Ｃ＋Ｎ）≦０．１２、１≦Ｎ／Ｃおよび１．５×１０－３≦（Ｖ×Ｎ）≦１．５×１０－２（Ｃ、Ｎ、Ｖはそれぞれ各元素の質量％を表す）を満たすことを特徴とする成形性に優れたフェライト系ステンレス鋼が開示されている。しかし、本発明者らが特許文献１に記載の手法でフェライト系ステンレス鋼を製造したところ、鋼板の圧延方向については優れた破断伸びが得られた。しかし、プレス加工により張出成形性を主体とした排気ダクトの作製を試みたところ、所定の形状に成形することができず、破断伸びから期待されるほどの張出成形性が得られなかった。さらに、特許文献１の実施例では熱間圧延後にいわゆる箱焼鈍（例えば、８６０℃で８時間の焼鈍）を行っている。このような箱焼鈍は加熱や冷却の過程を含めると一週間程度掛かり、生産性が低いという問題がある。また、高額な遷移金属元素であるＶを添加することによる固溶Ｎの低減技術を用いているため、製造コストが高くなるという問題もある。さらに、熱延板焼鈍をフェライト単相温度域で箱焼鈍を行っているために、フェライトコロニーがほとんど破壊されずに残存するため、耐リジング性が著しく低下するという問題もある。 In Patent Document 1, by mass, C: 0.02 to 0.06%, Si: 1.0% or less, Mn: 1.0% or less, P: 0.05% or less, S: 0.01% Hereinafter, Al: 0.005% or less, Ti: 0.005% or less, Cr: 11 to 30%, Ni: 0.7% or less, and 0.06 ≦ (C + N) ≦ 0.12, 1 ≦ N / C and 1.5 × 10 −3 ≦ (V × N) ≦ 1.5 × 10 −2 (C, N, and V each represent mass% of each element) An excellent ferritic stainless steel is disclosed. However, when the present inventors manufactured ferritic stainless steel by the method described in Patent Document 1, excellent elongation at break was obtained in the rolling direction of the steel sheet. However, when an attempt was made to produce an exhaust duct mainly composed of bulging formability by press working, it was not possible to form the predetermined shape, and the bulging formability expected from the elongation at break could not be obtained. . Furthermore, in the Example of patent document 1, what is called box annealing (for example, annealing for 8 hours at 860 degreeC) is performed after hot rolling. Such box annealing has a problem of low productivity because it takes about one week when heating and cooling processes are included. In addition, there is also a problem that the manufacturing cost increases because a technique for reducing solid solution N by adding V, which is an expensive transition metal element, is used. Further, since box annealing is performed in the single-phase temperature range of ferrite for hot-rolled sheet annealing, ferrite colonies remain with almost no destruction, and there is also a problem that ridging resistance is significantly lowered.
　特許文献２では、質量％で、Ｃ：０．０１～０．１０％、Ｓｉ：０．０５～０．５０％、Ｍｎ：０．０５～１．００％、Ｎｉ：０．０１～０．５０％、Ｃｒ：１０～２０％、Ｍｏ：０．００５～０．５０％、Ｃｕ：０．０１～０．５０％、Ｖ：０．００１～０．５０％、Ｔｉ：０．００１～０．５０％、Ａｌ：０．０１～０．２０％、Ｎｂ：０．００１～０．５０％、Ｎ：０．００５～０．０５０％およびＢ：０．０００１０～０．００５００％を含有した鋼を熱間圧延後、箱型炉あるいはＡＰライン(Ａｎｎｅａｌｉｎｇ　ａｎｄ　Ｐｉｃｋｌｉｎｇ　ｌｉｎｅ)の連続炉を用いてフェライト単相温度域で熱延板焼鈍を行い、さらに冷間圧延および仕上げ焼鈍を行うことを特徴とする加工性と表面性状に優れたフェライト系ステンレス鋼が開示されている。しかし、箱型炉を用いた場合には上記の特許文献１と同様に生産性が低いという問題がある。これに加えて、特許文献２でも特許文献１と同様に張出成形を主体とするプレス部材の作製を試みたところ、所定の形状に成形することができず、破断伸びから期待されるほどの張出成形性が得られなかった。さらに、一般に特許文献２のようなフェライト系ステンレス鋼は、鋳造あるいは熱間圧延時に類似した結晶方位を有する結晶粒群（コロニー）が生成し、熱延板焼鈍をフェライト単相温度域で行うとフェライト相のコロニーを十分に破壊することができない。そのため、コロニーは熱延板焼鈍後の冷間圧延によって圧延方向に展伸して残存し、成形後に著しいリジングが生じるという問題がある。 In Patent Document 2, by mass, C: 0.01 to 0.10%, Si: 0.05 to 0.50%, Mn: 0.05 to 1.00%, Ni: 0.01 to 0.00. 50%, Cr: 10 to 20%, Mo: 0.005 to 0.50%, Cu: 0.01 to 0.50%, V: 0.001 to 0.50%, Ti: 0.001 to 0 .50%, Al: 0.01-0.20%, Nb: 0.001-0.50%, N: 0.005-0.050% and B: 0.00010-0.00500% After hot-rolling steel, hot-rolled sheet annealing is performed in the ferrite single-phase temperature range using a box furnace or a continuous furnace of AP line (Annealing and Pickling line), followed by cold rolling and finish annealing. Ferritic stainless steel with excellent workability and surface properties It is shown. However, when a box furnace is used, there is a problem that productivity is low as in Patent Document 1 described above. In addition to this, in Patent Document 2, as in Patent Document 1, when an attempt was made to produce a press member mainly composed of stretch forming, it could not be formed into a predetermined shape, and as expected from elongation at break. No stretch formability was obtained. Furthermore, in general, ferritic stainless steel as in Patent Document 2 generates a group of crystal grains (colony) having a similar crystal orientation during casting or hot rolling, and hot-rolled sheet annealing is performed in a ferrite single-phase temperature range. The ferrite phase colonies cannot be sufficiently destroyed. Therefore, there is a problem that the colony expands and remains in the rolling direction by cold rolling after hot-rolled sheet annealing, and significant ridging occurs after forming.
　特許文献３では、０．１５％以下のＣ、１３～２５％のＣｒを含有するフェライト系ステンレス鋼であって、この鋼の熱延板をオーステナイトおよびフェライト相が共存する９３０～９９０℃の範囲で１０分以内の焼鈍を行うことにより、組織をマルテンサイト相とフェライト相の二相組織とし、次いで、冷間圧延および冷延板焼鈍を行うことを特徴とする耐リジング性と加工性に優れるフェライト系ステンレス鋼板の製造方法が開示されている。特許文献３では、加工性として伸びにのみ言及している。しかしながら、本発明者らが特許文献３に記載の方法で鋼板を製造し、張出成形を主体とした換気フードの作製を試みたところ、プレス加工中に割れが生じ所定の形状に成形できないことが散発し、破断伸びから期待されるほどの張出成形性が発現しない場合があることが明らかとなった。このように、特許文献３に記載のフェライト系ステンレス鋼板では引張試験における破断伸びは高いものの、プレス成形において求められる張出成形性を十分に発現させることができず、本発明が課題とする十分な成形性が得られているとは言えない。 In Patent Document 3, a ferritic stainless steel containing 0.15% or less of C and 13 to 25% of Cr, and a range of 930 to 990 ° C. in which austenite and a ferrite phase coexist on the hot-rolled sheet of this steel. By carrying out annealing within 10 minutes, the structure becomes a two-phase structure of martensite phase and ferrite phase, and then cold rolling and cold-rolled sheet annealing are performed, which is excellent in ridging resistance and workability A method for producing a ferritic stainless steel sheet is disclosed. In Patent Document 3, only elongation is mentioned as workability. However, when the inventors of the present invention manufactured a steel plate by the method described in Patent Document 3 and attempted to produce a ventilation hood mainly composed of overhang forming, cracks occurred during press working and it could not be formed into a predetermined shape. It was clarified that the stretch formability as expected from the elongation at break may not appear. As described above, the ferritic stainless steel sheet described in Patent Document 3 has a high elongation at break in a tensile test, but cannot fully exhibit the stretch formability required in press forming, and is a sufficient problem that the present invention has a problem. It cannot be said that a good moldability is obtained.
　以上のように、十分な耐食性を有し、成形性および耐リジング性に優れたＳＵＳ４３０ 系ステンレス鋼板を、生産する技術は確立されていない。 As described above, the technology for producing a SUS430 stainless steel plate having sufficient corrosion resistance and excellent formability and ridging resistance has not been established.
特許第３５８４８８１号公報(再公表ＷＯ００/６０１３４号)Japanese Patent No. 3584881 (Republished WO 00/60134) 特許第３５８１８０１号公報(特開２００１－３１３４号)Japanese Patent No. 3581801 (Japanese Patent Laid-Open No. 2001-3134) 特公昭４７－１８７８号公報Japanese Patent Publication No.47-1878
　本発明は、かかる課題を解決し、十分な耐食性を有し、成形性および耐リジング性に優れたフェライト系ステンレス鋼板およびその製造方法を提供することを目的とする。 An object of the present invention is to solve such problems and to provide a ferritic stainless steel sheet having sufficient corrosion resistance, excellent formability and ridging resistance, and a method for producing the same.
　なお、本発明において、十分な耐食性とは、表面を#６００エメリーペーパーにより研磨仕上げした後に端面部をシールした鋼板にＪＩＳ　Ｈ　８５０２に規定された塩水噴霧サイクル試験（（塩水噴霧（３５℃、５質量％ＮａＣｌ、噴霧２ｈｒ）→乾燥（６０℃、相対湿度４０％、４ｈｒ）→湿潤（５０℃、相対湿度≧９５％、２ｈｒ））を１サイクルとする試験）を８サイクル行った場合の鋼板表面における発錆面積率（＝発錆面積／鋼板全面積×１００ [％]）が２５％以下であることを意味する。 In the present invention, sufficient corrosion resistance refers to a salt spray cycle test (salt spray (35 ° C., 5 ° C., JIS H 8502) applied to a steel plate whose surface is polished with # 600 emery paper and the end face is sealed. (Mass% NaCl, spraying 2 hr) → drying (60 ° C., relative humidity 40%, 4 hr) → wetting (50 ° C., relative humidity ≧ 95%, 2 hr)))) It means that the rusting area ratio on the surface (= rusting area / total area of steel sheet × 100% [%]) is 25% or less.
　また、優れた成形性とは、優れた張出成形性、破断伸び、および平均ｒ値を有することを言う。優れた張出成形性とは、鋼の成形限界線図（Ｆｏｒｍｉｎｇ　Ｌｉｍｉｔ　Ｄｉａｇｒａｍ、ＦＬＤ）に基づいて決定される成形限界の最大対数ひずみの最小値が０．１５以上であることを意味する。優れた破断伸びとは、ＪＩＳ　Ｚ　２２４１に準拠した引張試験における破断伸び（Ｅｌ）が圧延方向と直角方向の試験片で２８％以上であることを意味する。優れた平均ｒ値とは、ＪＩＳ　Ｚ　２２４１に準拠した引張試験において１５％のひずみを付与した際の下記（１）式により算出される平均ランクフォード値（以下、平均ｒ値と称す）が０.７５以上であることを意味する。
平均ｒ値＝（ｒL＋２×ｒD＋ｒC）／４　　　（１）
ここで、ｒLは圧延方向に平行な方向に引張試験した際のｒ値、ｒDは圧延方向に対して４５°の方向に引張試験した際のｒ値、ｒCは圧延方向と直角方向に引張試験した際のｒ値である。 Further, excellent moldability means having excellent stretch formability, elongation at break, and average r value. Excellent stretch formability means that the minimum value of the maximum logarithmic strain of the forming limit determined based on the forming limit diagram (FLD) of steel is 0.15 or more. The excellent elongation at break means that the elongation at break (El) in a tensile test according to JIS Z 2241 is 28% or more in a test piece perpendicular to the rolling direction. An excellent average r value is an average rankford value (hereinafter referred to as an average r value) calculated by the following equation (1) when a strain of 15% is applied in a tensile test based on JIS Z 2241: It means that it is more than .75.
Average r value = (r L + 2 × r D + r C ) / 4 (1)
Here, r L is an r value when a tensile test is performed in a direction parallel to the rolling direction, r D is an r value when a tensile test is performed in a direction of 45 ° with respect to the rolling direction, and r C is a direction perpendicular to the rolling direction. The r value when a tensile test is performed.
　さらに、優れた耐リジング特性とは、以下の方法で測定したリジング高さが２．５μｍ以下であることを意味する。リジング高さの測定は、まず、圧延方向に平行にＪＩＳ　５号引張試験片を採取する。次いで、採取した試験片の表面を＃６００のエメリーペーパーを用いて研磨した後、２０％の引張ひずみを付与する。次いで、試験片の平行部中央の研磨面で、圧延方向に直角の方向に、表面粗度計でＪＩＳ　Ｂ　０６０１（２００１年）で規定される算術平均うねり（Ｗａ）を測定する。測定条件は、測定長１６ｍｍ、ハイカットフィルター波長０．８ｍｍ、ローカットフィルター波長８ｍｍである。この算術平均うねりをリジング高さとする。 Furthermore, excellent ridging resistance means that the ridging height measured by the following method is 2.5 μm or less. For measuring the ridging height, first, a JIS No. 5 tensile test piece is taken in parallel with the rolling direction. Next, after polishing the surface of the collected test piece using # 600 emery paper, 20% tensile strain is applied. Next, the arithmetic average waviness (Wa) defined by JIS B 0601 (2001) is measured with a surface roughness meter in the direction perpendicular to the rolling direction on the polishing surface at the center of the parallel part of the test piece. The measurement conditions are a measurement length of 16 mm, a high cut filter wavelength of 0.8 mm, and a low cut filter wavelength of 8 mm. This arithmetic mean swell is defined as the ridging height.
　課題を解決するために検討した。その結果、以下の知見を得た。適切な成分のフェライト系ステンレス鋼板に対して、熱間圧延後、冷間圧延する前に、フェライト相とオーステナイト相の二相域の好適な温度域で焼鈍（以下、熱延板焼鈍と称する）を行い、さらに、冷間圧延後の鋼板をフェライト単相域となる温度で焼鈍（以下、冷延板焼鈍と称する）を行うことにより、フェライト単相組織ではあるが、粒内の炭窒化物が多いフェライト粒と粒内の炭窒化物が少ないフェライト粒の混粒組織とする。その結果、十分な耐食性を有し、成形性および耐リジング特性に優れたフェライト系ステンレス鋼板が得られることを見出した。 Investigated to solve the problem. As a result, the following knowledge was obtained. An appropriate component ferritic stainless steel sheet is annealed at a suitable temperature range of a ferrite phase and an austenite phase after hot rolling and before cold rolling (hereinafter referred to as hot rolled sheet annealing). Further, by annealing the cold-rolled steel sheet at a temperature that becomes a ferrite single-phase region (hereinafter referred to as cold-rolled sheet annealing), it is a ferrite single-phase structure, but an intragranular carbonitride. A mixed grain structure of ferrite grains having a large amount of ferrite and ferrite grains having a small amount of carbonitride in the grains. As a result, it has been found that a ferritic stainless steel sheet having sufficient corrosion resistance and excellent formability and ridging resistance can be obtained.
　本発明は以上の知見に基づいてなされたものであり、以下を要旨とするものである。
［１］質量％で、Ｃ：０．００５～０．０２５％、Ｓｉ：０．０２～０．５０％、Ｍｎ：０．５５～１．００％、Ｐ：０．０４％以下、Ｓ：０．０１％以下、Ａｌ：０．００１～０．１０％、Ｃｒ：１５．５～１８．０％、Ｎｉ：０．１～１．０％、Ｎ：０．００５～０．０２５％を含有し、残部がＦｅおよび不可避的不純物からなり、破断伸びが２８％以上、平均r値が０．７５以上、かつ、ＦＬＤ（成形限界線図）に基づく成形限界の最大対数ひずみの最小値が０．１５以上であるフェライト系ステンレス鋼板。
［２］質量％で、さらに、Ｃｕ：０．１～１．０％、Ｖ：０．０１～０．１０％、Ｔｉ：０．００１～０．０５％、Ｎｂ：０．００１～０．０５％、Ｍｏ：０．１～０．５％、Ｃｏ：０．０１～０．２％のうちから選ばれる１種または２種以上を含む上記［１］に記載のフェライト系ステンレス鋼板。
［３］質量％で、さらに、Ｍｇ：０．０００２～０．００５０％、Ｃａ：０．０００２～０．００２０％、Ｂ：０．０００２～０．００５０％、ＲＥＭ：０．０１～０．１０％のうちから選ばれる１種または２種以上を含む上記［１］または［２］に記載のフェライト系ステンレス鋼板。
［４］上記［１］～［３］のいずれかに記載のフェライト系ステンレス鋼板の製造方法であって、鋼スラブに対して、熱間圧延を施した後、９００～１１００℃の温度範囲で５秒～１５分間保持する焼鈍を行い、次いで冷間圧延を施した後、８００～９００℃の温度範囲で５秒～５分間保持する焼鈍を行うフェライト系ステンレス鋼板の製造方法。
なお、本明細書において、鋼の成分を示す％はすべて質量％である。 This invention is made | formed based on the above knowledge, and makes the following a summary.
[1] By mass%, C: 0.005 to 0.025%, Si: 0.02 to 0.50%, Mn: 0.55 to 1.00%, P: 0.04% or less, S: 0.01% or less, Al: 0.001 to 0.10%, Cr: 15.5 to 18.0%, Ni: 0.1 to 1.0%, N: 0.005 to 0.025% And the balance is Fe and inevitable impurities, the elongation at break is 28% or more, the average r value is 0.75 or more, and the minimum value of the maximum logarithmic strain at the forming limit based on the FLD (molding limit diagram) is Ferritic stainless steel sheet that is 0.15 or more.
[2] By mass%, Cu: 0.1-1.0%, V: 0.01-0.10%, Ti: 0.001-0.05%, Nb: 0.001-0. The ferritic stainless steel sheet according to the above [1], including one or more selected from 05%, Mo: 0.1 to 0.5%, and Co: 0.01 to 0.2%.
[3] By mass%, Mg: 0.0002 to 0.0050%, Ca: 0.0002 to 0.0020%, B: 0.0002 to 0.0050%, REM: 0.01 to 0.00. The ferritic stainless steel sheet according to the above [1] or [2], containing one or more selected from 10%.
[4] A method for producing a ferritic stainless steel sheet according to any one of the above [1] to [3], wherein the steel slab is hot-rolled and then subjected to a temperature range of 900 to 1100 ° C. A method for producing a ferritic stainless steel sheet that is annealed for 5 seconds to 15 minutes, then cold-rolled, and then annealed at a temperature range of 800 to 900 ° C. for 5 seconds to 5 minutes.
In addition, in this specification, all% which shows the component of steel is the mass%.
　本発明によれば、十分な耐食性を有し、成形性および耐リジング性に優れたフェライト系ステンレス鋼板が得られる。 According to the present invention, a ferritic stainless steel sheet having sufficient corrosion resistance and excellent formability and ridging resistance can be obtained.
　以下、本発明を詳細に説明する。 Hereinafter, the present invention will be described in detail.
　本発明のフェライト系ステンレス鋼板は、プレス加工で建材部品、家電製品の部品、厨房器具、または、自動車部品などのさまざまな用途に使用されることを目的としている。これらの用途に適用するためには、十分な成形性が求められる。 The ferritic stainless steel sheet of the present invention is intended to be used in various applications such as building material parts, home appliance parts, kitchen appliances, or automobile parts by press working. In order to apply to these uses, sufficient moldability is required.
　しかし、十分な耐食性と優れた成形性および優れた耐リジング特性を同時に満足するＳＵＳ４３０系フェライト系ステンレス鋼の製造技術は十分には確立されていないのが現状である。 However, at present, the manufacturing technology of SUS430 ferritic stainless steel that satisfies both sufficient corrosion resistance, excellent formability, and excellent ridging resistance properties has not been established.
　そこで、本発明者らは成分や製造方法が異なる各種フェライト系ステンレス鋼板（特許文献１～３に該当するものも含まれる）を用いて換気フードを想定した張出成形試験を行った。その結果、破断伸びが高い鋼板でも破断伸びが低い鋼板に比べて張出成形性劣る場合があり、張出成形性の優劣が必ずしも破断伸びの大きさでは決まらないことが明らかとなった。そこで、上記張出成形試験に用いた鋼板について、ＦＬＤ（成形限界線図）を作成して張出成形性を詳細に評価したところ、上記換気フード想定の張出成形で良好な成形性を得るにはＦＬＤに基づく成形限界の最大対数ひずみの最小値で０．１５以上、好ましくは０．１８以上の張出成形性が必要であることが明らかとなった。 Therefore, the present inventors conducted an overhang forming test assuming a ventilation hood using various ferritic stainless steel sheets (including those corresponding to Patent Documents 1 to 3) having different components and production methods. As a result, it has been clarified that even if a steel sheet having a high elongation at break is inferior in stretch formability compared with a steel sheet having a low elongation at break, the superiority or inferiority of the stretch formability is not necessarily determined by the magnitude of the elongation at break. Therefore, when the steel sheet used in the bulge forming test was prepared and the bulge formability was evaluated in detail by creating an FLD (formation limit diagram), good formability was obtained by the bulge hood assumed above. It has become clear that a stretchable formability of 0.15 or more, preferably 0.18 or more is required as the minimum value of the maximum logarithmic strain at the forming limit based on FLD.
　次に、本発明者らは、従来の技術によって得られたフェライト系ステンレス鋼板の張出成形性の優劣が破断伸びの大小と対応しない場合が生じる原因を調査した。その結果、箱焼鈍または連続焼鈍を用いた従来技術の場合、冷延焼鈍後の組織がいずれも炭窒化物が多量にかつ均一に分散したフェライト単相組織であり、これが原因であることを突き止めた。鋼板を加工した場合、ひずみ量の増大に伴って組織中にボイドが生成し、このボイドが連結することで亀裂となって最終的に破断へと至る。このボイドは金属組織中の炭窒化物を起点として生成するため、従来の技術によって得られたフェライト系ステンレス鋼板では炭窒化物が多量にかつ均一に分散したフェライト単相組織であるため、金属組織全面から極めて多量のボイドが生成する。すなわち、従来の技術では、ボイドの連結に起因した亀裂が発生しやすい。その結果、引張試験のような単軸変形においては高い破断伸びを示していても、多軸的な応力およびひずみが加わる張出成形では、全方位的にボイドの連結が生じるために破断が生じやすく、十分な張出成形性が得られない場合があることを見出した。 Next, the present inventors investigated the cause of the case where the superiority or inferiority of the stretch formability of the ferritic stainless steel plate obtained by the conventional technique does not correspond to the magnitude of the elongation at break. As a result, in the case of the prior art using box annealing or continuous annealing, the structure after cold rolling annealing is a ferrite single-phase structure in which carbonitrides are dispersed in a large amount and uniformly, and this is the cause. It was. When a steel plate is processed, voids are generated in the structure as the amount of strain increases, and when these voids are connected, they become cracks and eventually break. Since these voids are generated starting from carbonitride in the metal structure, the ferritic stainless steel sheet obtained by the conventional technique is a ferrite single-phase structure in which carbonitride is dispersed in a large amount and uniformly. A very large amount of voids are generated from the entire surface. That is, in the prior art, cracks due to the connection of voids are likely to occur. As a result, even in the case of uniaxial deformation such as a tensile test, rupture occurs because void connection occurs in all directions in stretch forming in which multiaxial stress and strain are applied, even though it shows high elongation at break. It was easy to find out that sufficient stretchability could not be obtained.
　そこで本発明者らは、適切な成分の鋼板に対し熱延板焼鈍をフェライト相とオーステナイト相の二相域で行った後に、常法で冷間圧延を行い、さらに冷延板焼鈍をフェライト単相温度域で行い、最終的に再度フェライト単相組織とする技術を考案した。この技術により、本発明が目標とする優れた張出成形性、破断伸び、平均ｒ値および耐リジング性の全てを満たすことができることを見出した。 Accordingly, the present inventors performed hot rolling sheet annealing on a steel sheet having an appropriate component in a two-phase region of a ferrite phase and an austenite phase, and then cold-rolled the steel sheet in a conventional manner, and further performed cold rolling sheet annealing on a single ferrite sheet. We devised a technique to perform in the phase temperature range and finally to make a single phase ferrite structure again. It has been found that this technique can satisfy all of the excellent stretch formability, break elongation, average r value and ridging resistance which are the targets of the present invention.
　以下、得られた知見を基に、詳細に説明する。 The following is a detailed explanation based on the obtained knowledge.
　熱延板焼鈍をフェライト単相温度域よりも高温のフェライト相とオーステナイトの二相域で行うことにより、熱延板焼鈍にて面積率で３～２０％のオーステナイト相が生成する。このオーステナイト相は熱延板焼鈍後の冷却過程においてほぼ全てがマルテンサイト相へと変態する。フェライト相とマルテンサイト相からなる二相組織を冷間圧延および冷延板焼鈍した場合、冷延板焼鈍においてマルテンサイト相がフェライト相と炭窒化物に分解する。この組織変化により、冷延板焼鈍後の組織は当初からフェライト相であったフェライト粒とマルテンサイト相の分解によって生成したフェライト粒となる。すなわち、マルテンサイト相の分解によって生成したフェライト相の粒界および粒内には多量の炭窒化物が存在しており、金属組織全体では、粒内および粒界上の炭窒化物が極めて多いフェライト粒と炭窒化物が少ないフェライト粒からなる混粒組織となる。炭窒化物が多いフェライト粒と炭窒化物が少ないフェライト粒の間では、炭窒化物が多いフェライト粒の方が相対的に硬質となり、金属組織中に粒単位の硬度差が生じる。このような鋼板を張出成形した場合、ボイドは主に炭窒化物が多いフェライト粒と少ないフェライト粒の界面上の炭窒化物から生成し、その他の部位におけるボイド発生量は少ないことを見出した。すなわち、本発明鋼では、炭窒化物が多いフェライト粒が連続して位置している部位、炭窒化物が少ないフェライト粒が連続している部位、およびフェライト粒内ではボイド生成が少ない。そのため、従来技術によって得られるフェライト系ステンレス鋼板に比べてボイド間距離が長くなり、張出成形時のボイド連結に起因した亀裂が発生しにくく、ＦＬＤに基づく成形限界の最大対数ひずみの最小値が０．１５以上という高い張出成形性が発現する。 By performing hot-rolled sheet annealing in a two-phase region of ferrite phase and austenite that is higher than the ferrite single-phase temperature range, an austenite phase with an area ratio of 3 to 20% is generated by hot-rolled sheet annealing. Almost all of the austenite phase is transformed into a martensite phase in the cooling process after hot-rolled sheet annealing. When a two-phase structure consisting of a ferrite phase and a martensite phase is cold-rolled and cold-rolled sheet annealed, the martensite phase is decomposed into a ferrite phase and a carbonitride during cold-rolled sheet annealing. Due to this structural change, the structure after the cold-rolled sheet annealing becomes ferrite grains formed by the decomposition of ferrite grains that were ferrite phases and martensite phases from the beginning. That is, there are a large amount of carbonitrides in the grain boundaries and in the grains of the ferrite phase formed by the decomposition of the martensite phase, and in the entire metal structure, ferrite with an extremely large amount of carbonitrides in and on the grain boundaries. It becomes a mixed grain structure composed of ferrite grains with few grains and carbonitrides. Between the ferrite grains having a high carbonitride and the ferrite grains having a low carbonitride, the ferrite grains having a high carbonitride are relatively hard, and a hardness difference of a grain unit occurs in the metal structure. It was found that when such a steel sheet is stretched, voids are mainly generated from carbonitrides on the interface between ferrite grains with a large amount of carbonitride and few ferrite grains, and the amount of voids generated in other parts is small. . That is, in the steel according to the present invention, there is little void formation in a portion where ferrite grains having a large amount of carbonitride are continuously located, a portion where ferrite grains having a small amount of carbonitride are continuous, and a ferrite grain. Therefore, the distance between the voids is longer than that of the ferritic stainless steel plate obtained by the prior art, cracks due to void connection during overhang forming are less likely to occur, and the minimum value of the maximum logarithmic strain at the forming limit based on FLD is small. High stretch formability of 0.15 or more is manifested.
　また、本発明者らがさらに調査したところ、本発明の効果を得るためには、鋼中のＣ含有量およびＮ含有量、ならびに熱延板焼鈍温度を適切に制御することが肝要であることを知見した。すなわち、フェライト相とオーステナイト相の二相域で熱延板焼鈍を行い、３～２０％のオーステナイト相を生成させるためには、オーステナイト生成元素であるＣおよびＮをそれぞれ最低でも０．００５％以上含有させる必要がある。一方、Ｃ含有量およびＮ含有量のいずれかが０．０２５％を超えると、熱延板焼鈍時に生成するオーステナイト相が２０％超と過度に増加する。その結果、その後の冷延板焼鈍によって生成する炭窒化物の多いフェライト粒が増加し、加工時のボイド起点となる炭窒化物が多いフェライト粒と少ないフェライト粒の界面面積が増加するために所定の張出成形性を発現させることができない。よって、Ｃ含有量、Ｎ含有量をそれぞれの上限は０．０２５％とする必要がある。 Moreover, when the present inventors further investigated, in order to acquire the effect of this invention, it is important to control C content and N content in steel, and hot-rolled sheet annealing temperature appropriately. I found out. That is, in order to perform hot-rolled sheet annealing in a two-phase region of a ferrite phase and an austenite phase and to generate an austenite phase of 3 to 20%, at least 0.005% or more of C and N, which are austenite forming elements, respectively. It is necessary to contain. On the other hand, if either the C content or the N content exceeds 0.025%, the austenite phase generated during the hot-rolled sheet annealing excessively increases to more than 20%. As a result, the ferrite grains with many carbonitrides generated by subsequent cold-rolled sheet annealing increase, and the interface area between the ferrite grains with many carbonitrides and the few ferrite grains that become the starting point of voids during processing increases. The overhang moldability cannot be expressed. Therefore, the upper limit of each of the C content and the N content needs to be 0.025%.
　熱延板焼鈍温度に関しては、フェライト相とオーステナイト相の二相域、特に９００～１１００℃の範囲で焼鈍を行うことによって所定量のオーステナイト相を安定的に確保できるとともに、冷延板焼鈍後の粒径を過度に粗大化させることなく、良好な表面品質が得られる。 Regarding the hot-rolled sheet annealing temperature, a predetermined amount of austenite phase can be stably secured by annealing in the two-phase region of the ferrite phase and the austenite phase, particularly in the range of 900 to 1100 ° C. Good surface quality can be obtained without excessively coarsening the particle size.
　さらに、上記Ｃ含有量およびＮ含有量を有する鋼に本発明の技術的特徴の１つであるフェライト相とオーステナイト相の二相域温度で熱延板焼鈍を行うことによって、破断伸び、平均ｒ値および耐リジング性に関しても有益な効果が得られることを見出した。従来の技術では熱延板焼鈍をフェライト単相域温度で行っていたが、本発明ではフェライト相とオーステナイト相の二相域となる高温で熱延板焼鈍を行うため、フェライト相の再結晶と粒成長が一層促進され、結晶粒径が適度に大きくなる。これにより、破断伸びの向上効果、および焼鈍集合組織の発達が一層促進されることによる平均ｒ値の向上効果が得られる。なお、破断伸びは、以下の理由によっても向上する。Ｃ含有量およびＮ含有量を本発明が推奨するレベルにまで低減することによって、冷延板焼鈍後に生成する炭窒化物量が減少し、引張変形時のボイドの発生およびボイドの連結が抑制される。これによっても破断伸びは向上する。 Further, the steel having the C content and the N content is subjected to hot-rolled sheet annealing at a two-phase region temperature of a ferrite phase and an austenite phase, which is one of the technical features of the present invention, to thereby obtain an elongation at break and an average r It has been found that beneficial effects are also obtained with respect to value and ridging resistance. In the prior art, hot-rolled sheet annealing was performed at a ferrite single-phase temperature, but in the present invention, hot-rolled sheet annealing is performed at a high temperature that is a two-phase region of a ferrite phase and an austenite phase. Grain growth is further promoted and the crystal grain size is increased appropriately. Thereby, the improvement effect of breaking elongation and the improvement effect of average r value by the further acceleration | stimulation of the development of an annealing texture are acquired. The elongation at break is also improved for the following reasons. By reducing the C content and N content to the levels recommended by the present invention, the amount of carbonitride produced after cold-rolled sheet annealing is reduced, and the generation of voids and the connection of voids during tensile deformation are suppressed. . This also improves the elongation at break.
　耐リジング性に関して、有益な効果が得られる理由は以下の通りである。熱延板焼鈍でフェライト相からオーステナイト相が生成する際に、オーステナイト相が焼鈍前のフェライト相とは異なった結晶方位を有して生成する。さらに、熱延板焼鈍後の金属組織がマルテンサイト相とフェライト相の二相組織となる。その後の冷間圧延時に、マルテンサイト相に挟まれたフェライト相内に圧延ひずみが局所的に集中し、フェライト相内に方位差が形成される。フェライト相内に方位差が形成されることにより、その後の冷延板焼鈍において方位差が導入された部位にて再結晶が優先的に生じる。その結果、フェライト相のコロニーが効果的に破壊され、リジング高さで２．５μｍ以下の優れた耐リジング特性が得られる。 The reason why a beneficial effect can be obtained with respect to ridging resistance is as follows. When an austenite phase is generated from a ferrite phase by hot-rolled sheet annealing, the austenite phase is generated with a crystal orientation different from that of the ferrite phase before annealing. Furthermore, the metal structure after hot-rolled sheet annealing becomes a two-phase structure of a martensite phase and a ferrite phase. During the subsequent cold rolling, rolling strain is locally concentrated in the ferrite phase sandwiched between the martensite phases, and an orientation difference is formed in the ferrite phase. By forming an orientation difference in the ferrite phase, recrystallization occurs preferentially at a site where the orientation difference is introduced in the subsequent cold-rolled sheet annealing. As a result, the ferrite phase colonies are effectively destroyed, and excellent ridging resistance with a ridging height of 2.5 μm or less is obtained.
　以上より、十分な張出成形性、破断伸び、平均ｒ値、および耐リジング性の全てを具有させるためには、以下の条件が必要となる。まず、鋼成分がオーステナイト相が生成するＣ含有量およびＮ含有量であることが前提となる。その上で、所定量のオーステナイト相が生成できる範囲でＣ含有量、Ｎ含有量を低減させる。このような成分を有する鋼について、熱延板焼鈍をフェライト相とオーステナイトの二相域温度で行った後に、冷間圧延および冷延板焼鈍を行う。これにより、炭窒化物が多いフェライト粒と少ないフェライト粒からなるフェライト単相組織とすることが必要である。 From the above, the following conditions are necessary in order to have sufficient stretch formability, elongation at break, average r value, and ridging resistance. First, it is premised that the steel components have a C content and an N content generated by the austenite phase. In addition, the C content and the N content are reduced within a range in which a predetermined amount of austenite phase can be generated. About steel which has such a component, after performing hot-rolled sheet annealing at the two-phase region temperature of a ferrite phase and austenite, cold rolling and cold-rolled sheet annealing are performed. As a result, it is necessary to obtain a ferrite single-phase structure composed of ferrite grains containing a large amount of carbonitride and few ferrite grains.
　次に、本発明のフェライト系ステンレス鋼板の成分組成について説明する。
以下、特に断らない限り％は質量％を意味する。 Next, the component composition of the ferritic stainless steel sheet of the present invention will be described.
Hereinafter, unless otherwise specified,% means mass%.
　Ｃ: ０.００５～０.０２５％
Ｃはオーステナイト相の生成を促進し、熱延板焼鈍時にフェライト相とオーステナイト相が出現する二相温度域を拡大する効果がある。この効果を得るためには０.００５％以上の含有が必要である。しかし、Ｃ量が０.０２５％を超えると熱延板焼鈍におけるオーステナイト相の生成量が過剰となって、冷延板焼鈍後に炭窒化物の多いフェライト粒の生成量が過剰となる。その結果、金属組織中のボイド間距離が小さくなり、成形時にボイド連結に起因した破断が生じやすくなり、十分な張出成形性が得られなくなる。そのため、Ｃ量は０.００５～０.０２５％の範囲とする。好ましくは０．０１０～０．０２０％の範囲である。 C: 0.005 to 0.025%
C promotes the formation of the austenite phase and has the effect of expanding the two-phase temperature range where the ferrite phase and the austenite phase appear during hot-rolled sheet annealing. In order to obtain this effect, a content of 0.005% or more is necessary. However, if the amount of C exceeds 0.025%, the amount of austenite phase produced in the hot-rolled sheet annealing becomes excessive, and the amount of ferrite grains containing many carbonitrides becomes excessive after cold-rolled sheet annealing. As a result, the distance between voids in the metal structure is reduced, and breakage due to void connection is likely to occur at the time of molding, and sufficient stretch formability cannot be obtained. Therefore, the C content is in the range of 0.005 to 0.025%. Preferably it is 0.010 to 0.020% of range.
　Ｓｉ:０．０２～０．５０％
Ｓｉは鋼溶製時に脱酸剤として作用する元素である。この効果を得るためには０．０２％以上の含有が必要である。しかし、Ｓｉ量が０．５０％を超えると、鋼板が硬質化して熱間圧延時の圧延負荷が増大するとともに、仕上げ焼鈍後の延性が低下する。そのため、Ｓｉ量は０．０２～０．５０％の範囲とする。好ましくは０．１０～０．３５％の範囲である。さらに好ましくは０．１０～０．２０％の範囲である。 Si: 0.02 to 0.50%
Si is an element that acts as a deoxidizer during steel melting. In order to acquire this effect, 0.02% or more needs to be contained. However, if the amount of Si exceeds 0.50%, the steel sheet becomes hard, the rolling load during hot rolling increases, and the ductility after finish annealing decreases. For this reason, the Si content is in the range of 0.02 to 0.50%. Preferably it is 0.10 to 0.35% of range. More preferably, it is in the range of 0.10 to 0.20%.
　Ｍｎ：０．５５～１．００％
ＭｎはＣと同様にオーステナイト相の生成を促進し、熱延板焼鈍時にフェライト相とオーステナイト相が出現する二相温度域を拡大する効果がある。この効果を得るためには０．５５％以上の含有が必要である。しかし、Ｍｎ量が１．００％を超えるとＭｎＳの生成量が増加して耐食性が低下する。そのため、Ｍｎ量は０．５５～１．００％の範囲とする。好ましくは０．６０～０．９０％の範囲である。さらに好ましくは０．７５～０．８５％の範囲である。 Mn: 0.55 to 1.00%
Mn, like C, promotes the formation of an austenite phase and has the effect of expanding the two-phase temperature range in which a ferrite phase and an austenite phase appear during hot-rolled sheet annealing. In order to acquire this effect, 0.55% or more needs to be contained. However, if the amount of Mn exceeds 1.00%, the amount of MnS produced increases and the corrosion resistance decreases. Therefore, the Mn content is set in the range of 0.55 to 1.00%. Preferably it is 0.60 to 0.90% of range. More preferably, it is in the range of 0.75 to 0.85%.
　Ｐ：０．０４％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ため低い方が望ましく、上限を０．０４％とする。好ましくは０．０３％以下である。さらに好ましくは０．０１％以下である。 P: 0.04% or less Since P is an element that promotes grain boundary fracture due to grain boundary segregation, the lower one is desirable, and the upper limit is made 0.04%. Preferably it is 0.03% or less. More preferably, it is 0.01% or less.
　Ｓ:０．０１％以下
ＳはＭｎＳなどの硫化物系介在物となって存在して延性や耐食性等を低下させる元素であり、特に含有量が０．０１％を超えた場合にそれらの悪影響が顕著に生じる。そのためＳ量は極力低い方が望ましく、本発明ではＳ量の上限を０．０１％とする。好ましくは０．００７％以下である。さらに好ましくは０．００５％以下である。 S: 0.01% or less S is an element that exists as sulfide inclusions such as MnS and lowers ductility, corrosion resistance, etc., and particularly when the content exceeds 0.01%, their adverse effects Is noticeable. For this reason, the S amount is desirably as low as possible. In the present invention, the upper limit of the S amount is set to 0.01%. Preferably it is 0.007% or less. More preferably, it is 0.005% or less.
　Ａｌ：０．００１～０．１０％
ＡｌはＳｉと同様に脱酸剤として作用する元素である。この効果を得るためには０．００１％以上の含有が必要である。しかし、Ａｌ量が０．１０％を超えると、Ａｌ２Ｏ３等のＡｌ系介在物が増加し、表面性状が低下しやすくなる。そのため、Ａｌ量は０．００１～０．１０％の範囲とする。好ましくは０．００１～０．０７％の範囲である。さらに好ましくは０．００１～０．０５％の範囲である。 Al: 0.001 to 0.10%
Al is an element that acts as a deoxidizing agent like Si. In order to acquire this effect, 0.001% or more needs to be contained. However, when the Al content exceeds 0.10%, Al-based inclusions such as Al 2 O 3 increase, and the surface properties tend to decrease. Therefore, the Al content is set in the range of 0.001 to 0.10%. Preferably it is 0.001 to 0.07% of range. More preferably, it is in the range of 0.001 to 0.05%.
　Ｃｒ:１５．５～１８．０％
Ｃｒは鋼板表面に不動態皮膜を形成して耐食性を向上させる効果を有する元素である。この効果を得るためにはＣｒ量を１５．５％以上とする必要がある。しかし、Ｃｒ量が１８．０％を超えると、熱延板焼鈍時にオーステナイト相の生成が不十分となり、所望の材料特性が得られない。そのため、Ｃｒ量は１５．５～１８．０％の範囲とする。好ましくは１６．０～１７．０％の範囲である。さらに好ましくは１６．０～１６．５％の範囲である。 Cr: 15.5 to 18.0%
Cr is an element having an effect of improving the corrosion resistance by forming a passive film on the surface of the steel sheet. In order to obtain this effect, the Cr amount needs to be 15.5% or more. However, if the Cr content exceeds 18.0%, the austenite phase is not sufficiently generated during hot-rolled sheet annealing, and desired material characteristics cannot be obtained. Therefore, the Cr content is in the range of 15.5 to 18.0%. Preferably it is 16.0 to 17.0% of range. More preferably, it is in the range of 16.0 to 16.5%.
　Ｎｉ：０．１～１．０％
Ｎｉは耐食性を向上させる元素であり、特に高い耐食性が要求される場合には含有することが有効である。また、Ｎｉにはオーステナイト相の生成を促進し、熱延板焼鈍時にフェライト相とオーステナイト相が出現する二相温度域を拡大する効果がある。これらの効果は０．１％以上の含有で顕著となる。しかし、Ｎｉ含有量が１．０％を超えると成形性が低下するため好ましくない。そのためＮｉを含有する場合は０．１～１．０％とする。好ましくは０．１～０．３％の範囲である。 Ni: 0.1 to 1.0%
Ni is an element that improves corrosion resistance, and it is effective to contain it particularly when high corrosion resistance is required. Ni also has the effect of promoting the formation of the austenite phase and expanding the two-phase temperature range in which the ferrite phase and austenite phase appear during hot-rolled sheet annealing. These effects become significant when the content is 0.1% or more. However, if the Ni content exceeds 1.0%, the formability deteriorates, which is not preferable. Therefore, when Ni is contained, the content is made 0.1 to 1.0%. Preferably it is 0.1 to 0.3% of range.
　Ｎ:０．００５～０．０２５％
Ｎは、Ｃ、Ｍｎと同様にオーステナイト相の生成を促進し、熱延板焼鈍時にフェライト相とオーステナイト相が出現する二相温度域を拡大する効果がある。この効果を得るためにはＮ量を０．００５％以上とする必要がある。しかし、Ｎ量が０．０２５％を超えると延性が著しく低下する上、熱延板焼鈍におけるオーステナイト相の生成量が過剰となって、冷延板焼鈍後に炭窒化物の多いフェライト粒の生成量が過剰となる。その結果、金属組織中のボイド間距離が小さくなり、成形時にボイド連結に起因した破断が生じやすくなり、十分な張出成形性が得られなくなる。そのため、Ｎ量は０．００５～０．０２５％の範囲とする。好ましくは０．０１０～０．０２０％の範囲である。 N: 0.005 to 0.025%
N, like C and Mn, promotes the formation of the austenite phase and has the effect of expanding the two-phase temperature range in which the ferrite phase and austenite phase appear during hot-rolled sheet annealing. In order to obtain this effect, the N amount needs to be 0.005% or more. However, if the N content exceeds 0.025%, the ductility is remarkably lowered, and the amount of austenite phase generated in hot-rolled sheet annealing becomes excessive, and the amount of ferrite grains that are rich in carbonitride after cold-rolled sheet annealing is increased. Becomes excessive. As a result, the distance between voids in the metal structure is reduced, and breakage due to void connection is likely to occur at the time of molding, and sufficient stretch formability cannot be obtained. Therefore, the N content is set in the range of 0.005 to 0.025%. Preferably it is 0.010 to 0.020% of range.
　残部はFeおよび不可避的不純物である。 The balance is Fe and inevitable impurities.
　以上の成分組成により本発明の効果は得られるが、さらに製造性あるいは材料特性を向上させる目的で以下の元素を含有することができる。 Although the effects of the present invention can be obtained by the above component composition, the following elements can be contained for the purpose of further improving manufacturability or material properties.
　Ｃｕ：０．１～１．０％、Ｖ：０．０１～０．１０％、Ｔｉ：０．００１～０．０５％、Ｎｂ：０．００１～０．０５％、Ｍｏ：０．１～０．５％、Ｃｏ：０．０１～０．２％のうちから選ばれる１種または２種以上
　Ｃｕ：０．１～１．０％
Ｃｕは耐食性を向上させる元素であり、特に高い耐食性が要求される場合には含有することが有効である。また、Ｃｕにはオーステナイト相の生成を促進し、熱延板焼鈍時にフェライト相とオーステナイト相が出現する二相温度域を拡大する効果がある。これらの効果は０．１％以上の含有で顕著となる。しかし、Ｃｕ含有量が１．０％を超えると成形性が低下する場合があり好ましくない。そのためＣｕを含有する場合は０．１～１．０％とする。好ましくは０．２～０．３％の範囲である。 Cu: 0.1 to 1.0%, V: 0.01 to 0.10%, Ti: 0.001 to 0.05%, Nb: 0.001 to 0.05%, Mo: 0.1 to One or more selected from 0.5%, Co: 0.01 to 0.2% Cu: 0.1 to 1.0%
Cu is an element that improves corrosion resistance, and it is effective to contain it particularly when high corrosion resistance is required. In addition, Cu has an effect of promoting the generation of an austenite phase and expanding a two-phase temperature range in which a ferrite phase and an austenite phase appear during hot-rolled sheet annealing. These effects become significant when the content is 0.1% or more. However, if the Cu content exceeds 1.0%, formability may be deteriorated, which is not preferable. Therefore, when Cu is contained, the content is made 0.1 to 1.0%. Preferably it is 0.2 to 0.3% of range.
　Ｖ:０．０１～０．１０％
Ｖは鋼中のＣおよびＮと化合して、固溶Ｃおよび固溶Ｎを低減する。これにより、平均ｒ値を向上させる。この効果を得るためにはＶ量を０．０１％以上含有する必要がある。しかし、Ｖ量が０．１０％を超えると加工性が低下するとともに、製造コストの上昇を招く。そのため、Ｖを含有する場合は０．０１～０．１０％の範囲とする。好ましくは０．０２～０．０８％の範囲である。 V: 0.01-0.10%
V combines with C and N in the steel to reduce solute C and solute N. This improves the average r value. In order to acquire this effect, it is necessary to contain V amount 0.01% or more. However, if the amount of V exceeds 0.10%, the workability is lowered and the manufacturing cost is increased. Therefore, when V is contained, the content is made 0.01 to 0.10%. Preferably it is 0.02 to 0.08% of range.
　Ｔｉ：０．００１～０．０５％、Ｎｂ：０．００１～０．０５％
ＴｉおよびＮｂはＶと同様に、ＣおよびＮとの親和力の高い元素であり、熱間圧延時に炭化物あるいは窒化物として析出し、母相中の固溶Ｃおよび固溶Ｎを低減させ、冷延板焼鈍後の加工性を向上させる効果がある。この効果を得るためには、０．００１％以上のＴｉ、０．００１％以上のＮｂを含有する必要がある。しかし、Ｔｉ量が０．０５％を超えると、あるいはＮｂ量が０．０５％を超えると、過剰なＴｉＮおよびＮｂＣの析出により良好な表面性状を得ることができない。そのため、Ｔｉを含有する場合は０．００１～０．０５％の範囲、Ｎｂを含有する場合は０．００１～０．０５％の範囲とする。Ｔｉ量は好ましくは０．００３～０．０３％の範囲である。さらに好ましくは０．００５～０．０１５％の範囲である。Ｎｂ量は好ましくは０．００３～０．０３％の範囲である。さらに好ましくは０．００５～０．０１５％の範囲である。 Ti: 0.001 to 0.05%, Nb: 0.001 to 0.05%
Ti and Nb are elements having a high affinity with C and N, like V, and precipitate as carbide or nitride during hot rolling, reducing the solid solution C and solid solution N in the matrix, and cold rolling. There is an effect of improving workability after sheet annealing. In order to obtain this effect, it is necessary to contain 0.001% or more of Ti and 0.001% or more of Nb. However, when the Ti content exceeds 0.05% or the Nb content exceeds 0.05%, good surface properties cannot be obtained due to excessive precipitation of TiN and NbC. Therefore, when Ti is contained, the range is 0.001 to 0.05%, and when Nb is contained, the range is 0.001 to 0.05%. The amount of Ti is preferably in the range of 0.003 to 0.03%. More preferably, it is in the range of 0.005 to 0.015%. The amount of Nb is preferably in the range of 0.003 to 0.03%. More preferably, it is in the range of 0.005 to 0.015%.
　Ｍｏ：０．１～０．５％
Ｍｏは耐食性を向上させる元素であり、特に高い耐食性が要求される場合には含有することが有効である。この効果は０．１％以上の含有で顕著となる。しかし、Ｍｏ量が０．５％を超えると熱延板焼鈍時にオーステナイト相の生成が不十分となり、所望の材料特性が得られなくなり好ましくない。そのため、Ｍｏを含有する場合は０．１～０．５％とする。好ましくは０．２～０．３％の範囲である。 Mo: 0.1 to 0.5%
Mo is an element that improves corrosion resistance, and it is effective to contain it particularly when high corrosion resistance is required. This effect becomes remarkable when the content is 0.1% or more. However, if the Mo content exceeds 0.5%, the austenite phase is not sufficiently generated during hot-rolled sheet annealing, and desired material characteristics cannot be obtained. Therefore, when it contains Mo, it is 0.1 to 0.5%. Preferably it is 0.2 to 0.3% of range.
　Ｃｏ：０．０１～０．２％
Ｃｏは靭性を向上させる元素である。この効果は０．０１％以上の含有によって得られる。一方、含有量が０．２％を超えると成形性を低下させる．そのため、Ｃｏを含有する場合の含有量は０．０１～０．２％の範囲とする。 Co: 0.01 to 0.2%
Co is an element that improves toughness. This effect is obtained when the content is 0.01% or more. On the other hand, if the content exceeds 0.2%, the moldability is lowered. Therefore, if Co is contained, the content is made 0.01 to 0.2%.
　Ｍｇ：０．０００２～０．００５０％、Ｃａ：０．０００２～０．００２０％、Ｂ：０．０００２～０．００５０％、ＲＥＭ：０．０１～０．１０％のうちから選ばれる１種または２種以上
　Ｍｇ:０．０００２～０．００５０％
Ｍｇは熱間加工性を向上させる効果がある元素である。この効果を得るためには０．０００２％以上の含有が必要である。しかし、Ｍｇ量が０．００５０％を超えると表面品質が低下する。そのため、Ｍｇを含有する場合は０．０００２～０．００５０％の範囲とする。好ましくは０．０００５～０．００３５％の範囲である。さらに好ましくは０．０００５～０．００２０％の範囲である。 One selected from Mg: 0.0002 to 0.0050%, Ca: 0.0002 to 0.0020%, B: 0.0002 to 0.0050%, REM: 0.01 to 0.10% Or two or more Mg: 0.0002 to 0.0050%
Mg is an element that has an effect of improving hot workability. In order to acquire this effect, 0.0002% or more needs to be contained. However, when the amount of Mg exceeds 0.0050%, the surface quality deteriorates. Therefore, when Mg is contained, the content is made 0.0002 to 0.0050%. Preferably it is 0.0005 to 0.0035% of range. More preferably, it is in the range of 0.0005 to 0.0020%.
　Ｃａ：０．０００２～０．００２０％
Ｃａは連続鋳造の際に発生しやすい介在物の晶出によるノズルの閉塞を防止するのに有効な成分である。その効果を得るためには０．０００２％以上の含有が必要である。しかし、Ｃａ量が０．００２０％を超えるとＣａＳが生成して耐食性が低下する。そのため、Ｃａを含有する場合は０．０００２～０．００２０％の範囲とする。好ましくは０．０００５～０．００１５％の範囲である。さらに好ましくは０．０００５～０．００１０％の範囲である。 Ca: 0.0002 to 0.0020%
Ca is an effective component for preventing nozzle clogging due to crystallization of inclusions that are likely to occur during continuous casting. In order to acquire the effect, 0.0002% or more needs to be contained. However, if the Ca content exceeds 0.0020%, CaS is generated and the corrosion resistance is lowered. Therefore, when Ca is contained, the content is made 0.0002 to 0.0020%. Preferably it is 0.0005 to 0.0015% of range. More preferably, it is in the range of 0.0005 to 0.0010%.
　Ｂ：０．０００２～０．００５０％
Ｂは低温二次加工脆化を防止するのに有効な元素である。この効果を得るためには０．０００２％以上の含有が必要である。しかし、Ｂ量が０．００５０％を超えると熱間加工性が低下する。そのため、Ｂを含有する場合は０．０００２～０．００５０％の範囲とする。好ましくは０．０００５～０．００３５％の範囲である。さらに好ましくは０．０００５～０．００２０％の範囲である。 B: 0.0002 to 0.0050%
B is an element effective for preventing embrittlement at low temperature secondary work. In order to acquire this effect, 0.0002% or more needs to be contained. However, when the amount of B exceeds 0.0050%, hot workability deteriorates. Therefore, when B is contained, the content is made 0.0002 to 0.0050%. Preferably it is 0.0005 to 0.0035% of range. More preferably, it is in the range of 0.0005 to 0.0020%.
　ＲＥＭ:０．０１～０．１０％
ＲＥＭ（Ｒａｒｅ Ｅａｒｔｈ Ｍｅｔａｌｓ）は耐酸化性を向上させる元素であり、特に溶接部の酸化皮膜の形成を抑制し溶接部の耐食性を向上させる効果がある。この効果を得るためには０．０１％以上の含有が必要である。しかし、０．１０％を超えて含有すると冷延焼鈍時の酸洗性などの製造性を低下させる。また、ＲＥＭは高価な元素であるため、過度な含有は製造コストの増加を招くため好ましくない。そのため、ＲＥＭを含有する場合は０．０１～０．１０％の範囲とする。好ましくは０．０１～０．０５％の範囲である。 REM: 0.01-0.10%
REM (Rare Earth Metals) is an element that improves oxidation resistance, and is particularly effective in suppressing the formation of an oxide film on the welded portion and improving the corrosion resistance of the welded portion. In order to obtain this effect, a content of 0.01% or more is necessary. However, if the content exceeds 0.10%, productivity such as pickling at the time of cold rolling annealing is lowered. Moreover, since REM is an expensive element, excessive inclusion is not preferable because it causes an increase in manufacturing cost. Therefore, when REM is contained, the content is made 0.01 to 0.10%. Preferably it is 0.01 to 0.05% of range.
　次に本発明のフェライト系ステンレス鋼板の製造方法について説明する。
本発明のフェライト系ステンレス鋼板は上記成分組成を有する鋼スラブに対して、熱間圧延を施した後、９００～１１００℃の温度範囲で５秒～１５分間保持する熱延板焼鈍を行い、次いで冷間圧延を施した後、８００～９００℃の温度範囲で５秒～５分間保持する冷延板焼鈍を行うことで得られる。 Next, the manufacturing method of the ferritic stainless steel plate of this invention is demonstrated.
The ferritic stainless steel sheet of the present invention is subjected to hot rolling on a steel slab having the above component composition, followed by hot rolling sheet annealing at a temperature range of 900 to 1100 ° C. for 5 seconds to 15 minutes, After cold rolling, it is obtained by annealing a cold-rolled sheet that is held at a temperature range of 800 to 900 ° C. for 5 seconds to 5 minutes.
　まずは、上記した成分組成からなる溶鋼を、転炉、電気炉、真空溶解炉等の公知の方法で溶製し、連続鋳造法あるいは造塊－分塊法により鋼素材（スラブ）とする。このスラブを、１１００～１２５０℃で１～２４時間加熱するか、あるいは加熱することなく鋳造まま直接、熱間圧延して熱延板とする。 First, the molten steel having the above component composition is melted by a known method such as a converter, electric furnace, vacuum melting furnace or the like, and is made into a steel material (slab) by a continuous casting method or an ingot-bundling method. The slab is heated at 1100 to 1250 ° C. for 1 to 24 hours, or directly hot-rolled as cast without heating to form a hot-rolled sheet.
　次いで、熱間圧延を行う。巻取りでは、巻取り温度を５００℃以上８５０℃以下とすることが好ましい。５００℃未満では巻取り後の熱延板組織中にマルテンサイト相が生成し、その後の熱延板焼鈍における再結晶および粒成長が遅滞する。これにより、熱延板焼鈍組織中の微細粒が増加し、当該微細粒が冷延板焼鈍組織注にも残存するために冷延板焼鈍後の延性が低下する場合があるため好ましくない。８５０℃超で巻き取ると粒径が大きくなり、プレス加工時に肌荒れが発生してしまう場合がある。したがって、巻取り温度は５００～８５０℃の範囲が好ましい。 Next, hot rolling is performed. In winding, the winding temperature is preferably 500 ° C. or higher and 850 ° C. or lower. If it is less than 500 degreeC, a martensite phase will produce | generate in the hot-rolled sheet | seat structure after winding, and the recrystallization and grain growth in subsequent hot-rolled sheet | seat annealing will be overdue. Thereby, since the fine grain in a hot-rolled sheet annealing structure | tissue increases and the said fine grain remains also in a cold-rolled sheet annealing structure | tissue injection | pouring, since the ductility after cold-rolled sheet annealing may fall, it is unpreferable. When it winds up above 850 degreeC, a particle size will become large and rough skin may generate | occur | produce at the time of press work. Therefore, the winding temperature is preferably in the range of 500 to 850 ° C.
　９００～１１００℃の温度範囲で５秒～１５分間保持する熱延板焼鈍
その後、フェライト相とオーステナイト相の二相温度域となる９００～１１００℃の温度範囲で５秒～１５分間保持する熱延板焼鈍を行う。
熱延板焼鈍は本発明が優れた成形性および耐リジング特性を得るために極めて重要な工程である。熱延板焼鈍温度が９００℃未満では十分な再結晶が生じないうえ、フェライト単相域となるため、二相温度域での焼鈍によって発現する本発明の効果が得られない場合がある。一方、焼鈍温度が１１００℃を超えるとオーステナイト相の生成量が著しく低下し、所定の耐リジング性が得られない場合がある。焼鈍時間が５秒未満の場合、所定の温度で焼鈍したとしてもオーステナイト相の生成とフェライト相の再結晶が十分に生じないため、所望の成形性が得られない場合がある。一方、焼鈍時間が１５分を超えるとオーステナイト相中へのＣ濃化が助長されてマルテンサイト相が過度に硬質化する。その結果、その後の冷間圧延において鋼板表面に過度に硬質なマルテンサイトに起因した表面疵が発生し、冷延板焼鈍後の表面性状が悪化する場合がある。そのため、熱延板焼鈍は９００～１１００℃の温度範囲で、５秒～１５分間保持する。好ましくは、９２０～１０８０℃の温度範囲で１５秒～５分間保持である。さらに好ましくは９４０～１０４０℃の温度範囲で３０秒～３分間保持である。 Hot-rolled sheet annealing held for 5 seconds to 15 minutes in the temperature range of 900 to 1100 ° C. Thereafter, hot rolling for 5 seconds to 15 minutes in the temperature range of 900 to 1100 ° C. which is a two-phase temperature range of the ferrite phase and austenite phase Sheet annealing is performed.
Hot-rolled sheet annealing is an extremely important process for the present invention to obtain excellent formability and ridging resistance. When the hot-rolled sheet annealing temperature is less than 900 ° C., sufficient recrystallization does not occur and the ferrite single-phase region is obtained, so that the effects of the present invention that are manifested by annealing in the two-phase temperature region may not be obtained. On the other hand, if the annealing temperature exceeds 1100 ° C., the amount of austenite phase produced is significantly reduced, and the predetermined ridging resistance may not be obtained. When the annealing time is less than 5 seconds, even if annealing is performed at a predetermined temperature, generation of austenite phase and recrystallization of the ferrite phase do not occur sufficiently, so that desired formability may not be obtained. On the other hand, if the annealing time exceeds 15 minutes, C concentration in the austenite phase is promoted and the martensite phase becomes excessively hard. As a result, surface defects resulting from excessively hard martensite are generated on the surface of the steel sheet in the subsequent cold rolling, and the surface properties after the cold-rolled sheet annealing may be deteriorated. Therefore, the hot-rolled sheet annealing is held in the temperature range of 900 to 1100 ° C. for 5 seconds to 15 minutes. Preferably, the temperature is maintained at 920 to 1080 ° C. for 15 seconds to 5 minutes. More preferably, the temperature is kept at 940 to 1040 ° C. for 30 seconds to 3 minutes.
　次いで、必要に応じて酸洗を施し、冷間圧延を行う。冷間圧延は伸び性、曲げ性、プレス成形性および形状矯正の観点から、５０％以上の圧下率で行うことが好ましい。また、本発明では、冷延－焼鈍を２回以上繰り返しても良い。さらに表面性状を向上させるために、研削や研磨等を施してもよい。 Next, pickling is performed as necessary, and cold rolling is performed. Cold rolling is preferably performed at a rolling reduction of 50% or more from the viewpoints of extensibility, bendability, press formability, and shape correction. In the present invention, cold rolling and annealing may be repeated twice or more. Further, in order to improve the surface properties, grinding or polishing may be performed.
　８００～９００℃の温度範囲で５秒～５分間保持する冷延板焼鈍
次いで、冷延板焼鈍を行う。冷延板焼鈍は熱延板焼鈍で形成したフェライト相とマルテンサイト相の二相組織をフェライト単相組織とするために重要な工程である。冷延板焼鈍温度が８００℃未満では再結晶が十分に生じず所定の成形性を得ることができない。一方、冷延板焼鈍温度が９００℃を超えた場合、９００℃を超える温度がフェライト相とオーステナイト相の二相温度域となる鋼成分では冷延板焼鈍後にマルテンサイト相が生成するために鋼板が硬質化し、所定の破断伸びおよび張出成形性を得ることができない。また、９００℃を超える温度がフェライト単相温度域となる鋼成分であったとしても、結晶粒の著しい粗大化により、鋼板の光沢度が低下するため表面品質の観点で好ましくない。焼鈍時間が５秒未満の場合、所定の温度で焼鈍したとしてもフェライト相の再結晶が十分に生じないため、所定の成形性を得ることができない。焼鈍時間が５分を超えると、結晶粒が著しく粗大化し、鋼板の光沢度が低下するため表面品質の観点で好ましくない。そのため、冷延板焼鈍は８００～９００℃の温度範囲で５秒～５分間保持とする。好ましくは、８５０℃～９００℃の温度範囲で１５秒～３分間保持である。より光沢を求めるためにＢＡ焼鈍（光輝焼鈍）を行っても良い。 Cold-rolled sheet annealing is carried out for 5 seconds to 5 minutes in a temperature range of 800 to 900 ° C. Next, cold-rolled sheet annealing is performed. Cold-rolled sheet annealing is an important process for making a two-phase structure of a ferrite phase and a martensite phase formed by hot-rolled sheet annealing into a ferrite single-phase structure. If the cold-rolled sheet annealing temperature is less than 800 ° C., sufficient recrystallization does not occur and a predetermined formability cannot be obtained. On the other hand, when the cold-rolled sheet annealing temperature exceeds 900 ° C., the steel component in which the temperature exceeding 900 ° C. becomes the two-phase temperature range of the ferrite phase and the austenite phase generates a martensite phase after the cold-rolled sheet annealing. Becomes hard, and the predetermined elongation at break and stretchability cannot be obtained. Moreover, even if it is a steel component in which the temperature exceeding 900 ° C. is the ferrite single phase temperature range, the glossiness of the steel sheet is lowered due to the remarkable coarsening of crystal grains, which is not preferable from the viewpoint of surface quality. When the annealing time is less than 5 seconds, even if annealing is performed at a predetermined temperature, the ferrite phase is not sufficiently recrystallized, and therefore, a predetermined formability cannot be obtained. When the annealing time exceeds 5 minutes, the crystal grains become extremely coarse and the glossiness of the steel sheet is lowered, which is not preferable from the viewpoint of surface quality. Therefore, cold-rolled sheet annealing is held for 5 seconds to 5 minutes in a temperature range of 800 to 900 ° C. Preferably, the temperature is maintained at 850 ° C. to 900 ° C. for 15 seconds to 3 minutes. In order to obtain more gloss, BA annealing (bright annealing) may be performed.
　さらに、必要に応じて酸洗を施して製品とする。 Furthermore, pickle the product if necessary to make it a product.
　以下、本発明を実施例により詳細に説明する。
表１に示す化学組成を有するステンレス鋼を５０kg小型真空溶解炉にて溶製した。これらの鋼塊を１１５０℃で１ｈｒ加熱後、熱間圧延を施して板厚３.５ｍｍの熱延板とした。次いで、これらの熱延板に表２に記載の条件で熱延板焼鈍を施した後、表面にショットブラスト処理と酸洗による脱スケールを行った。さらに、板厚０.８ｍｍまで冷間圧延した後、表２に記載の条件で冷延板焼鈍を行った。さらに、酸洗による脱スケール処理を行い、冷延酸洗焼鈍板（フェライト系ステンレス鋼板）を得た。 Hereinafter, the present invention will be described in detail with reference to examples.
Stainless steel having the chemical composition shown in Table 1 was melted in a 50 kg small vacuum melting furnace. These steel ingots were heated at 1150 ° C. for 1 hour and then hot rolled to form hot rolled sheets having a thickness of 3.5 mm. Subsequently, these hot-rolled sheets were subjected to hot-rolled sheet annealing under the conditions shown in Table 2, and then the surfaces were descaled by shot blasting and pickling. Further, after cold rolling to a plate thickness of 0.8 mm, cold rolled sheet annealing was performed under the conditions shown in Table 2. Furthermore, descaling treatment by pickling was performed to obtain a cold-rolled pickling annealed plate (ferritic stainless steel plate).
　かくして得られた冷延酸洗焼鈍板（フェライト系ステンレス鋼板）について以下の評価を行った。 The cold roll pickling annealed plate (ferritic stainless steel plate) thus obtained was evaluated as follows.
　（１）張出成形性の評価
冷延酸洗焼鈍板の表面に、評点間距離が１ｍｍとなるように、直径５ｍｍのスクライブドサークルをマーキングしたものを試験片とし、圧延平行方向、圧延４５°方向および圧延直行方向をそれぞれ最大対数ひずみ方向として、中島法によりＦＬＤ（成形限界線図）を作成した。得られたＦＬＤから成形限界の最大対数ひずみの最小値を求め、最大対数ひずみの最小値が０．１５以上の場合を合格（○）、０．１８以上の場合を特に優れる合格（◎）、０．１５未満の場合を不合格（×）とした。 (1) Evaluation of bulging formability The surface of a cold-rolled pickled and annealed plate is marked with a scribed circle having a diameter of 5 mm so that the distance between the scores is 1 mm. An FLD (formation limit diagram) was prepared by the Nakajima method with the maximum direction and the direction orthogonal to the rolling direction as the maximum logarithmic strain direction. The minimum value of the maximum logarithmic strain at the molding limit is obtained from the obtained FLD, and the case where the minimum value of the maximum logarithmic strain is 0.15 or more is acceptable (◯), and the case where it is 0.18 or more is particularly excellent (◎). The case where it was less than 0.15 was determined to be rejected (x).
　（２）延性の評価
冷延酸洗焼鈍板（フェライト系ステンレス鋼板）から、圧延方向と直角にＪＩＳ　１３Ｂ号引張試験片を採取し、引張試験をＪＩＳ　Ｚ２２４１に準拠して行い、破断伸びを測定し、破断伸びが２８％以上の場合を合格（○）、３０％以上の場合を特に優れる合格（◎）、２８％未満の場合を不合格（×）とした。 (2) Evaluation of ductility From a cold-rolled pickled and annealed sheet (ferritic stainless steel sheet), a JIS No. 13B tensile specimen was taken at right angles to the rolling direction, and the tensile test was performed in accordance with JIS Z2241, and the elongation at break was measured. When the elongation at break was 28% or more, it was judged as acceptable (◯), when it was 30% or more, particularly excellent (A), and when it was less than 28%, it was judged as unacceptable (X).
　（３）平均ｒ値の評価
冷延酸洗焼鈍板（フェライト系ステンレス鋼板）から、圧延方向に対して平行（Ｌ方向）、４５°（Ｄ方向）およびに直角（Ｃ方向）となる方向にＪＩＳ　１３Ｂ号引張試験片を採取し、ＪＩＳ　Ｚ２４１１に準拠した引張試験をひずみ１５％まで行って中断し、各方向のｒ値を測定し平均ｒ値（＝（ｒＬ＋２ｒＤ＋ｒＣ）／４）を算出した。ここで、ｒＬ、ｒＤ、ｒＣはそれぞれＬ方向、Ｄ方向およびＣ方向のｒ値である。平均ｒ値は０．７５以上を合格（○）、０．７５未満を不合格（×）とした。 (3) Evaluation of average r value From cold-rolled pickled and annealed sheet (ferritic stainless steel sheet), parallel to rolling direction (L direction), 45 ° (D direction) and perpendicular to C direction (C direction) A JIS No. 13B tensile test piece was collected, the tensile test according to JIS Z2411 was interrupted to a strain of 15%, the r value in each direction was measured, and the average r value (= (r L + 2r D + r C ) / 4. ) Was calculated. Here, r L , r D , and r C are r values in the L direction, the D direction, and the C direction, respectively. As for the average r value, 0.75 or more was regarded as acceptable (◯), and less than 0.75 was regarded as unacceptable (x).
　（４）耐リジング特性の評価
冷延酸洗焼鈍板（フェライト系ステンレス鋼板）から、圧延方向に平行にＪＩＳ　５号引張試験片を採取し、その表面を＃６００のエメリーペーパーを用いて研磨した後、２０％の引張ひずみを付与し、その試験片の平行部中央の研磨面で圧延方向に直角の方向に、表面粗度計を用いて、ＪＩＳ　Ｂ　０６０１（２００１年）で規定される算術平均うねり（Ｗａ）を、測定長１６ｍｍ、ハイカットフィルター波長０．８ｍｍ、ローカットフィルター波長８ｍｍで測定した。算術平均うねり（Ｗａ）が２．５μｍ以下の場合を合格（○）、２．５μｍ超の場合を不合格（×）とした。 (4) Evaluation of Ridging Resistance Characteristics A JIS No. 5 tensile specimen was taken in parallel with the rolling direction from a cold rolled pickled and annealed sheet (ferritic stainless steel sheet), and the surface was polished using # 600 emery paper. Thereafter, 20% tensile strain was applied, and the arithmetic defined in JIS B 0601 (2001) was conducted using a surface roughness meter in a direction perpendicular to the rolling direction on the polished surface at the center of the parallel part of the test piece. Average waviness (Wa) was measured at a measurement length of 16 mm, a high cut filter wavelength of 0.8 mm, and a low cut filter wavelength of 8 mm. The case where the arithmetic average waviness (Wa) was 2.5 μm or less was determined to be acceptable (◯), and the case where it was greater than 2.5 μm was determined to be unacceptable (x).
　（５）耐食性の評価
冷延酸洗焼鈍板から、６０×１００ｍｍの試験片を採取し、表面を＃６００エメリーペーパーにより研磨仕上げした後に端面部をシールした試験片を作製し、ＪＩＳ　Ｈ　８５０２に規定された塩水噴霧サイクル試験に供した。塩水噴霧サイクル試験は、塩水噴霧（５質量％ＮａＣｌ、３５℃、噴霧２ｈr）→乾燥（６０℃、４ｈr、相対湿度４０％）→湿潤（５０℃、２ｈr、相対湿度≧９５％）を１サイクルとして、８サイクル行った。
塩水噴霧サイクル試験を８サイクル実施後の試験片表面を写真撮影し、画像解析により試験片表面の発錆面積を測定し、試験片全面積との比率から発錆率（（試験片中の発錆面積／試験片全面積）×１００　［％］）を算出した。発錆率が１０％以下を特に優れた耐食性で合格（◎）、１０％超２５％以下を合格（○）、２５％超を不合格（×）とした。 (5) Evaluation of corrosion resistance A 60 × 100 mm test piece was collected from a cold rolled pickled and annealed plate, and a test piece was prepared by polishing the surface with # 600 emery paper and then sealing the end face part. Subjected to the prescribed salt spray cycle test. In the salt spray cycle test, salt spray (5 mass% NaCl, 35 ° C., spray 2 hr) → dry (60 ° C., 4 hr, relative humidity 40%) → wet (50 ° C., 2 hr, relative humidity ≧ 95%) is one cycle. As a result, 8 cycles were performed.
Photograph the surface of the specimen after 8 cycles of salt spray cycle test, measure the rusting area on the specimen surface by image analysis, and calculate the rusting rate (( Rust area / total area of test piece) × 100 [%]) was calculated. A rusting rate of 10% or less was determined to pass with excellent corrosion resistance (）), more than 10% to 25% or less passed (◯), and more than 25% to reject (x).
　評価結果を熱延板焼鈍条件および冷延板焼鈍条件と併せて表２に示す。 The evaluation results are shown in Table 2 together with hot-rolled sheet annealing conditions and cold-rolled sheet annealing conditions.
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000002
　鋼成分が本発明の範囲を満たすＮｏ．１～３２（鋼Ｓ１～Ｓ２４）では、破断伸び２８％以上、平均ｒ値が０．７５以上、リジング高さが２．５μｍ以下、耐食性に関しては塩水噴霧サイクル試験を８サイクル実施後の試験片表面の発錆率がいずれも２５％以下、かつ張出し成形性の評価としてＦＬＤに基づく成形限界の最大対数ひずみの最小値が０．１５以上と優れた成形性と耐食性および耐リジング特性が確認された。 No. in which steel components satisfy the scope of the present invention. For 1 to 32 (steel S1 to S24), specimens after breaking elongation of 28% or more, average r value of 0.75 or more, ridging height of 2.5 μm or less, and salt spray cycle test for 8 cycles for corrosion resistance The surface rusting rate is 25% or less, and the minimum value of the maximum logarithmic strain of the molding limit based on FLD is 0.15 or more as an evaluation of the stretch moldability. Excellent moldability, corrosion resistance and ridging resistance are confirmed. It was.
　特に、Ｃｒを１７．８０％含有したＮｏ．１０（鋼Ｎｏ．Ｓ１０）、Ｎｉを０．４％含有したＮｏ.１７(鋼Ｎｏ．Ｓ１７)、Ｃｕを０．４％含有したＮｏ．１８(鋼Ｎｏ．Ｓ１８)およびＭｏを０．３％含有したＮｏ．１９(鋼Ｎｏ．Ｓ１９)では、塩水噴霧サイクル試験後の発錆率が１０％以下（◎）となっており、耐食性が一層向上した。 Especially, No. containing 17.80% Cr. No. 10 (steel No. S10), No. 17 containing 0.4% Ni (steel No. S17), No. 17 containing Cu 0.4%. 18 (steel No. S18) and No. containing 0.3% Mo. In 19 (steel No. S19), the rusting rate after the salt spray cycle test was 10% or less (以下), and the corrosion resistance was further improved.
　一方、Ｃｒ含有量が本発明の範囲を下回るＮｏ．３８（鋼Ｎｏ．Ｓ３０）では、所定の成形性および耐リジング特性は得られたものの、Ｃｒ含有量が不足したために所定の耐食性が得られなかった。 On the other hand, the Cr content falls below the scope of the present invention. With 38 (steel No. S30), although predetermined formability and ridging characteristics were obtained, the predetermined corrosion resistance was not obtained because the Cr content was insufficient.
　Ｃｒ含有量が本発明の範囲を上回るＮｏ．３９（鋼Ｎｏ．Ｓ３１）では、十分な耐食性は得られたが、過剰にＣｒを含有したために熱延板焼鈍時にオーステナイト相が生成せず、所定の耐リジング特性を得ることができなかった。さらに、熱延板焼鈍を二相温度域で行うことによって得られる、粒内炭窒化物の多いフェライト粒と少ないフェライト粒からなる冷延板焼鈍組織を得ることができず、所定の張出成形性が得られなかった。 No. Cr content exceeding the range of the present invention. In 39 (steel No. S31), sufficient corrosion resistance was obtained, but since an excessive amount of Cr was contained, an austenite phase was not formed during hot-rolled sheet annealing, and a predetermined ridging resistance characteristic could not be obtained. Furthermore, it is not possible to obtain a cold-rolled sheet annealed structure consisting of ferrite grains with many intragranular carbonitrides and few ferrite grains, which is obtained by performing hot-rolled sheet annealing in a two-phase temperature range. Sex was not obtained.
　Ｃ含有量が本発明の範囲を下回るＮｏ．３３（鋼Ｎｏ．Ｓ２５）では、所定の破断伸びおよび平均ｒ値は得られたが、オーステナイト生成能が不足したために熱延板焼鈍においてオーステナイト相が生成せず、所定の耐リジング特性および張出成形性を得ることができなかった。これに対して、Ｃ含有量が本発明の範囲を上回るＮｏ．３４（鋼Ｎｏ．Ｓ２６）では、所定の耐リジング特性や張出成形性は得られたが、鋼板が硬質化したために伸びが低下し、所定の破断伸びが得られなかった。 No. C content is below the scope of the present invention. In No. 33 (steel No. S25), a predetermined elongation at break and an average r value were obtained, but since austenite generation ability was insufficient, an austenite phase was not formed in hot-rolled sheet annealing, and predetermined ridging resistance and overhanging properties were obtained. Formability could not be obtained. On the other hand, the C content exceeds the range of the present invention. In 34 (steel No. S26), predetermined ridging resistance and stretch formability were obtained, but the steel sheet was hardened, so that the elongation was lowered and the predetermined breaking elongation was not obtained.
　Ｓｉ含有量が本発明の範囲を上回るＮｏ．２７(鋼Ｎｏ．Ｓ２７)は、過度のＳｉ含有によって鋼板が硬質化し、所定の破断伸びを得ることができなかった。 No. with Si content exceeding the range of the present invention. No. 27 (steel No. S27) was hardened by excessive Si content, and could not obtain a predetermined elongation at break.
　Ｍｎ含有量が本発明の範囲を下回るＮｏ．３６（鋼Ｎｏ．Ｓ２８）では、所定の破断伸びおよび平均ｒ値は得られたが、オーステナイト生成能が不足したために熱延板焼鈍においてオーステナイト相が生成せず、所定の耐リジング特性および張出成形性を得ることができなかった。これに対して、Ｍｎ含有量が本発明の範囲を上回るＮｏ．３７（鋼Ｎｏ．Ｓ２９）では、組織中に多量のＭｎＳが生成したために所定の耐食性が得られなかった。 No. with Mn content below the range of the present invention. In 36 (steel No. S28), a predetermined elongation at break and an average r value were obtained, but since austenite generation ability was insufficient, an austenite phase was not formed in hot-rolled sheet annealing, and a predetermined ridging resistance and overhanging property were obtained. Formability could not be obtained. On the other hand, the Mn content exceeds the range of the present invention. In No. 37 (steel No. S29), a large amount of MnS was generated in the structure, so that the predetermined corrosion resistance was not obtained.
　Ｎ含有量が本発明の範囲を下回るＮｏ．４０（鋼Ｎｏ．Ｓ３２）では、所定の破断伸びおよび平均ｒ値は得られたが、オーステナイト生成能が不足したために熱延板焼鈍においてオーステナイト相が生成せず、所定の耐リジング特性および張出成形性を得ることができなかった。これに対して、Ｎ含有量が本発明の範囲を上回るＮｏ．４１（鋼Ｎｏ．Ｓ３３）では、所定の耐リジング特性や張出成形性は得られたが、鋼板が硬質化したために所定の破断伸びが得られなかった。さらに、組織中に多量のＣｒ窒化物が析出することに起因した鋭敏化が生じ、所定の耐食性が得られなかった。 No. N content is below the range of the present invention. In 40 (steel No. S32), a predetermined elongation at break and an average r value were obtained, but since austenite forming ability was insufficient, an austenite phase was not formed in hot-rolled sheet annealing, and a predetermined ridging resistance and overhanging property were obtained. Formability could not be obtained. On the other hand, the N content exceeds the range of the present invention. In No. 41 (steel No. S33), predetermined ridging resistance characteristics and stretch formability were obtained, but the predetermined breaking elongation was not obtained because the steel plate was hardened. Furthermore, sensitization caused by the precipitation of a large amount of Cr nitride in the structure occurred, and the predetermined corrosion resistance could not be obtained.
　Ｎｏ．４２～４７では、所定の成形性および耐リジング特性は得られたもののＣｒ含有量が不足したために所定の耐食性が得られなかった鋼Ｓ３０を用い、熱延板焼鈍および冷延板焼鈍の条件が成形性および耐リジング特性に対する影響を検討した。熱延板焼鈍温度が本発明を下回るＮｏ．４２では熱延板焼鈍温度がフェライト単相域となったためにオーステナイト相が生成せず、所定の耐リジング性および張出成形性が得られなかったことに加え、十分な再結晶が生じなかったために所定の破断伸びおよび平均ｒ値も得られなかった。熱延板焼鈍温度が本発明の範囲を上回るＮｏ．４３では、オーステナイト相の生成量が低下したために所定の耐リジング特性が得られなかった。熱延板焼鈍の時間が本発明の範囲を下回るＮｏ．４４では、オーステナイト相が十分に生成しなかったことに加え、再結晶が不十分であったため、所定の破断伸び、平均ｒ値および張出成形性が得られなかった。冷延板焼鈍温度、あるいは冷延板焼鈍時間が本発明の範囲を下回るＮｏ．４５およびＮｏ．４７では、熱延板焼鈍で生成したマルテンサイト相が残存するとともに、十分な再結晶が生じなかったために、所定の破断伸びおよび張出成形性が得られなかった。冷延板焼鈍温度が本発明の範囲を上回るＮｏ．４６では、冷延板焼鈍温度がフェライト相とオーステナイト相の二相域となってマルテンサイト相が生成したために鋼板が硬質化し、所定の破断伸びおよび張出成形性が得られなかった。 No. In Nos. 42 to 47, steel S30 was used, which had predetermined formability and ridging resistance but did not have predetermined corrosion resistance due to insufficient Cr content, and the conditions for hot-rolled sheet annealing and cold-rolled sheet annealing were as follows. The effects on moldability and ridging resistance were investigated. No. of hot-rolled sheet annealing temperature lower than the present invention. In No. 42, since the hot-rolled sheet annealing temperature was in the ferrite single-phase region, the austenite phase was not generated, and the predetermined ridging resistance and stretch formability were not obtained, and sufficient recrystallization did not occur. Also, the predetermined breaking elongation and average r value were not obtained. No. of hot-rolled sheet annealing temperature exceeding the range of the present invention. In No. 43, since the amount of austenite phase produced decreased, predetermined ridging resistance characteristics could not be obtained. No. in which the time of hot-rolled sheet annealing is below the range of the present invention. In No. 44, the austenite phase was not sufficiently formed, and recrystallization was insufficient, so that the predetermined elongation at break, average r value and stretch formability could not be obtained. Cold-rolled sheet annealing temperature or cold-rolled sheet annealing time falls below the scope of the present invention. 45 and no. In No. 47, the martensite phase produced by hot-rolled sheet annealing remained and sufficient recrystallization did not occur, so that the predetermined elongation at break and stretchability were not obtained. No. of cold-rolled sheet annealing temperature exceeding the range of the present invention. In No. 46, the cold-rolled sheet annealing temperature became a two-phase region of a ferrite phase and an austenite phase, and a martensite phase was generated. Therefore, the steel plate was hardened, and the predetermined elongation at break and stretchability were not obtained.
　本発明で得られるフェライト系ステンレス鋼板は、張出成形を主体としたプレス成形品を要求される用途、例えば厨房器具や食器への適用に特に好適である。 The ferritic stainless steel sheet obtained by the present invention is particularly suitable for applications requiring press-formed products mainly composed of overhang forming, such as kitchen appliances and tableware.

Claims (4)

  1. 　質量％で、Ｃ：０．００５～０．０２５％、Ｓｉ：０．０２～０．５０％、Ｍｎ：０．５５～１．００％、Ｐ：０．０４％以下、Ｓ：０．０１％以下、Ａｌ：０．００１～０．１０％、Ｃｒ：１５．５～１８．０％、Ｎｉ：０．１～１．０％、Ｎ：０．００５～０．０２５％を含有し、残部がＦｅおよび不可避的不純物からなり、
    破断伸びが２８％以上、平均r値が０．７５以上、かつ、ＦＬＤ（成形限界線図）に基づく成形限界の最大対数ひずみの最小値が０．１５以上であるフェライト系ステンレス鋼板。     By mass%, C: 0.005 to 0.025%, Si: 0.02 to 0.50%, Mn: 0.55 to 1.00%, P: 0.04% or less, S: 0.01 %: Al: 0.001 to 0.10%, Cr: 15.5 to 18.0%, Ni: 0.1 to 1.0%, N: 0.005 to 0.025%, The balance consists of Fe and inevitable impurities,
    A ferritic stainless steel sheet having an elongation at break of 28% or more, an average r value of 0.75 or more, and a minimum value of maximum logarithmic strain at the forming limit based on FLD (forming limit diagram) is 0.15 or more.
  2. 　質量％で、さらに、Ｃｕ：０．１～１．０％、Ｖ：０．０１～０．１０％、Ｔｉ：０．００１～０．０５％、Ｎｂ：０．００１～０．０５％、Ｍｏ：０．１～０．５％、Ｃｏ：０．０１～０．２％のうちから選ばれる１種または２種以上を含む請求項１に記載のフェライト系ステンレス鋼板。 Further, Cu: 0.1 to 1.0%, V: 0.01 to 0.10%, Ti: 0.001 to 0.05%, Nb: 0.001 to 0.05%, The ferritic stainless steel sheet according to claim 1, comprising one or more selected from Mo: 0.1 to 0.5% and Co: 0.01 to 0.2%.
  3. 　質量％で、さらに、Ｍｇ：０．０００２～０．００５０％、Ｃａ：０．０００２～０．００２０％、Ｂ：０．０００２～０．００５０％、ＲＥＭ：０．０１～０．１０％のうちから選ばれる１種または２種以上を含む請求項１または２に記載のフェライト系ステンレス鋼板。 Further, Mg: 0.0002 to 0.0050%, Ca: 0.0002 to 0.0020%, B: 0.0002 to 0.0050%, REM: 0.01 to 0.10%. The ferritic stainless steel sheet according to claim 1 or 2, comprising one or more selected from among them.
  4. 　請求項１～３のいずれか一項に記載のフェライト系ステンレス鋼板の製造方法であって、鋼スラブに対して、熱間圧延を施した後、９００～１１００℃の温度範囲で５秒～１５分間保持する焼鈍を行い、次いで冷間圧延を施した後、８００～９００℃の温度範囲で５秒～５分間保持する焼鈍を行うフェライト系ステンレス鋼板の製造方法。 The method for producing a ferritic stainless steel sheet according to any one of claims 1 to 3, wherein the steel slab is hot-rolled and then subjected to a temperature range of 900 to 1100 ° C for 5 seconds to 15 seconds. A method for producing a ferritic stainless steel sheet, in which annealing is performed for a minute, followed by cold rolling, followed by annealing in a temperature range of 800 to 900 ° C. for 5 seconds to 5 minutes.
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Cited By (3)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN109722508A (en) * 2017-10-27 2019-05-07 杰富意钢铁株式会社 Ferrite series stainless steel plate and its manufacturing method
CN110546293A (en) * 2017-04-25 2019-12-06 杰富意钢铁株式会社 Ferritic stainless steel sheet and method for producing same
JPWO2022085708A1 (en) * 2020-10-23 2022-04-28

Families Citing this family (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
KR101949629B1 (en) * 2014-12-11 2019-02-18 제이에프이 스틸 가부시키가이샤 Stainless steel and production method therefor
TWI801538B (en) * 2018-03-27 2023-05-11 日商日鐵不銹鋼股份有限公司 Ferritic stainless steel, method for producing the same, ferritic stainless steel sheet, method for producing the same, and members for fuel cell
CN111936654B (en) * 2018-03-30 2022-01-18 日铁不锈钢株式会社 Ferritic stainless steel having excellent ridging resistance
CN108315651B (en) * 2018-04-11 2020-02-04 山西太钢不锈钢股份有限公司 Continuous cold rolling annealing pickling method for ultra-pure ferrite stainless steel cold-rolled strip steel
US11377702B2 (en) 2018-07-18 2022-07-05 Jfe Steel Corporation Ferritic stainless steel sheet and method of producing same
JP6617858B1 (en) * 2018-07-18 2019-12-11 Ｊｆｅスチール株式会社 Ferritic stainless steel sheet and manufacturing method thereof
JP6881666B2 (en) * 2018-10-19 2021-06-02 Ｊｆｅスチール株式会社 Manufacturing method of ferritic stainless steel sheet

Citations (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPH0892652A (en) * 1994-09-22 1996-04-09 Nikko Kinzoku Kk Production of stainless steel sheet
WO2000060134A1 (en) * 1999-03-30 2000-10-12 Kawasaki Steel Corporation Ferritic stainless steel plate
JP2007119847A (en) * 2005-10-27 2007-05-17 Jfe Steel Kk Cold-rolled ferritic stainless steel sheet having excellent press formability and its production method
KR20100058849A (en) * 2008-11-25 2010-06-04 주식회사 포스코 Hot-rolled ferritic stainless steel sheet with excellent surface quality and method of manufacturing the same
JP2013227659A (en) * 2012-03-22 2013-11-07 Nippon Steel & Sumikin Stainless Steel Corp Ferritic stainless steel sheet excellent in scale peeling resistance and method for producing the same

Family Cites Families (18)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS471878Y1 (en) 1969-02-03 1972-01-22
JPS581801A (en) 1981-06-26 1983-01-07 Sony Corp Record player
JPS584881A (en) 1981-06-26 1983-01-12 三菱レイヨン株式会社 Dyeing of triacetate fiber
JP3468156B2 (en) * 1999-04-13 2003-11-17 住友金属工業株式会社 Ferritic stainless steel for automotive exhaust system parts
JP2001271143A (en) 2000-03-28 2001-10-02 Nisshin Steel Co Ltd Ferritic stainless steel excellent in ridging resistance and its production method
JP2005271261A (en) * 2004-03-23 2005-10-06 Riso Kagaku Corp Stencil printing method and water-based ink for stencil printing
JP2006274436A (en) * 2005-03-30 2006-10-12 Jfe Steel Kk Ferritic stainless sheet sheet and steel tube for bent tube having sectional shape for components
JP5453747B2 (en) 2008-08-25 2014-03-26 Ｊｆｅスチール株式会社 Stainless cold-rolled steel sheet excellent in punching processability and manufacturing method thereof
KR101279052B1 (en) * 2009-12-23 2013-07-02 주식회사 포스코 Ferritic stainless steel sheet with excellent ridging resistance and manufacturing method thereof
JP5659061B2 (en) * 2011-03-29 2015-01-28 新日鐵住金ステンレス株式会社 Ferritic stainless steel sheet excellent in heat resistance and workability and manufacturing method thereof
EP2722411B1 (en) * 2011-06-16 2020-04-08 Nippon Steel & Sumikin Stainless Steel Corporation Ferritic stainless steel plate which has excellent ridging resistance and method of production of same
ES2651071T3 (en) * 2012-01-30 2018-01-24 Jfe Steel Corporation Ferritic Stainless Steel Sheet
IN2015DN01886A (en) * 2012-12-07 2015-08-07 Jfe Steel Corp
KR101522077B1 (en) * 2012-12-20 2015-05-20 주식회사 포스코 Manufacturing method of ferritic stainless steel sheet with excellent ridging resistance
PL2952602T3 (en) * 2013-02-04 2020-09-07 Nippon Steel Stainless Steel Corporation Ferritic stainless steel sheet which is excellent in workability and method of production of same
US10450623B2 (en) * 2013-03-06 2019-10-22 Nippon Steel & Sumikin Stainless Steel Corporation Ferritic stainless steel sheet having excellent heat resistance
CN104685086B (en) * 2013-03-18 2017-03-08 杰富意钢铁株式会社 Ferrite series stainless steel plate
WO2014157066A1 (en) * 2013-03-25 2014-10-02 新日鐵住金ステンレス株式会社 Ferritic stainless steel sheet with excellent blanking workability and process for manufacturing same

Patent Citations (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPH0892652A (en) * 1994-09-22 1996-04-09 Nikko Kinzoku Kk Production of stainless steel sheet
WO2000060134A1 (en) * 1999-03-30 2000-10-12 Kawasaki Steel Corporation Ferritic stainless steel plate
JP2007119847A (en) * 2005-10-27 2007-05-17 Jfe Steel Kk Cold-rolled ferritic stainless steel sheet having excellent press formability and its production method
KR20100058849A (en) * 2008-11-25 2010-06-04 주식회사 포스코 Hot-rolled ferritic stainless steel sheet with excellent surface quality and method of manufacturing the same
JP2013227659A (en) * 2012-03-22 2013-11-07 Nippon Steel & Sumikin Stainless Steel Corp Ferritic stainless steel sheet excellent in scale peeling resistance and method for producing the same

Cited By (9)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN110546293A (en) * 2017-04-25 2019-12-06 杰富意钢铁株式会社 Ferritic stainless steel sheet and method for producing same
CN110546293B (en) * 2017-04-25 2022-07-29 杰富意钢铁株式会社 Ferritic stainless steel sheet and method for producing same
US11401573B2 (en) 2017-04-25 2022-08-02 Jfe Steel Corporation Ferritic stainless steel sheet and method for manufacturing the same
CN109722508A (en) * 2017-10-27 2019-05-07 杰富意钢铁株式会社 Ferrite series stainless steel plate and its manufacturing method
CN109722508B (en) * 2017-10-27 2020-10-02 杰富意钢铁株式会社 Ferritic stainless steel sheet and method for producing same
JPWO2022085708A1 (en) * 2020-10-23 2022-04-28
WO2022085708A1 (en) * 2020-10-23 2022-04-28 日鉄ステンレス株式会社 Ferritic stainless steel, and method for manufacturing ferritic stainless steel
CN115917029A (en) * 2020-10-23 2023-04-04 日铁不锈钢株式会社 Ferritic stainless steel and method for producing ferritic stainless steel
JP7374338B2 (en) 2020-10-23 2023-11-06 日鉄ステンレス株式会社 Ferritic stainless steel and manufacturing method of ferritic stainless steel

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