Classifications

• • 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
  • C21D8/0205—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
  • 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
  • C22C38/001—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
  • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
• • 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
  • C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
• • 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
  • C22C38/10—Ferrous alloys, e.g. steel alloys containing 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
  • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
• • 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
  • C22C38/14—Ferrous alloys, e.g. steel alloys containing 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
  • C22C38/16—Ferrous alloys, e.g. steel alloys containing 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
• • 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
  • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
• • 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 steel sheet with excellent cold workability during forming and a method for manufacturing the sheet.
• Automotive parts, knives, and other mechanical parts are manufactured through working processes such as punching, bending, and pressing.
• working processes improvement of workability is required for a material carbon steel sheet, in order to improve product quality and stability and/or cost reduction.
• a carbon steel sheet is subjected to cold rolling and spheroidizing annealing, so as to produce a soft carbon steel sheet with excellent workability made of ferrite and spheroidized carbide.
• Many technologies for improving the workability of carbon steel sheets have been proposed so far.
• Patent Document 1 discloses a high-carbon steel sheet for precision punching and a method for producing the sheet, wherein the sheet comprises, in terms of % by mass, C: 0.15 to 0.90%, Si: 0.40% or less, Mn: 0.3 to 1.0%, P: 0.03% or less, total Al: 0.1% or less, Ti: 0.01 to 0.05%, B: 0.0005 to 0.0050%, N: 0.01% or less, and Cr: 1.2% or less, has a structure in which carbides having an average carbide grain size of 0.4 to 1.0 ⁇ m and a carbide spheroidization ratio of 80% or more are dispersed in a ferrite matrix, and has a notched tensile elongation of 20% or more.
• Patent Document 2 discloses a medium- to high-carbon steel sheet with excellent workability and a method for producing the sheet, wherein the sheet comprises C: 0.3 to 1.3 wt %, Si: 1.0 wt % or less, Mn: 0.2 to 1.5 wt %, P: 0.02 wt% or less, and S: 0.02 wt % or less, has a structure in which carbides are dispersed so that the relationship C GB /C IG ⁇ 0.8 holds between the carbide number C GB on the ferrite crystal grain boundary and the carbide number C IG in the ferrite crystal grains, and has a cross-sectional hardness of 160 HV or less.
• Patent Document 3 discloses a medium- to high-carbon steel sheet with excellent workability, wherein the sheet comprises C: 0.30 to 1.00 wt %, Si: 1.0 wt % or less, Mn: 0.2 to 1.5 wt %, P: 0.02 wt % or less, and S: 0.02 wt % or less, has a structure in which carbides are dispersed in ferrite so that the relationship C GB /C IG ⁇ 0.8 holds between the carbide number C GB on the ferrite crystal grain boundary and the carbide number C IG in the ferrite crystal grains, and simultaneously 90% or more of the total carbides are occupied by spheroidized carbides having a long axis/short axis of 2 or less.
• Patent Documents 1 to 3 describe that the greater the proportion of carbides in ferrite grains, the more the workability is improved.
• Patent Documents 2 and 3 consider that the deterioration of workability is caused by the low carbide spheroidization ratio of carbides precipitated on the grain boundary, but do not take into account the problem of improving the spheroidization ratio of grain boundary carbides.
• Techniques described in Patent Document 4 only specify the tissue factor, and Patent Document 4 does not discuss the relationship between workability and mechanical properties.
• Patent Document 5 The technology described in Patent Document 5 is an invention made by focusing on the relationship between fine blanking workability and the amount of carbide present in ferrite grains and ferrite grain size. However, Patent Document 5 does not discuss what effect the aggregate structure has on the plastic anisotropy.
• Patent Document 6 discloses a hot-rolled steel sheet in which the development of an aggregate structure otherwise developed by rolling is suppressed and a method for manufacturing the sheet. However, Patent Document 6 does not discuss the relationship between the aggregate structure other than the aggregate structure developed by rolling and the cold forgeability.
• Patent Document 7 is an invention made by considering that the hardness and the total elongation of a high-carbon hot-rolled steel sheet prior to quenching are greatly influenced by the cementite density in the ferrite grains.
• the hot-rolled steel sheet described in Patent Document 7 is characterized in that it has a microstructure composed of ferrite and cementite, said microstructure having a cementite density of 0.10 strips/ ⁇ m 2 or less in the ferrite grains.
• Patent Document 7 does not discuss what effect the aggregate texture has on the plastic anisotropy.
• Patent Document 8 The technology described in Patent Document 8 is an invention made by considering that the C eq value is related not only to mechanical properties and weldability but also to the fatigue crack growth rate in steels having a fine structure. Patent Document 8 discloses that by limiting the range of the C eq value to a range of 0.28% to 0.65%, the fatigue resistance of the steel material is improved and simultaneously weldability is secured. However, Patent Document 8 does not discuss what effect the aggregate texture has on the plastic anisotropy.
• Patent Document 1 Japanese Patent No. 4465057
• Patent Document 2 Japanese Patent No. 4974285
• Patent Document 3 Japanese Patent No. 5197076
• Patent Document 4 Japanese Patent No. 5194454
• Patent Document 5 Japanese Unexamined Patent Publication No. 2007-270331
• Patent Document 6 Japanese Unexamined Patent Publication No. 2009-263718
• Patent Document 7 Japanese Unexamined Patent Publication No. 2015-17294
• Patent Document 8 Japanese Unexamined Patent Publication No. 2004-27355
• the present inventors have conducted intensive and extensive studies on methods for solving the above-mentioned problems. As a result, the present inventors have found that by controlling the dispersion state of the carbide in the structure of the steel sheet before cold working through the optimization of the manufacturing conditions in the steps from hot rolling to annealing, the carbide can be precipitated on the ferrite boundary and simultaneously the aggregate structure in the hot rolled steel plate can be controlled, thereby leading to enhanced cold workability.
• a steel sheet having an excellent cold workability during forming comprising, in terms of % by mass:
• a balance of Fe and inevitable impurities wherein (a) a ratio of the number of carbides at a ferrite grain boundary relative to the number of carbides in the ferrite grain is more than 1, wherein (b) a diameter of the ferrite grain is 5 ⁇ m or more and 50 ⁇ m or less, wherein (c) an in-plane anisotropy
• V 0.001 to 0.10%
• Ta 0.001 to 0.10%
• a steel sheet with excellent cold workability during forming can be manufactured and provided.
• a steel sheet with excellent cold workability during forming according to the present invention (hereinafter may be referred to as “the inventive steel sheet”) comprises, in terms of % by mass:
• the ferrite grain diameter is 5 ⁇ m or more and 50 ⁇ m or less
• of the r value standardized according to JIS Z 2254 is 0.2 or less
• the Vickers hardness is 100 HV or more and 150 HV or less
• the method (hereinafter may be referred to as “the inventive method”) of the present invention for producing a steel sheet with excellent cold workability during forming is a method for producing the inventive steel sheet
• a hot-rolled steel strip that has been obtained by subjecting a steel strip having an ingredient composition of the inventive steel sheet to hot rolling by heating, followed by completing the finish hot rolling at a temperature range of 800° C. or higher and 900° C. or lower, and by coiling the resulting hot-rolled steel sheet at a temperature of 400° C. or higher and 550° C. or lower is, after pickling, subjected to two-step type annealing in which the sheet is retained in two temperature ranges, whereupon
• the hot-rolled steel sheet is subjected to a first step annealing performed by retaining said hot-rolled steel at a temperature range of 650° C. or higher and 720° C. or lower for 3 hours or longer and 60 hours or shorter, and then subjected to a second step annealing performed by retaining the hot-rolled steel at a temperature range of 725° C. or higher and 790° C. or lower for 3 hours or longer and 50 hours or shorter, and thereafter
• the sheet is cooled down to 650° C. or lower at a cooling rate of 1° C./hour or more and 30° C./hour or less.
• the percentage relating to the ingredient composition means % by mass.
• C is an element that forms carbide in steel, and is effective for strengthening steel and refining ferrite grains.
• C is set to 0.10% or more, and preferably 0.12% or more.
• C is set to 0.40% or less, and preferably 0.38% or less.
• Si is an element that acts as a deoxidizing agent and also affects the form of the carbide.
• it is necessary to generate an austenite phase during annealing in the two-step type annealing, and, after transiently dissolving the carbides, to cool gradually to promote the precipitation of carbides at the ferrite grain boundaries.
• the amount of Si may preferably be as small as possible. However, when it is reduced to less than 0.01%, the manufacturing cost increases. Therefore, Si is set to 0.01% or more.
• Si is set to 0.30% or less, and preferably 0.28% or less.
• Mn is an element that controls the figuration of carbides in the two-step type annealing. When its content is less than 0.30%, it is difficult to precipitate carbides at the ferrite grain boundaries in slow cooling after the second-step annealing. Therefore, Mn is set to 0.30% or more, and preferably 0.33% or more.
• Mn is set to 1.00% or less, and preferably 0.96% or less.
• P is an element that segregates at the ferrite grain boundaries and suppresses the formation of grain boundary carbides.
• the amount of P may preferably be as small as possible. However, when P is reduced to less than 0.0001% in the refining process, the refining cost may greatly increase. Therefore, it is set to 0.0001% or more, and preferably 0.0013% or more.
• P is set to 0.020% or less, and preferably 0.018% or less.
• S is an element that forms a non-metallic inclusion such as MnS. Since a non-metallic inclusion serves as the starting point for break generation during cold forging, the amount of S may preferably be as small as possible. However, when it is reduced to less than 0.0001%, the refining cost greatly increases. Therefore, S is set to 0.0001% or more, and preferably 0.0012% or more.
• S is set to 0.010% or less, and preferably 0.007% or less.
• Al is an element that acts as a deoxidizing agent for steel and stabilizes ferrite. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Al is set to 0.001% or more, and preferably 0.004% or more.
• Al is set to 0.10% or less, and preferably 0.08% or less.
• the inventive steel sheet may contain one or a plurality of N: 0.0001 to 0.010%, 0: 0.0001 to 0.020%, Cr: 0.001 to 0.50%, Mo: 0.001 to 0.10%, Nb: 0.001 to 0.10%, V: 0.001 to 0.10%, Cu: 0.001 to 0.10%, W: 0.001 to 0.10%, Ta: 0.001 to 0.10%, Ni: 0.001 to 0.10%, Sn: 0.001 to 0.050%, Sb: 0.001 to 0.050%, As: 0.001 to 0.050%, Mg: 0.0001 to 0.050%, Ca: 0.001 to 0.050%, Y: 0.001 to 0.050%, Zr: 0.001 to 0.050%, La: 0.001 to 0.050%, and Ce: 0.001 to 0.050%, in order to improve the properties of the inventive steel sheet.
• N is an element that, when present in large amounts, causes the embrittlement of ferrite.
• the amount of N may preferably be as small as possible. However, when it is reduced to less than 0.0001%, the refining cost greatly increases. Therefore, N should be 0.0001% or more, and preferably 0.0006% or more. On the other hand, when it exceeds 0.010%, ferrite embrittles and the cold forgeability deteriorates. Therefore, N should be 0.010% or less, and preferably 0.007% or less.
• O is an element that, when present in large amounts, forms coarse oxides in steel.
• the amount of O may preferably be as small as possible. However, when it is reduced to less than 0.0001%, the refining cost increases greatly. Therefore, O is set to 0.0001% or more, and preferably 0.0011% or more.
• O is set to 0.020% or less, and preferably 0.017% or less.
• Cr is an element which enhances quenchability and contributes to the improvement of strength and which is thickened to carbide and forms stable carbide even in the austenitic phase. When its content is less than 0.001%, the sufficient effect of improving quenchability cannot be obtained. Therefore, Cr is set to 0.001% or more, and preferably 0.007% or more. On the other hand, when it exceeds 0.50%, the carbide becomes stabilized thereby delaying the dissolution of the carbide during quenching, and thus, it is feared that the desired quenching strength may not be achieved. Therefore, Cr is set to 0.50% or less, and preferably 0.45% or less.
• Mo is an element effective for controlling the figuration of carbides. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Mo is set to 0.001% or more, and preferably 0.010% or more. On the other hand, when it exceeds 0.10%, the in-plane anisotropy of the r value deteriorates and the cold workability deteriorates. Therefore, Mo is set to 0.10% or less, and preferably 0.08% or less.
• Nb is an element which is effective for controlling the figuration of carbides and which refines the structure, thereby contributing to the enhancement of its toughness.
• Nb should be 0.001% or more, and preferably 0.004% or more.
• Nb is set to 0.10 or less, and preferably 0.08% or less.
• V 0.001 to 0.10%
• V is an element which is effective for controlling the figuration of carbides and which refines the structure, thereby contributing to the enhancement of its toughness.
• V is set to 0.001% or more, and preferably 0.004% or more.
• V is set to 0.10 or less, and preferably 0.08% or less.
• Cu is an element which segregates at the ferrite crystal grain boundary and forms fine precipitates thereby to contribute to the enhancement of strength. When its content is less than 0.001%, a sufficient effect of enhancing strength cannot be obtained. Therefore, Cu is set to 0.001% or more, and preferably 0.005% or more. On the other hand, when it exceeds 0.10%, red heat embrittlement occurs and the productivity by hot rolling decreases. Therefore, Cu is set to 0.10% or less, and preferably 0.08% or less.
• W is also an element effective for controlling the figuration of carbides.
• W is set to 0.001% or more, and preferably 0.003% or more.
• W is set to 0.10 or less, and preferably 0.08% or less.
• Ta 0.001 to 0.10%
• Ta is also an element effective for controlling the figuration of carbides.
• W is set to 0.001% or more, and preferably 0.005% or more.
• Ta is set to 0.10 or less, and preferably 0.08% or less.
• Ni is an element effective for improving the toughness of parts. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Ni is set to 0.001% or more, and preferably 0.003% or more. On the other hand, when it exceeds 0.10%, the number ratio of grain boundary carbides decreases and the cold forgeability deteriorates. Therefore, Ni is set to 0.10% or less, and preferably 0.08% or less.
• Sn is an element contaminated from a steel raw material (scrap). It segregates at the grain boundary, leading to the decreased number ratio of grain boundary carbides. Therefore, its content may preferably be as small as possible. However, when it is reduced to less than 0.001%, the refining cost will be greatly increased. Therefore, Sn is set to 0.001% or more, and preferably 0.002% or more. On the other hand, when it exceeds 0.050%, ferrite embrittles and cold forgeability deteriorates. Therefore, Sn is set to 0.050% or less, and preferably 0.040% or less.
• Sb is an element contaminated from a steel raw material (scrap). It segregates at the grain boundary, leading to the decreased number ratio of grain boundary carbides. Therefore, its content may preferably be as small as possible. However, when it is reduced to less than 0.001%, the refining cost will be greatly increased. Therefore, Sb is set to 0.001% or more, preferably 0.002% or more. On the other hand, when it exceeds 0.050%, the cold forgeability deteriorates. Therefore, Sb is set to 0.050% or less, and preferably 0.040% or less.
• Sn and Sb As is an element contaminated from a steel raw material (scrap). It segregates at the grain boundary, thereby leading to a decrease in the number ratio of grain boundary carbides. Therefore, its content may preferably be as small as possible. However, when it is reduced to less than 0.001%, the refining cost increases greatly. Therefore, As is set to 0.001% or more, and preferably 0.002% or more. On the other hand, when it exceeds 0.050%, the number ratio of the grain boundary carbides decreases and the cold forgeability deteriorates. Therefore, As is set to 0.050% or less, and preferably 0.040% or less.
• Mg is an element that can control the figuration of sulfides with the addition of its trace amount. When its content is less than 0.0001%, a sufficient addition effect cannot be obtained. Therefore, Mg is set to 0.0001% or more, and preferably 0.0008% or more. On the other hand, when it exceeds 0.050%, ferrite embrittles and the cold forgeability deteriorates. Therefore, Mg is set to 0.050% or less, and preferably 0.040% or less.
• Ca is an element that can control the figuration of sulfides with the addition of its trace amount. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Ca is set to 0.001% or more, and preferably 0.003% or more. On the other hand, when it exceeds 0.050%, coarse Ca oxides are formed, which serve as starting points of break generation during cold forging. Therefore, Ca is set to 0.050% or less, and preferably 0.040% or less.
• Y is an element that can control the figuration of sulfides with the addition of its trace amount. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Y is set to 0.001% or more, and preferably 0.003% or more. On the other hand, when it exceeds 0.050%, coarse Y oxides are formed, which serve as starting points of break generation during cold working. Therefore, Y is set to 0.050% or less, and preferably 0.035% or less.
• Zr is an element that can control the figuration of sulfides with the addition of its trace amount.
• Zr is set to 0.001% or more, and preferably 0.004% or more.
• Zr is set to 0.050% or less, and preferably 0.045% or less.
• La is an element that can control the figuration of sulfides with the addition of its trace amount, but it is also an element that segregates at the grain boundary and causes a decrease in the number ratio of grain boundary carbides.
• its content is less than 0.001%, a sufficient effect of controlling figuration cannot be obtained. Therefore, La is set to 0.001% or more, and preferably 0.004% or more.
• La is set to 0.050% or less, and preferably 0.045% or less.
• Ce is an element that can control the figuration of sulfides with the addition of its trace amount, but it is also an element that segregates at the grain boundary and causes a decrease in the number ratio of grain boundary carbides.
• Ce is set to 0.001% or more, and preferably 0.004% or more.
• Ce is set to 0.050% or less, and preferably 0.045% or less.
• the remainder of the ingredient composition of the inventive steel sheet is Fe and unavoidable impurities.
• the inventive steel sheet has excellent cold workability during forming, because, in addition to the above ingredient composition, it was found, as a result of optimum hot rolling and annealing, that
• the ferrite grain diameter is 5 ⁇ m or more and 50 ⁇ m or less
• the Vickers hardness is 100 HV or more and 150 HV or less
• the inventive steel sheet has a structure which is substantially composed of ferrite and carbide, and in which the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain exceeds 1.
• Carbides are, in addition to cementite (Fe 3 C) that is a compound of iron and carbon, compounds obtained by replacing Fe in cementite with an element such as Mn and Cr, and alloy carbides (M 23 C 6 , M 6 Co, MC, etc., wherein M is Fe and another additive metal element).
• a shear band is formed in the macrostructure of the steel sheet, and slip deformation is generated and concentrated in the vicinity of the shear band.
• the slip deformation involves propagation of dislocations, and regions with high dislocation density are formed in the vicinity of the shear band. As the strain amount applied to the steel sheet increases, the slip deformation is promoted and thereby the dislocation density increases. In cold forging, strong processing exceeding an equivalent strain of 1 is applied.
• shear band formation is a phenomenon in which a slip generated in one crystal grain crosses the crystal grain boundary and propagates continuously to an adjacent crystal grain. Therefore, in order to suppress the formation of a shear band, it is necessary to prevent the propagation of slippage beyond crystal the grain boundary.
• Carbides in the steel sheet are tenacious particles that hinder slippage. Therefore, the presence of carbides at the ferrite grain boundaries would make it possible, for the first time, to suppress the formation of a shear band and thereby to improve cold forgeability.
• the present inventors have found that cold forgeability can be evaluated by using, as an index, the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain, and that cold forgeability can be remarkably improved when the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain is more than 1.
• any of buckling, folding and convolution of a steel sheet that occurs during cold working is caused by the localization of strain accompanying the formation of a shear band. Therefore, by allowing the carbide to exist at the ferrite grain boundaries, the formation of the shear band and the localization of strain can be alleviated and the generation of buckling, folding and convolution can be suppressed.
• the carbide spheroidization ratio on the crystal grain boundary may preferably be 80% or more, and more preferably 90% or more.
• the average particle diameter of the carbide in the ferrite grain and the carbide at the ferrite grain boundaries is less than 0.1 ⁇ m, the hardness of the steel sheet remarkably increases and the workability deteriorates. Therefore, the average particle diameter of the carbide may preferably be 0.1 ⁇ m or more, and more preferably 0.17 ⁇ m or more. On the other hand, when the average particle diameter of the carbide exceeds 2.0 ⁇ m, fissures occur with the coarse carbide serving as a starting point during cold working, and thus the cold workability deteriorates. Therefore, the average particle diameter of the carbide may preferably be 2.0 ⁇ m or less, and more preferably 1.95 ⁇ m or less.
• Observation of the carbide is carried out by a scanning electron microscope. Prior to observation, samples for structure observation are polished by wet polishing with emery paper and polishing with diamond abrasive grains having an average particle size of 1 ⁇ m. After polishing the observation surface to a mirror finish, the structure is etched with a 3% nitric acid-alcohol solution.
• magnification for observation, within 3000 times, a magnification capable of discriminating between ferrite and carbide is selected. At the selected magnification, eight images with a viewing field of 30 ⁇ m ⁇ 40 ⁇ m are randomly photographed at the 1 ⁇ 4 plate layer thickness.
• the area of each carbide contained in the region is measured in detail by an image analysis software represented by Mitsuya Shoji Co. Ltd. (Win ROOF).
• the spheroidization ratio of the carbide was determined by approximating the carbide to an ellipse having an equal area and equal moment of inertia, and then by calculating the proportion of the carbides in which the ratio of the maximum length to the maximum length in the perpendicular direction is less than 3.
• carbides having an area of 0.01 ⁇ m 2 or more among the carbides in grains and grain boundaries were counted and the carbides having an area of 0.01 ⁇ m 2 or less were excluded from evaluation.
• the number of carbides present on the ferrite grain boundary was counted, and from the total number of carbides the number of carbides in the ferrite grain was determined by subtracting the number of carbides on the ferrite grain boundary. Based on the measured number, the ratio of the number of carbides on the grain boundary relative to the number of carbides in the ferrite grain was determined.
• the ferrite grain diameter is 5 ⁇ m or more and 50 ⁇ m or less:
• the cold workability can be improved by setting the ferrite grain diameter to 5 ⁇ m or more.
• the ferrite grain size is set to 5 ⁇ m or more, and preferably 7 ⁇ m or more.
• the ferrite grain size is set to 50 ⁇ m or less, and preferably 37 ⁇ m or less.
• the ferrite grain diameter is measured in the above-described polishing method, wherein the observation surface of the sample is polished to a mirror surface, followed by etching with a 3% nitric acid-alcohol solution. The structure of the observation surface is then examined with an optical microscope or a scanning electron microscope, and a line segment method is then applied to the image photographed to determine the ferrite grain diameter.
• of the plastic strain ratio (r value) of the steel sheet is measured in a method in accordance with JIS Z 2254.
• the r value (0° direction: r 0 , 45° direction: r 45 , 90° direction: r 90 ) measured by taking test strips from each direction of 0° direction, 45° direction and 90° direction with respect to the rolling direction was used to calculate the following equation.
• of the plastic strain ratio (r value) of the steel sheet is set to 0.2 or less.
• the Vickers hardness is 100 HV or more and 150 HV or less:
• the Vickers hardness of the steel sheet By setting the Vickers hardness of the steel sheet to 100 HV or more and 150 HV or less, the cold workability can be improved. When the Vickers hardness is less than 100 HV, buckling can easily occur during cold working. Therefore, the Vickers hardness is set to 100 HV or more, and preferably 110 HV or more.
• the Vickers hardness exceeds 150 HV, the ductility decreases and the internal breaking tends to occur easily during cold forging. Therefore, the Vickers hardness is set to 150 HV or less, and preferably 146 HV or less.
• the draw formability during cold forging In cold forging, in addition to controlling the figuration of carbides, the draw formability during cold forging must be secured. In order to improve the draw formability during cold forging, plastic anisotropy such as in-plane anisotropy
• One surface of a hot-rolled steel plate is ground to a 1 ⁇ 2 plate thickness surface in parallel to the surface to expose a 1 ⁇ 2 plate thickness surface, followed by the analysis of the 1 ⁇ 2 plate thickness surface by X-ray diffraction.
• X-ray diffraction X-ray diffraction by Mo bulb may be used.
• Diffraction intensities of diffraction orientations ⁇ 110 ⁇ , ⁇ 220 ⁇ , ⁇ 211 ⁇ and ⁇ 310 ⁇ by reflection are obtained, and based thereon, the orientation distribution function (ODF) is created.
• the X-ray diffraction intensity ratio is determined by using the diffraction intensity data of the 1 ⁇ 2 plate thickness surface obtained from the ODF and the diffraction intensity data of random orientation of the hot-rolled steel sheet. Specifically, as a standard sample in which the metallic structure has no accumulation in a specific direction, a sample obtained by sintering powder iron of a hot-rolled steel sheet to be measured or the powder before sintering is used to determine the diffraction intensity under the same conditions as when the diffraction intensity data of the 1 ⁇ 2 plate thickness surface was obtained.
• the part to be collected as the standard sample is not particularly limited and may be any part of the hot-rolled steel sheet.
• the X-ray diffraction intensity ratio in a specific orientation is a numerical value obtained by dividing the diffraction intensity in the specific direction of the 1 ⁇ 2 plate thickness surface obtained from the ODF by the diffraction intensity of the standard sample.
• the X-ray diffraction intensity ratio of the ⁇ 311 ⁇ ⁇ 011> orientation obtained by the above-described ODF analysis is set to I1, it is necessary that this I1 is 3.0 or less, and preferably 2.5 or less for the random aggregate structure during hot rolling.
• this I1 is 3.0 or less, and preferably 2.5 or less for the random aggregate structure during hot rolling.
• the manufacturing method according to the present invention is characterized in that the hot rolling and the annealing are consistently managed to control the structure.
• the steel strip After continuously casting a steel strip having a predetermined ingredient composition, the steel strip is subjected to hot rolling by heating to complete finish hot rolling at a temperature range of 800° C. or higher to 900° C. or lower, coiled at 400° C. or higher and 550° C. or lower to obtain a hot-rolled steel sheet.
• the hot-rolled steel sheet is, after pickling, subjected to a two-step type annealing in which the hot-rolled steel sheet is maintained in two temperature ranges, whereupon
• the hot-rolled steel sheet is subjected to a first step annealing performed by retaining said hot-rolled steel at a temperature range of 650° C. or higher and 720° C. or lower for 3 hours or longer and 60 hours or shorter, and then subjected to a second step annealing performed by retaining the hot-rolled steel at a temperature range of 725° C. or higher and 790° C. or lower for 3 hours or longer and 50 hours or shorter, and thereafter
• the hot-rolled steel sheet is cooled to 650° C. or lower at a cooling rate of 1° C./hour or more and 30° C./hour or less,
• a structure composed of fine pearlite and bainite can be formed as the structure of the steel sheet.
• Heating temperature of a steel strip 1000° C. or higher and 1250° C. or lower
• the heating temperature of the steel strip subjected to hot rolling may preferably be 1000° C. or higher and 1250° C. or lower, and the heating time may preferably be 0.5 hour or longer and 3 hours or shorter.
• the heating temperature When the heating temperature is lower than 1000° C. or the heating time is shorter than 0.5 hour, the microsegregation and/or macrosegregation formed by casting are not eliminated, and regions in which Si, Mn, etc., are locally concentrated inside the steel material may remain, and thus the impact resistance property of the steel material is lowered. Therefore, the heating temperature may preferably be 1000° C. or higher, and preferably 0.5 hour or longer.
• the heating temperature exceeds 1250° C. or the heating time exceeds 3 hours, decarburization from the surface layer of the steel strip becomes conspicuous, and austenite grains in the surface layer grow abnormally during heating before carburizing and quenching, and the impact resistance property of the steel strip is deteriorated.
• the heating temperature may preferably be 1250° C. or lower, and the heating time may preferably be 3 hours or shorter.
• Finish hot rolling temperature 800° C. or higher and 900° C. or lower
• Finish hot rolling is completed at 800° C. or higher and 900° C. or lower.
• the finish hot rolling temperature is set to 800° C. or higher, and preferably 820° C. or higher.
• the finish hot rolling temperature exceeds 900° C.
• thick scales are generated during plate passing on the ROT (Run Out Table)
• scratches are generated on the surface of the steel sheet due to the scale
• cracks are generated starting from scratches when an impact load is applied after cold forging and carburizing and annealing, leading to reduced impact resistance property of the steel sheet. Therefore, the finish hot rolling temperature is set to 900° C. or lower, and preferably 880° C. or lower.
• Cooling rate on ROT 10° C./sec or more and 100° C./sec or less
• the cooling rate at the time of cooling the hot-rolled steel sheet on the ROT after finish hot rolling may preferably be 10° C./sec or more and 100° C./sec or less.
• the cooling rate is set to 10° C./sec or more, and more preferably 20° C./sec or more.
• the steel sheet is cooled at a cooling rate exceeding 100° C./sec from the surface layer to the inside of the steel sheet, the outermost layer part of the steel sheet is excessively cooled, and a low-temperature transformed structure such as bainite or martensite is formed.
• the cooling rate may preferably be 100° C./sec or less.
• the above cooling rate refers to the cooling capacity from the cooling facility at each water injection zone from the point at which the hot-rolled steel sheet after the finish hot rolling is cooled at the water injection zone after passing through the water-free zone to a point at which it is cooled to the coiling target temperature on the ROT, and does not refer to the average cooling rate from the water injection starting point to the temperature at which it is coiled by the coiling device.
• Coiling temperature 400° C. or higher and 550° C. or lower
• the coiling temperature is set to 400° C. or higher and 550° C. or lower.
• the coiling temperature is set to 400° C. or higher, and preferably 430° C. or higher.
• the coiling temperature exceeds 550° C.
• pearlite having a large lamellar spacing is generated and thick needle-shaped carbides having high thermal stability are formed, and even after the two-step type annealing, needle-shaped carbides remain. Since fissures are generated during cold working with these needle-shaped carbides as a starting point, the coiling temperature is set to 550° C. or lower, and preferably 520° C. or lower.
• the hot-rolled coil manufactured under the above conditions is annealed, after pickling, in a two-step type annealing which retains the coil in two temperature ranges.
• the first-step annealing and the second-step annealing may be either box annealing or continuous annealing.
• the first step annealing is carried out in a temperature range of the A ci point or lower to coarsen carbides and enrich alloy elements to increase the thermal stability of carbides. Thereafter, the temperature is raised to a range from A C1 point or higher to A 3 point or lower to generate austenite in the structure.
• Temperature range 650° C. or higher and 720° C. or lower
• Retention time 3 hours or longer and 60 hours or shorter
• the annealing temperature is set to 650° C. or higher and 720° C. or lower.
• the annealing temperature of the first step is lower than 650° C., the stability of the carbide becomes insufficient and it becomes difficult to allow the carbide to remain in the austenite in the second step annealing. Therefore, the temperature of the first step annealing is set to 650° C. or higher, and preferably 670° C. or higher.
• the first step annealing temperature is set to 720° C. or lower, and preferably 700° C. or lower.
• the retention time at the first step is 3 hours or longer and 60 hours or shorter.
• the retention time of the first step is set to 3 hours or longer.
• the retention time of the first step exceeds 60 hours, improvement of the stability of the carbide cannot be expected and furthermore the productivity is lowered. Therefore, the retention time of the first step is set to 60 hours or shorter, and preferably 55 hours or shorter.
• the annealing atmosphere is not limited to a specific atmosphere.
• it may be either a nitrogen atmosphere having a nitrogen content of 95% or more, a hydrogen atmosphere having a hydrogen content of 95% or more, or an atmospheric atmosphere.
• Temperature range 725° C. or higher and 790° C. or lower
• Retention time 3 hours or longer and 50 hours or shorter
• the annealing temperature is set to 725° C. or higher and 790° C. or lower.
• the second-step annealing temperature is set to 725° C. or higher, and preferably 745° C. or higher.
• the second-step annealing temperature is set to 790° C. or lower, and preferably 770° C. or lower.
• the retention time of the second step is set to 3 hours or longer and 50 hours or shorter.
• the retention time of the second step is set to 3 hours or longer, and preferably 5 hours or longer.
• the retention time of the second step exceeds 50 hours, it becomes difficult to allow the carbide to remain in the austenite. Therefore, the retention time of the second step is set to 50 hours or shorter, and preferably is 46 hours or shorter.
• the annealing atmosphere is not limited to a specific atmosphere.
• it may be either a nitrogen atmosphere having a nitrogen content of 95% or more, a hydrogen atmosphere having a hydrogen content of 95% or more, or an atmospheric atmosphere.
• the hot-rolled steel sheet After completion of the two-step type annealing, the hot-rolled steel sheet is cooled, whereupon it is cooled to 650° C. at a cooling rate of 1° C./hour or more to 30° C./hour or less.
• Cooling rate to a temperature of 650° C. or lower: 1° C./hour or more and 30° C./hour or less
• the temperature range for controlling the structure change by slow cooling is sufficient up to 650° C., it is only necessary to control the cooling rate in the temperature range up to 650° C. After reaching a temperature of 650° C. or lower, it may be cooled to room temperature within the above range without controlling the cooling rate.
• the cooling rate is slow in order to gradually cool the austenite produced in the second step annealing to transform into ferrite and allow carbon to be adsorbed to the carbides remaining in the austenite.
• the cooling rate is 1° C./hour or more, and preferably 5° C./hour.
• the cooling rate exceeds 30° C./hour, austenite transforms to pearlite, the hardness of the steel sheet increases, the cold forgeability deteriorates, and the impact resistance property of the steel sheet after carburizing quenching and tempering decreases. Therefore, the cooling rate is set to 30° C./hour or less, and preferably 26° C./hour or less.
• a steel sheet with excellent cold workability during forming can be produced in which the ingredient composition is, in terms of % by mass, comprising: C: 0.10 to 0.40%, Si: 0.01 to 0.30%, Mn: 0.30 to 1.00% , P: 0.0001 to 0.020%, S: 0.0001 to 0.010%, and Al: 0.001 to 0.10%, the balance being Fe and unavoidable impurities, the metal structure is substantially composed of ferrite and spheroidized carbides, and (a) the ratio of the number of carbides at the ferrite grain boundary to the number of carbides in the ferrite grain exceeds 1, (b) the ferrite grain size is 5 ⁇ m or more and 50 ⁇ m or less, (c) the in-plane anisotropy
• the cross-sectional shrinkage percentage is defined by the following formula (1).
• a large value of this ratio means that the local deformability is high, and as the value of the formula (1) increases, the workability of the steel sheet increases.
• Sectional shrinkage percentage (%) 100-(cross-sectional area at tensile fracture/initial cross-sectional area) ⁇ 100 Equation (1)
• the present invention is characterized in that by rolling control and heat treatment after rolling, a structure in which carbides (that is, cementite) are uniformly dispersed is formed, so that the crystal anisotropy can be eliminated. Therefore, in the present invention, the random intensity ratio of the ⁇ 311 ⁇ ⁇ 011> orientation at the 1 ⁇ 2 plate thickness portion of the steel sheet can be made 3.0 or less.
• a continuous cast strip (steel ingot) having the ingredient composition shown in Table 1 was subjected to hot rolling under the conditions shown in Table 2 to produce a hot-rolled coil having a thickness of 3.0 mm.
• the steel type described as “Developed steel” in the column of “Remarks” in Table 1 has a composition included in the composition range of the steel sheet according to the present invention.
• the steel type described as “Comparative steel” in the column of “Remarks” in Table 1 has a composition outside the composition range of the steel sheet according to the present invention.
• the ingredients that do not satisfy the composition conditions of the steel sheet according to the present invention are underlined.
• a sample for characterization was prepared as follows: a hot-rolled coil, after pickling, was placed in a box-type annealing furnace, the atmosphere was controlled to 95% hydrogen-5% nitrogen, the coil was heated from room temperature to 705° C. and was retained for 36 hours to make the temperature distribution uniform in the hot-rolled coil. The coil was then heated to 760° C. and retained at 760° C. for 10 hours, and was then cooled to 650° C. at a cooling rate of 10° C./hour, then furnace-cooled to room temperature to prepare the sample for characterization. The structure of the sample was measured by the method described above.
• the cold workability was evaluated using the notched tensile test and the in-plane anisotropy of the r value.
• a notched tensile test strip was taken from an as-annealed material with a thickness of 3 mm, and a tensile test was performed in the rolling direction to determine the cross-sectional shrinkage percentage, and the local deformability was evaluated. When the cross-sectional shrinkage percentage is 40% or more, it was rated as superior.
• the in-plane anisotropy of the r value was rated as superior when the in-plane anisotropy
• I1 X-ray diffraction intensity ratio
• Table 2 shows, for each of the samples prepared, the results of the carbide diameter, the ferrite grain diameter, the Vickers hardness, the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain, the cross-sectional shrinkage percentage, the X-ray diffraction intensity ratio of ⁇ 311 ⁇ ⁇ 011> and in-plane anisotropy.
• those indicated as “Inventive steel” in the Remarks column satisfy the requirements of the steel sheet according to the present invention
• those indicated as “Comparative steel” in the Remarks column do not satisfy the requirements of the steel sheet according to the present invention.
• the measurement results that do not satisfy the requirements of the steel sheet according to the present invention and the manufacturing conditions that do not satisfy the requirements of the steel sheet manufacturing method according to the present invention are underlined.
• the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain exceeds 1, and the Vickers hardness is 150 HV or less.
• the cross-sectional shrinkage percentage exceeds 40% and the in-plane anisotropy
• these steels can be suitably used for cold working.
• the Comparative steel A-1 since the Al content is high and the A3 point decreased, recrystallization during finish hot rolling was inhibited and
• the contents of Mo and Cr are high, recrystallization during finish hot rolling was inhibited, and
• the comparative steels K-1 and N-1 the content of S or Mn is high, coarse MnS was formed in the steel, and the cold workability is low.
• the Comparative steel M-1 the content of Si was high and hardness increased, and thus cold workability is low. Also, in the Comparative steel M-1, since the A3 point rose, recrystallization during finish hot rolling was hindered and
• Comparative steel O-1 C is high, the volume fraction of carbides increased, a large amount of cracks as the starting point of fractures were generated, and the cross-sectional shrinkage percentage was low. Thus, the cold workability is low.
• the finish temperature of hot rolling was low and the productivity decreased.
• the finish temperature of hot rolling was high, and scale scratches were generated on the surface of the steel sheet.
• the coiling temperature of hot rolling was low, the low-temperature transformation structure such as bainite and martensite increased resulting in brittled steel, and breaks frequently occurred when the hot-rolled coil was discharged resulting in a decrease in productivity.
• the coiling temperature of hot rolling was high, thick pearlite with lamellar spacing and needle-shaped coarse carbides with high thermal stability were produced in the hot rolled structure. Since these carbides remained in the steel sheet even after the two-step type annealing, the cross-sectional shrinkage percentage was low and thus the cold workability is low.
• steel strips having the ingredient composition shown in Table 1 were heated at 1240° C. for 1.8 hours and then subjected to hot rolling. After completing finish hot rolling at 890° C., they were cooled to 520° C. at a cooling rate of 45° C./sec on ROT and coiled at 510° C. to produce a hot-rolled coil with a thickness of 3.0 mm. And under the conditions shown in Table 3, a hot-rolled sheet-annealed sample with a thickness of 3.0 mm was prepared.
• the carbide diameter, the ferrite grain diameter, the Vickers hardness, the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain, the cross-sectional shrinkage percentage, the X-ray diffraction intensity ratio of ⁇ 311 ⁇ ⁇ 011> and the in-plane anisotropy were determined in the same manner as the inventive steels and the comparative steels in Table 2. The results are shown in Table 3.
• the ratio of the number of carbides at the ferrite grain boundary to the number of carbides in the ferrite grain exceeds 1, and the Vickers hardness is 150 HV or less.
• the cross-sectional shrinkage percentage exceeds 40% and the in-plane anisotropy
• the retention time in the first step annealing during annealing of the two-step type is short, the treatment of the carbide coarsening at the Ac1 temperature or lower is insufficient, and the thermal stability of the carbide is insufficient, and thus the carbide remaining at the second step of annealing decreases and the pearlite transformation cannot be suppressed in the structure after the slow cooling, and the cross-sectional shrinkage percentage is low. And thus the cold workability is low.
• the retention time during the first stage box annealing of the two-step type is long and the productivity is low.
• the annealing temperature during the second-step annealing during the two-step type box annealing is low, and the amount of austenite produced is small, so that the proportion of the number of carbides in the grain boundary cannot be increased.
• the cold workability is low.
• the annealing temperature during the second-step annealing during the two-step type annealing is high, the amount of the carbide remaining is decreased due to the promoted dissolution of carbides, and pearlite transformation cannot be suppressed in the structure after the slow cooling, the Vickers hardness is too high, and the cross-sectional shrinkage percentage is low.
• the cold forgeability is low.
• the annealing temperature during the second-step annealing during the two-step type annealing is low, and the amount of austenite produced is small, so that the proportion of the number of carbides in the grain boundary cannot be increased.
• the cold workability is low.
• the retention time during the second-step annealing during the two-step type annealing is long, the amount of carbides remaining is decreased due to the promoted dissolution of carbides, and pearlite transformation cannot be suppressed in the structure after the slow cooling, and the cross-sectional shrinkage percentage is low.
• the cold forgeability is low.
• the cooling rate from the second-step annealing during the two-step type annealing to 650° C. is fast, pearlite transformation occurred during cooling, the Vickers hardness is too high, and the cross-sectional shrinkage percentage is low.
• the cold workability is low.
• the X-ray diffraction intensity ratio of ⁇ 311 ⁇ ⁇ 011> is greater than 3.0.
• exceeds 0.2, and thus the cold workability is low.
• the degree of plastic anisotropy such as the in-plane anisotropy
• the steel sheet of the present invention is a steel sheet suitable as a material for automotive parts, blades, and other mechanical parts manufactured through processing steps such as punching, bending, pressing, etc. Therefore, the present invention has excellent industrial applicability.

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