EP1888800B1 - Feuille d acier laminee a froid ayant une formabilite superieure et un rapport de rendement eleve et son procede de production - Google Patents

Feuille d acier laminee a froid ayant une formabilite superieure et un rapport de rendement eleve et son procede de production Download PDF

Info

Publication number
EP1888800B1
EP1888800B1 EP06732897.1A EP06732897A EP1888800B1 EP 1888800 B1 EP1888800 B1 EP 1888800B1 EP 06732897 A EP06732897 A EP 06732897A EP 1888800 B1 EP1888800 B1 EP 1888800B1
Authority
EP
European Patent Office
Prior art keywords
steel sheet
rolled steel
less
cold rolled
precipitates
Prior art date
Legal status (The legal status 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 status listed.)
Active
Application number
EP06732897.1A
Other languages
German (de)
English (en)
Other versions
EP1888800A1 (fr
EP1888800A4 (fr
Inventor
Man-Young c/o Posco PARK
Jin-Hee c/o Posco CHUNG
Noi-Ha c/o Posco CHO
Jeong-Bong c/o POSCO YOON
Kwang-Geun c/o POSCO CHIN
Ho-Seok c/o Posco KIM
Sung-Il c/o Posco KIM
Sang-Ho c/o POSCO HAN
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Posco Holdings Inc
Original Assignee
Posco Co Ltd
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
Priority claimed from KR1020050129239A external-priority patent/KR100723181B1/ko
Application filed by Posco Co Ltd filed Critical Posco Co Ltd
Publication of EP1888800A1 publication Critical patent/EP1888800A1/fr
Publication of EP1888800A4 publication Critical patent/EP1888800A4/fr
Application granted granted Critical
Publication of EP1888800B1 publication Critical patent/EP1888800B1/fr
Active legal-status Critical Current
Anticipated expiration legal-status Critical

Links

Classifications

    • 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
    • C21D11/00—Process control or regulation for heat treatments
    • 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
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • 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
    • C22C38/002—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
    • C22C38/04—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
    • C22C38/14—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

Definitions

  • the present invention relates to niobium (Nb) and titanium (Ti)-added interstitial free (IF) cold rolled steel sheets that are used as materials for automobiles, household electronic appliances, etc. More specifically, the present invention relates to highly formable IF cold rolled steel sheets whose yield strength is enhanced due to the distribution of fine precipitates, and a process for producing the IF cold rolled steel sheets.
  • cold rolled steel sheets for use in automobiles and household electronic appliances are required to have excellent room-temperature aging resistance and bake hardenability, together with high strength and superior formability.
  • Aging is a strain aging phenomenon that arises from hardening caused by dissolved elements, such as C and N, fixed to dislocations. Since aging causes defect, called “stretcher strain", it is important to secure excellent room-temperature aging resistance.
  • Bake hardenability means increase in strength due to the presence of dissolved carbon after press formation, followed by painting and drying, by leaving a slight small amount of carbon in a solid solution state. Steel sheets with excellent bake hardenability can overcome the difficulties of press formability resulting from high strength.
  • Room-temperature aging resistance and bake hardenability can be imparted to aluminum (Al)-killed steels by batch annealing of the Al-killed steels.
  • extended time of the batch annealing causes low productivity of the Al-killed steels and severe variation in steel materials at different sites.
  • Al-killed steels have a bake hardening (BH) value (a difference in yield strength before and after painting) of 10-20 MPa, which demonstrates that an increase in yield strength is low.
  • BH bake hardening
  • interstitial free (IF) steels with excellent room-temperature aging resistance and bake hardenability have been developed by adding carbide and nitride-forming elements, such as Ti and Nb, followed by continuous annealing.
  • Japanese Unexamined Patent Publication No. Sho 57-041349 describes an enhancement in the strength of a Ti-based IF steel by adding 0.4-0.8% of manganese (Mn) and 0.04-0.12% of phosphorus (P).
  • Mn manganese
  • P phosphorus
  • Japanese Unexamined Patent Publication No. Hei 5-078784 describes an enhancement in strength by the addition of Mn as a solid solution strengthening element in an amount exceeding 0.9% and not exceeding 3.0%.
  • Korean Patent Laid-open No. 2003-0052248 describes an improvement in secondary working embrittlement resistance as well as strength and workability by the addition of 0.5-2.0% of Mn instead of P, together with aluminum (Al) and boron (B).
  • Japanese Unexamined Patent Publication No. Hei 10-158783 describes an enhancement in strength by reducing the content of P and using Mn and Si as solid solution strengthening elements.
  • Mn is used in an amount of up to 0.5%
  • Al as a deoxidizing agent is used in an amount of 0.1%
  • nitrogen (N) as an impurity is limited to 0.01% or less. If the Mn content is increased, the plating characteristics are worsened.
  • Japanese Unexamined Patent Publication No. Hei 6-057336 discloses an enhancement in the strength of an IF steel by adding 0.5-2.5% of copper (Cu) to form ⁇ -Cu precipitates. High strength of the IF steel is achieved due to the presence of the ⁇ -Cu precipitates, but the workability of the IF steel is worsened.
  • Japanese Unexamined Patent Publication Nos. Hei 9-227951 and Hei 10-265900 suggest technologies associated with improvement in workability or surface defects due to carbides by the use of Cu as a nucleus for precipitation of the carbides.
  • 0.005-0.1% of Cu is added to precipitate CuS during temper rolling of an IF steel, and the CuS precipitates are used as nuclei to form Cu-Ti-C-S precipitates during hot rolling.
  • the former publication states that the number of nuclei forming a ⁇ 111 ⁇ plane parallel to the surface of a plate increases in the vicinity of the Cu-Ti-C-S precipitates during recrystallization, which contributes to an improvement in workability.
  • Japanese Unexamined Patent Publication Nos. Hei 6-240365 and Hei 7-216340 describe the addition of a combination of Cu and P to improve the corrosion resistance of baking hardening type IF steels.
  • Cu is added in an amount of 0.05-1.0% to ensure improved corrosion resistance.
  • Cu is added in an excessively large amount of 0.2% or more.
  • Japanese Unexamined Patent Publication Nos. Hei 10-280048 and Hei 10-287954 suggest the dissolution of carbosulfide (Ti-C-S based) in a carbide at the time of reheating and annealing to obtain a solid solution in crystal grain boundaries, thereby achieving a bake hardening (BH) value (a difference in yield strength before and after baking) of 30 MPa or more.
  • BH bake hardening
  • EP-A-1136575 discloses a method for manufacturing cold-rolled steel sheet which comprises the steps of: rough-rolling a slab using a rough-rolling unit; finish-rolling the sheet bar using a continuous hot finishing-rolling mill; cooling the hot-rolled steel strip on a runout table; coiling thus cooled hot-rolled steel strip; and applying picking, cold-rolling the hot-rolled steel strip, and final annealing to the cold-rolled steel strip.
  • the cold rolled steel sheets of the present invention have characteristics of soft cold rolled steel sheets of the order of 280 MPa and high-strength cold rolled steel sheets of the order of 340 MPa or more.
  • soft cold rolled steel sheets of the order of 280 MPa are produced.
  • the soft cold rolled steel sheets further contain at least one solid solution strengthening element selected from Si and Cr, or the P content is in the range of 0.015-0.2%, a high strength of 340 MPa or more is attained.
  • the P content in the high-strength steels containing P alone is preferably in the range of 0.03% to 0.2%.
  • the Si content in the high-strength steels is preferably in the range of 0.1 to 0.8%.
  • the Cr content in the high-strength steels is preferably in the range of 0.2 to 1.2.
  • the P content may be freely designed in an amount of 0.2% or less.
  • the cold rolled steel sheets of the present invention may further contain 0.01-0.2 wt% of Mo.
  • Fine precipitates having a size of 0.2 ⁇ m or less are distributed in the cold rolled steel sheets of the present invention.
  • examples of such precipitates include MnS precipitates, CuS precipitates, and composite precipitates of MnS and CuS. These precipitates are referred to simply as "(Mn,Cu)S".
  • the present inventors have found that when fine precipitates are distributed in Nb and Ti-added IF steels (also referred to simply as "Nb-Ti composite IF steels"), the yield strength of the IF steels is enhanced and the in-plane anisotropy index of the IF steels is lowered, thus leading to an improvement in workability.
  • the present invention has been achieved based on this finding.
  • the precipitates used in the present invention have drawn little attention in conventional IF steels. Particularly, the precipitates have not been actively used from the viewpoint of yield strength and in-plane anisotropy index.
  • Nb-Ti composite IF steels Regulation of the components in the Nb-Ti composite IF steels is required to obtain (Mn,Cu)S precipitates and/or AlN precipitates. If the IF steels contain Ti, Zr and other elements, S and N preferentially react with Ti and Zr. Since the cold rolled steel sheets of the present invention are Nb-Ti composite IF steels, Ti reacts with C, N and S. Accordingly, it is necessary to regulate the components so that S and N are precipitated into (Mn,Cu)S and AlN forms, respectively.
  • the fine precipitates thus obtained allow the formation of minute crystal grains. Minuteness in the size of crystal grains relatively increases the proportion of crystal grain boundaries. Accordingly, the dissolved carbon is present in a larger amount in the crystal grain boundaries than within the crystal grains, thus achieving excellent room-temperature non-aging properties. Since the dissolved carbon present within the crystal grains can more freely migrate, it binds to movable dislocations, thus affecting the room-temperature aging properties. In contrast, the dissolved carbon segregated in stable positions, such as in the crystal grain boundaries and in the vicinity of the precipitates, is activated at a high temperature, for example, a temperature for painting/baking treatment, thus affecting the bake hardenability.
  • the fine precipitates distributed in the steel sheets of the present invention have a positive influence on the increase of yield strength arising from precipitation enhancement, improvement in strength-ductility balance, in-plane anisotropy index, and plasticity anisotropy.
  • the fine (Mn,Cu)S precipitates and AlN precipitates must be uniformly distributed. According to the cold rolled steel sheets of the present invention, contents of components affecting the precipitation, composition between the components, production conditions, and particularly cooling rate after hot rolling, have a great influence on the distribution of the fine precipitates.
  • the content of carbon (C) is preferably limited to 0.01% or less.
  • Carbon (C) affects the room-temperature aging resistance and bake hardenability of the cold rolled steel sheets.
  • the carbon content exceeds 0.01%, the addition of the expensive agents Nb and Ti is required to remove the remaining carbon, which is economically disadvantageous and is undesirable in terms of formability.
  • the carbon is preferably added in an amount of 0.001% or more, and more preferably 0.005% to 0.01%.
  • the carbon content is less than 0.005%, room-temperature aging resistance can be ensured without increasing the amounts of Nb and Ti.
  • the content of copper (Cu) is preferably in the range of 0.01-0.2%.
  • Copper serves to form fine CuS precipitates, which make the crystal grains fine. Copper lowers the in-plane anisotropy index of the cold rolled steel sheets and enhances the yield strength of the cold rolled steel sheets by precipitation promotion.
  • the Cu content In order to form fine precipitates, the Cu content must be 0.01% or more. When the Cu content is more than 0.2%, coarse precipitates are obtained. The Cu content is more preferably in the range of 0.03 to 0.2%.
  • the content of manganese (Mn) is preferably in the range of 0.01-0.3%.
  • Manganese serves to precipitate sulfur in a solid solution state in the steels as MnS precipitates, thereby preventing occurrence of hot shortness caused by the dissolved sulfur, or is known as a solid solution strengthening element. From such a technical standpoint, manganese is generally added in a large amount. The present inventors have found that when the manganese content is reduced and the sulfur content is optimized, very fine MnS precipitates are obtained. Based on this finding, the manganese content is limited to 0.3% or less. In order to ensure this characteristic, the manganese content must be 0.01% or more. When the manganese content is less than 0.01%, i.e. the sulfur content remaining in a solid solution state is high, hot shortness may occur. When the manganese content is greater than 0.3%, coarse MnS precipitates are formed, thus making it difficult to achieve desired strength. A more preferable Mn content is within the range of 0.01 to 0.12%.
  • the content of sulfur (S) is preferably limited to 0.08% or less.
  • S Sulfur
  • Cu and/or MnS precipitates reacts with Cu and/or Mn to form CuS and MnS precipitates, respectively.
  • sulfur content is greater than 0.08%, the proportion of dissolved sulfur is increased. This increase of dissolved sulfur greatly deteriorates the ductility and formability of the steel sheets and increases the risk of hot shortness.
  • a sulfur content of 0.005% or more is preferred.
  • the content of aluminum (Al) is preferably limited to 0.1% or less.
  • Aluminum reacts with nitrogen (N) to form fine AlN precipitates, thereby completely preventing aging by dissolved nitrogen.
  • N nitrogen
  • AlN precipitates are sufficiently formed.
  • the distribution of the fine AlN precipitates in the steel sheets allows the formation of minute crystal grains and enhances the yield strength of the steel sheets by precipitation enhancement.
  • a more preferable Al content is in the range of 0.01 to 0.1%.
  • the content of nitrogen (N) is preferably limited to 0.02% or less.
  • nitrogen is added in an amount of up to 0.02%. Otherwise, the nitrogen content is controlled to 0.004% or less. When the nitrogen content is less than 0.004%, the number of the AlN precipitates is small, and therefore, the minuteness effects of crystal grains and the precipitation enhancement effects are negligible. In contrast, when the nitrogen content is greater than 0.02%, it is difficult to guarantee aging properties by use of dissolved nitrogen.
  • the content of phosphorus (P) is preferably limited to 0.2% or less.
  • Phosphorus is an element that has excellent solid solution strengthening effects while allowing a slight reduction in r-value. Phosphorus guarantees high strength of the steel sheets of the present invention in which the precipitates are controlled. It is desirable that the phosphorus content in steels requiring a strength of the order of 280 MPa be defined to 0.015% or less. It is desirable that the phosphorus content in high-strength steels of the order of 340 MPa be limited to a range exceeding 0.015% and not exceeding 0.2%. A phosphorus content exceeding 0.2% can lead to a reduction in ductility of the steel sheets. Accordingly, the phosphorus content is preferably limited to a maximum of 0.2%. When Si and Cr are added in the present invention, the phosphorus content can be appropriately controlled to be 0.2% or less to achieve the desired strength.
  • the content of boron (B) is preferably in the range of 0.0001 to 0.002%.
  • boron is added to prevent occurrence of secondary working embrittlement.
  • a preferable boron content is 0.0001% or more. When the boron content exceeds 0.002%, the deep drawability of the steel sheets may be markedly deteriorated.
  • the content of niobium (Nb) is preferably in the range of 0.002 to 0.04%.
  • Nb is added for the purpose of ensuring the non-aging properties and improving the formability of the steel sheets.
  • Nb which is a potent carbide-forming element, is added to steels to form NbC precipitates in the steels.
  • the NbC precipitates permit the steel sheets to be well textured during annealing, thus greatly improving the deep drawability of the steel sheets.
  • the content of Nb added is not greater than 0.002%, the NbC precipitates are obtained in very small amounts. Accordingly, the steel sheets are not well textured and thus there is little improvement in the deep drawability of the steel sheets.
  • the Nb content exceeds 0.04%, the NbC precipitates are obtained in very large amounts. Accordingly, the deep drawability and elongation of the steel sheets are lowered, and thus the formability of the steel sheets may be markedly deteriorated.
  • the content of titanium (Ti) is preferably in the range of 0.005 to 0.15%.
  • Titanium is added for the purpose of ensuring the non-aging properties and improving the formability of the steel sheets.
  • Ti which is a potent carbide-forming element, is added to steels to form TiC precipitates in the steels.
  • the TiC precipitates allow the precipitation of dissolved carbon to ensure non-aging properties.
  • the content of Ti added is less than 0.005%, the TiC precipitates are obtained in very small amounts. Accordingly, the steel sheets are not well textured and thus there is little improvement in the deep drawability of the steel sheets.
  • the titanium is added in an amount exceeding 0.15%, very large TiC precipitates are formed. Accordingly, minuteness effects of crystal grains are reduced, resulting in high in-plane anisotropy index, reduction of yield strength and marked worsening of plating characteristics.
  • S* which is determined by Relationship 2, represents the content of sulfur that does not react with Ti and thereafter reacts with Cu.
  • the value of (Cu/63.5)/(S*/32) be equal to or greater than 1. If the value of (Cu/63.5)/(S*/32) is greater than 30, coarse CuS precipitates are distributed, which is undesirable.
  • the value of (Cu/63.5)/(S*/32) is preferably in the range of 1 to 20, more preferably 1 to 9, and most preferably 1 to 6. 1 ⁇ Mn / 55 + Cu / 63.5 / S * / 32 ⁇ 30
  • Relationship 3 is associated with the formation of (Mn,Cu)S precipitates, and is obtained by adding a Mn content to Relationship 1.
  • the value of (Mn/55 + Cu/63.5)/(S*/32) must be 1 or greater.
  • the value of Relationship 3 is greater than 30, coarse (Mn,Cu)S precipitates are obtained.
  • a more preferable value of (Cu/63.5)/(S*/32) is preferably in the range of 1 to 20, more preferably 1 to 9, and most preferably 1 to 6.
  • the sum of Mn and Cu is more preferably 0.05-0.4%.
  • Relationship 4 is associated with the formation of fine (Mn,Cu)S precipitates.
  • N* which is determined by Relationship 5
  • the value of (Al/27)/(N*/14) be in the range of 1-10.
  • the value of (Al/27)/(N*/14) must be 1 or greater. If the value of (Al/27)/(N*/14) is greater than 10, coarse AlN precipitates are obtained and thus poor workability and low yield strength are caused. It is preferred that the value of (Al/27)/(N*/14) be in the range of 1 to 6.
  • one or more kinds selected from the group consisting of 0.01-0.2% of Cu, 0.01-0.3% of Mn and 0.004-0.2% of N lead to various combinations of (Mn,Cu)S and AlN precipitates having a size not greater than 0.2 ⁇ m.
  • Relationship 6 is associated with the formation of NbC and TiC precipitates to remove the carbon in a solid solution state, thereby achieving room-temperature non-aging properties.
  • Ti* which is determined by Relationship 7, represents the content of titanium that reacts with N and S and thereafter reacts with C.
  • Relationship 8 is associated with the achievement of bake hardenability.
  • Cs which is expressed in ppm by Relationship 8, represents the content of dissolved carbon that is not precipitated into NbC and TiC forms.
  • the Cs value In order to achieve a high bake hardening value, the Cs value must be 5 ppm or more. If the Cs value exceeds 30 ppm, the content of dissolved carbon is increased, making it difficult to attain room-temperature non-aging properties.
  • the fine precipitates are uniformly distributed in the compositions of the present invention. It is preferable that the precipitates have an average size of 0.2 ⁇ m or less. According to a study conducted by the present inventors, when the precipitates have an average size greater than 0.2 ⁇ m, the steel sheets have poor strength and low in-plane anisotropy index. Further, large amounts of precipitates having a size of 0.2 ⁇ m or less are distributed in the compositions of the present invention. While the number of the distributed precipitates is not particularly limited, it is more advantageous with higher number of the precipitates.
  • the number of the distributed precipitates is preferably 1 x 10 5 /mm 2 or more, more preferably 1 x 10 6 /mm 2 or more, and most preferably 1 x 10 7 /mm 2 or more.
  • the plasticity-anisotropy index is increased and the in-plane anisotropy index is lowered with increasing number of the precipitates, and as a result, the workability is greatly improved. It is commonly known that there is a limitation in increasing the workability because the in-plane anisotropy index is increased with increasing plasticity-anisotropy index.
  • the plasticity-anisotropy index of the steel sheets is increased and the in-plane anisotropy index of the steel sheets is lowered.
  • the steel sheets of the present invention in which the fine precipitates are formed satisfy a yield ratio (yield strength/tensile strength) of 0.58 or higher.
  • the steel sheets of the present invention When the steel sheets of the present invention are applied to high-strength steel sheets, they may further contain at least one solid solution strengthening element selected from P, Si and Cr.
  • P solid solution strengthening element
  • Si solid solution strengthening element
  • the content of silicon (Si) is preferably in the range of 0.1 to 0.8%.
  • Si is an element that has solid solution strengthening effects and shows a slight reduction in elongation. Si guarantees high strength of the steel sheets of the present invention in which the precipitates are controlled. Only when the Si content is 0.1% or more, high strength can be ensured. However, when the Si content is more than 0.8%, the ductility of the steel sheets is deteriorated.
  • the content of chromium (Cr) is preferably in the range of 0.2 to 1.2%.
  • Cr is an element that has solid solution strengthening effects, lowers the secondary working embrittlement temperature, and lowers the aging index due to the formation of Cr carbides. Cr guarantees high strength of the steel sheets of the present invention in which the precipitates are controlled and serves to lower the in-plane anisotropy index of the steel sheets. Only when the Cr content is 0.2% or more, high strength can be ensured. However, when the Cr content exceeds 1.2%, the ductility of the steel sheets is deteriorated.
  • the cold rolled steel sheets of the present invention may further contain molybdenum (Mo).
  • the content of molybdenum (Mo) in the cold rolled steel sheets of the present invention is preferably in the range of 0.01 to 0.2%.
  • Mo is added as an element that increases the plasticity-anisotropy index of the steel sheets. Only when the molybdenum content is not lower than 0.01%, the plasticity-anisotropy index of the steel sheets is increased. However, when the molybdenum content exceeds 0.2%, the plasticity-anisotropy index is not further increased and there is a danger of hot shortness.
  • the process of the present invention is characterized in that a steel satisfying one of the steel compositions defined above is processed through hot rolling and cold rolling to form precipitates having an average size of 0.2 ⁇ m or less in a cold rolled sheet.
  • the average size of the precipitates in the cold rolled plate is affected by the design of the steel composition and the processing conditions, such as reheating temperature and winding temperature. Particularly, cooling rate after hot rolling has a direct influence on the average size of the precipitates.
  • a steel satisfying one of the compositions defined above is reheated, and is then subjected to hot rolling.
  • the reheating temperature is preferably 1,100°C or higher.
  • the hot rolling is performed at a finish rolling temperature not lower than the Ar 3 transformation point.
  • finish rolling temperature is lower than the Ar 3 transformation point, rolled grains are created, which deteriorates the workability and causes poor strength.
  • the cooling is preferably performed at a rate of 300 °C/min or higher before winding and after hot rolling.
  • the composition of the components is controlled to obtain fine precipitates, the precipitates may have an average size greater than 0.2 ⁇ m at a cooling rate of less than 300 °C/min. That is, as the cooling rate is increased, many nuclei are created and thus the size of the precipitates becomes finer and finer. Since the size of the precipitates is decreased with increasing cooling rate, it is not necessary to define the upper limit of the cooling rate.
  • the cooling rate is preferably in the range of 300-1000 °C/min.
  • winding is performed at a temperature not higher than 700°C.
  • the winding temperature is higher than 700°C, the precipitates are grown too coarsely, thus making it difficult to ensure high strength.
  • the steel is cold rolled at a reduction rate of 50-90%. Since a cold reduction rate lower than 50 % leads to creation of a small amount of nuclei upon annealing recrystallization, the crystal grains are grown excessively upon annealing, thereby coarsening of the crystal grains recrystallized through annealing, which results in reduction of the strength and formability. A cold reduction rate higher than 90 % leads to enhanced formability, while creating an excessively large amount of nuclei, so that the crystal grains recrystallized through annealing become too fine, thus deteriorating the ductility of the steel.
  • Continuous annealing temperature plays an important role in determining the mechanical properties of the final product.
  • the continuous annealing is preferably performed at a temperature of 700 to 900°C.
  • the continuous annealing is performed at a temperature lower than 700°C, the recrystallization is not completed and thus a desired ductility cannot be ensured.
  • the continuous annealing is performed at a temperature higher than 900°C, the recrystallized grains become coarse and thus the strength of the steel is deteriorated.
  • the continuous annealing is maintained until the steel is completely recrystallized.
  • the recrystallization of the steel can be completed for about 10 seconds or more.
  • the continuous annealing is preferably performed for 10 seconds to 30 minutes.
  • the mechanical properties of steel sheets produced in the following examples were evaluated according to the ASTM E-8 standard test methods. Specifically, each of the steel sheets was machined to obtain standard samples. The yield strength, tensile strength, elongation, plasticity-anisotropy index (r m value) and in-plane anisotropy index ( ⁇ r value), and the aging index were measured using a tensile strength tester (available from INSTRON Company, Model 6025).
  • the aging index of the steel sheets is defined as a yield point elongation measured by annealing each of the samples, followed by 1.0% skin pass rolling and thermally processing at 100°C for 2 hours.
  • the bake hardening (BH) value of the standard samples was measured by the following procedure. After a 2% strain was applied to each of the samples, the strained sample was annealed at 170°C for 20 minutes. The yield strength of the annealed sample was measured. The BH value was calculated by subtracting the yield strength measured before annealing from the yield strength value measured after annealing.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • steel slabs were prepared in accordance with the compositions shown in the following tables.
  • the steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets.
  • the hot rolled steel sheets were cooled at a rate of 400 °C/min., wound at 650°C, cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets.
  • the finish hot rolling was performed at 910°C, which is above the Ar 3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10 °C/second to 830°C for 40 seconds to produce the final cold rolled steel sheets.
  • the distribution of fine precipitates in Nb-Ti composite IF steels allows the formation of minute crystal grains, and as a result, the in-plane anisotropy index is lowered and the yield strength is enhanced by precipitation enhancement.

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Heat Treatment Of Sheet Steel (AREA)

Claims (32)

  1. Feuille d'acier laminée à froid ayant une formabilité supérieure et un rapport d'élasticité élevé, la feuille d'acier laminée à froid ayant une composition comprenant 0,01 % de C ou moins, 0,01 à 0,2 % de Cu, 0,005 à 0,08 % de S, 0,1 % de Al ou moins, 0,004 % de N ou moins, 0,2 % de P ou moins, 0,0001 à 0,002 % de B, 0,002 à 0,04 % de Nb, 0,005 à 0,15 % de Ti, facultativement au moins un parmi 0,1 à 0,8 % de Si et 0,2 à 1,2 % de Cr, facultativement 0,01 à 0,2 % de Mo, en poids, et le reste étant Fe et autres impuretés inévitables,
    la composition satisfaisant les relations suivantes : 1 ≤ (Cu/63,5)/(S*/32) ≤ 30 et S* = S-0,8 x (Ti-0,8 x (48/14) x N) x (32/48), et
    la feuille d'acier comprenant des précipités de CuS ayant une taille moyenne de 0,2 µm ou moins.
  2. Feuille d'acier laminée à froid ayant une formabilité supérieure et un rapport d'élasticité élevé, la feuille d'acier laminée à froid ayant une composition comprenant 0,01 % de C ou moins, 0,01 à 0,2 % de Cu, 0,01 à 0,3 % de Mn, 0,005 à 0,08 % de S, 0,1 % de Al ou moins, 0,004 % de N ou moins, 0,2 % de P ou moins, 0,0001 à 0,002 % de B, 0,002 à 0,04 % de Nb, 0,005 à 0,15 % de Ti, en poids, et le reste étant Fe et autres impuretés inévitables,
    la composition satisfaisant les relations suivantes : 1 ≤ (Mn/55 + Cu/63,5)/(S*/32) ≤ 30 et S* = S-0,8 x (Ti-0,8 x (48/14) x N) x (32/48), et la feuille d'acier comprenant des précipités de (Mn,Cu)S ayant une taille moyenne de 0,2 µm ou moins.
  3. Feuille d'acier laminée à froid ayant une formabilité supérieure et un rapport d'élasticité élevé, la feuille d'acier laminée à froid ayant une composition comprenant 0,01 % de C ou moins, 0,01 à 0,2 % de Cu, 0,005 à 0,08 % de S, 0,1 % de Al ou moins, 0,004 à 0,02 % de N, 0,2 % de P ou moins, 0,0001 à 0,002 % de B, 0,002 à 0,04 % de Nb, 0,005 à 0,15 % de Ti, en poids, et le reste étant Fe et autres impuretés inévitables,
    la composition satisfaisant les relations suivantes : 1 ≤ (Cu/63,5)/(S*/32) ≤ 30, 1 ≤ (Al/27)/(N*/14) ≤ 10, S* = S-0,8 x (Ti-0,8 x (48/14) x N) x (32/48) et N* = N-0,8 x (Ti-0,8 x (48/32) x S) x (14/48), et la feuille d'acier comprenant des précipités de CuS et des précipités de AlN ayant une taille moyenne de 0,2 µm ou moins.
  4. Feuille d'acier laminée à froid ayant une formabilité supérieure et un rapport d'élasticité élevé, la feuille d'acier laminée à froid ayant une composition comprenant 0,01 % de C ou moins, 0,01 à 0,2 % de Cu, 0,01 à 0,3 % de Mn, 0,005 à 0,08 % de S, 0,1 % de Al ou moins, 0,004 à 0,02 % de N, 0,2 % de P ou moins, 0,0001 à 0,002 % de B, 0,002 à 0,04 % de Nb, 0,005 à 0,15 % de Ti, en poids, et le reste étant Fe et autres impuretés inévitables,
    la composition satisfaisant les relations suivantes : 1 ≤ (Mn/55 + Cu/63,5)/(S*/32) ≤ 30, 1 ≤ (Al/27)/(N*/14) ≤ 10, S* = S-0,8 x (Ti-0,8 x (48/14) x N) x (32/48) et N* = N-0,8 x (Ti-0,8 x (48/32) x S) x (14/48), et la feuille d'acier comprenant des précipités de (Mn,Cu)S et des précipités de AlN ayant une taille moyenne de 0,2 µm ou moins.
  5. Feuille d'acier laminée à froid selon la revendication 1, dans laquelle les teneurs en C, Ti, Nb et S satisfont les relations suivantes : 0,8 ≤ (Ti*/48 + Nb/93)/(C/12) ≤ 5,0 et Ti* = Ti-0,8 x ((48/14) x N + (48/32) x S).
  6. Feuille d'acier laminée à froid selon la revendication 5, dans laquelle la teneur en C est de 0,005 % ou moins.
  7. Feuille d'acier laminée à froid selon la revendication 1, dans laquelle du soluté de carbone (Cs) [Cs = (C - Nb x 12/93 - Ti* x 12/48) x 1000 où Ti* = Ti-0,8 x (((48/14) x N + (48/32) x S), à condition que, lorsque Ti* est inférieur à 0, Ti* soit défini comme étant 0], qui est déterminé par les teneurs en C et Ti, est compris entre 5 et 30.
  8. Feuille d'acier laminée à froid selon la revendication 7, dans laquelle la teneur en C est comprise entre 0,001 et 0,01 %.
  9. Feuille d'acier laminée à froid selon l'une quelconque des revendications 1 à 4, dans laquelle la feuille d'acier laminée à froid satisfait un rapport d'élasticité (limite d'élasticité/résistance à la traction) de 0,58 ou plus.
  10. Feuille d'acier laminée à froid selon l'une quelconque des revendications 1 à 4, dans laquelle le nombre des précipités est 1 x 106/mm2 ou plus.
  11. Feuille d'acier laminée à froid selon la revendication 1, dans laquelle la teneur en P est de 0,015 % ou moins.
  12. Feuille d'acier laminée à froid selon la revendication 1, dans laquelle la teneur en P est comprise entre 0,03 % et 0,2 %.
  13. Feuille d'acier laminée à froid selon la revendication 2 ou 4, dans laquelle la somme de Mn et Cu est comprise entre 0,05 % et 0,4 %.
  14. Feuille d'acier laminée à froid selon la revendication 2 ou 4, dans laquelle la teneur en Mn est comprise entre 0,01 et 0,12 %.
  15. Feuille d'acier laminée à froid selon la revendication 2 ou 4, dans laquelle la valeur de (Mn/55 + Cu/63,5)/(S*/32) est comprise dans la plage allant de 1 à 9.
  16. Feuille d'acier laminée à froid selon la revendication 3 ou 4, dans laquelle la valeur de (Al/27)/(N*/14) est comprise dans la plage allant de 1 à 6.
  17. Procédé de production d'une feuille d'acier laminée à froid ayant une formabilité supérieure et un rapport d'élasticité élevé, le procédé comprenant les étapes :
    réchauffer une plaque à une température de 1100 °C ou plus, la plaque ayant une composition comprenant 0,01 % de C ou moins, 0,01 à 0,2 % de Cu, 0,005 à 0,08 % de S, 0,1 % de Al ou moins, 0,004 % de N ou moins, 0,2 % de P ou moins, 0,0001 à 0,002 % de B, 0,002 à 0,04 % de Nb, 0,005 à 0,15 % de Ti, facultativement au moins un parmi 0,1 à 0,8 % de Si et 0,2 à 1,2 % de Cr, facultativement 0,01 à 0,2 % de Mo, en poids, et le reste étant Fe et autres impuretés inévitables, et la composition satisfaisant les relations suivantes : 1 ≤ (Cu/63,5)/(S*/32) ≤ 30 et S* = S-0,8 x (Ti-0,8 x (48/14) x N) x (32/48) ;
    laminer à chaud la plaque réchauffée à une température de laminage de finition du point de transformation Ar3 ou plus, pour fournir une feuille d'acier laminée à chaud ;
    refroidir la feuille d'acier laminée à chaud à une vitesse de 300°C/min ou plus ;
    enrouler la feuille d'acier refroidie à 700°C ou moins ;
    laminer à froid la feuille d'acier enroulée ; et
    recuire de manière continue la feuille d'acier laminée à froid, la feuille d'acier laminée à froid comprenant des précipités de CuS ayant une taille moyenne de 0,2 µm ou moins.
  18. Procédé de production d'une feuille d'acier laminée à froid ayant une formabilité supérieure et un rapport d'élasticité élevé, le procédé comprenant les étapes :
    réchauffer une plaque à une température de 1100°C ou plus, la plaque ayant une composition comprenant 0,01 % de C ou moins, 0,01 à 0,2 % de Cu, 0,01 à 0,3 % de Mn, 0,005 à 0,08 % de S, 0,1 % de Al ou moins, 0,004 % de N ou moins, 0,2 % de P ou moins, 0,0001 à 0,002 % de B, 0,002 à 0,04 % de Nb, 0,005 à 0,15 % de Ti, en poids, et le reste étant Fe et autres impuretés inévitables, et la composition satisfaisant les relations suivantes : 1 ≤ (Mn/55 + Cu/63,5)/(S*/32) ≤ 30 et S* = S-0,8 x (Ti-0,8 x (48/14) x N) x (32/48) ;
    laminer à chaud la plaque réchauffée à une température de laminage de finition du point de transformation Ar3 ou plus, pour fournir une feuille d'acier laminée à chaud ;
    refroidir la feuille d'acier laminée à chaud à une vitesse de 300°C/min ou plus ;
    enrouler la feuille d'acier refroidie à 700°C ou moins ;
    laminer à froid la feuille d'acier enroulée ; et
    recuire de manière continue la feuille d'acier laminée à froid, la feuille d'acier laminée à froid comprenant des précipités de (Mn,Cu)S ayant une taille moyenne de 0,2 µm ou moins.
  19. Procédé de production d'une feuille d'acier laminée à froid ayant une formabilité supérieure et un rapport d'élasticité élevé, le procédé comprenant les étapes :
    réchauffer une plaque à une température de 1100°C ou plus, la plaque ayant une composition comprenant 0,01 % de C ou moins, 0,01 à 0,2 % de Cu, 0,005 à 0,08 % de S, 0,1 % de Al ou moins, 0,004 à 0,02 % de N, 0,2 % de P ou moins, 0,0001 à 0,002 % de B, 0,002 à 0,04 % de Nb, 0,005 à 0,15 % de Ti, en poids, et le reste étant Fe et autres impuretés inévitables, et la composition satisfaisant les relations suivantes : 1 ≤ (Cu/63,5)/(S*/32) ≤ 30, 1 ≤ (Al/27)/(N*/14) ≤ 10, S* = S-0,8 x (Ti-0,8 x (48/14) x N) x (32/48) et N* = N-0,8 x (Ti-0,8 x (48/32) x S) x (14/48) ;
    laminer à chaud la plaque réchauffée à une température de laminage de finition du point de transformation Ar3 ou plus, pour fournir une feuille d'acier laminée à chaud ;
    refroidir la feuille d'acier laminée à chaud à une vitesse de 300°C/min ou plus ;
    enrouler la feuille d'acier refroidie à 700°C ou moins ;
    laminer à froid la feuille d'acier enroulée ; et
    recuire de manière continue la feuille d'acier laminée à froid, la feuille d'acier laminée à froid comprenant des précipités de CuS et des précipités de AlN ayant une taille moyenne de 0,2 µm ou moins.
  20. Procédé de production d'une feuille d'acier laminée à froid ayant une formabilité supérieure et un rapport d'élasticité élevé, le procédé comprenant les étapes :
    réchauffer une plaque à une température de 1100°C ou plus, la plaque ayant une composition comprenant 0,01 % de C ou moins, 0,01 à 0,2 % de Cu, 0,01 à 0,3 % de Mn, 0,005 à 0,08 % de S, 0,1 % de Al ou moins, 0,004 à 0,02 % de N, 0,2 % de P ou moins, 0,0001 à 0,002 % de B, 0,002 à 0,04 % de Nb, 0,005 à 0,15 % de Ti, en poids, et le reste étant Fe et autres impuretés inévitables, et la composition satisfaisant les relations suivantes : 1 ≤ (Mn/55 + Cu/63,5)/(S*/32) ≤ 30, 1 ≤ (Al/27)/(N*/14) ≤ 10, S* = S-0,8 x (Ti-0,8 x (48/14) x N) x (32/48) et N* = N-0,8 x (Ti-0,8 x (48/32) x S) x (14/48) ;
    laminer à chaud la plaque réchauffée à une température de laminage de finition du point de transformation Ar3 ou plus, pour fournir une feuille d'acier laminée à chaud ;
    refroidir la feuille d'acier laminée à chaud à une vitesse de 300°C/min ou plus ;
    enrouler la feuille d'acier refroidie à 700°C ou moins ;
    laminer à froid la feuille d'acier enroulée ; et
    recuire de manière continue la feuille d'acier laminée à froid, la feuille d'acier laminée à froid comprenant des précipités de (Mn,Cu)S et des précipités de AlN ayant une taille moyenne de 0,2 µm ou moins.
  21. Procédé selon la revendication 17, dans lequel les teneurs en C, Ti, Nb, N et S satisfont les relations suivantes : 0,8 ≤ (Ti*/48 + Nb/93)/(C/12) ≤ 5,0 et Ti* = Ti-0,8 x ((48/14) x N + (48/32) x S).
  22. Procédé selon la revendication 21, dans lequel la teneur en C est de 0,005 % ou moins.
  23. Procédé selon la revendication 17, dans lequel du soluté de carbone (Cs) [Cs = (C - Nb x 12/93 - Ti* x 12/48) x 1000 où Ti* = Ti-0,8 x (((48/14) x N + (48/32) x S), à condition que, lorsque Ti* est inférieur à 0, Ti* soit défini comme étant 0], qui est déterminé par les teneurs en C et Ti, est compris entre 5 et 30.
  24. Procédé selon la revendication 23, dans lequel la teneur en C est comprise entre 0,001 et 0,01 %.
  25. Procédé selon l'une quelconque des revendications 17 à 20, dans lequel la feuille d'acier laminée à froid satisfait un rapport d'élasticité (limite d'élasticité/résistance à la traction) de 0,58 ou plus.
  26. Procédé selon l'une quelconque des revendications 17 à 20, dans lequel le nombre des précipités est 1 x 106/mm2 ou plus.
  27. Procédé selon la revendication 17, dans lequel la teneur en P est de 0,015 % ou moins.
  28. Procédé selon la revendication 17, dans lequel la teneur en P est comprise entre 0,03 % et 0,2 %.
  29. Procédé selon la revendication 18 ou 20, dans lequel la somme de Mn et Cu est comprise entre 0,08 % et 0,4 %.
  30. Procédé selon la revendication 18 ou 20, dans lequel la teneur en Mn est comprise entre 0,01 et 0,12 %.
  31. Procédé selon la revendication 18 ou 20, dans lequel la valeur de (Mn/55 + Cu/63,5)/(S*/32) est comprise dans la plage allant de 1 à 9.
  32. Procédé selon la revendication 19 ou 20, dans lequel la valeur de (Al/27)/(N*/14) est comprise dans la plage allant de 1 à 6.
EP06732897.1A 2005-05-03 2006-05-03 Feuille d acier laminee a froid ayant une formabilite superieure et un rapport de rendement eleve et son procede de production Active EP1888800B1 (fr)

Applications Claiming Priority (6)

Application Number Priority Date Filing Date Title
KR20050037183 2005-05-03
KR1020050129239A KR100723181B1 (ko) 2005-05-03 2005-12-26 성형성이 우수한 냉연강판과 그 제조방법
KR1020050129237A KR100723163B1 (ko) 2005-05-03 2005-12-26 면내이방성이 우수한 냉연강판과 그 제조방법
KR1020050129236A KR100723164B1 (ko) 2005-05-03 2005-12-26 가공성이 우수한 냉연강판과 그 제조방법
KR1020050129235A KR100723165B1 (ko) 2005-05-03 2005-12-26 소성이방성이 우수한 냉연강판과 그 제조방법
PCT/KR2006/001670 WO2006118425A1 (fr) 2005-05-03 2006-05-03 Feuille d’acier laminee a froid ayant une formabilite superieure et un rapport de rendement eleve et son procede de production

Publications (3)

Publication Number Publication Date
EP1888800A1 EP1888800A1 (fr) 2008-02-20
EP1888800A4 EP1888800A4 (fr) 2011-01-26
EP1888800B1 true EP1888800B1 (fr) 2018-11-07

Family

ID=37308183

Family Applications (1)

Application Number Title Priority Date Filing Date
EP06732897.1A Active EP1888800B1 (fr) 2005-05-03 2006-05-03 Feuille d acier laminee a froid ayant une formabilite superieure et un rapport de rendement eleve et son procede de production

Country Status (2)

Country Link
EP (1) EP1888800B1 (fr)
WO (1) WO2006118425A1 (fr)

Families Citing this family (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
KR20070038730A (ko) 2005-10-06 2007-04-11 주식회사 포스코 항복비가 우수한 석출강화형 냉연강판 및 그 제조방법
CN114231845A (zh) * 2021-12-03 2022-03-25 本钢板材股份有限公司 一种屈服强度220Mpa级冷轧深冲高强钢及其生产方法

Family Cites Families (13)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS5967322A (ja) * 1982-10-08 1984-04-17 Kawasaki Steel Corp 深絞り用冷延鋼板の製造方法
JPH0747797B2 (ja) * 1989-03-10 1995-05-24 川崎製鉄株式会社 耐つまとび性、耐泡・黒点欠陥性及びプレス成形性に優れたほうろう用鋼板並びにその製造方法
US5200005A (en) * 1991-02-08 1993-04-06 Mcgill University Interstitial free steels and method thereof
TW415967B (en) * 1996-02-29 2000-12-21 Kawasaki Steel Co Steel, steel sheet having excellent workability and method of the same by electric furnace-vacuum degassing process
DE19628714C1 (de) * 1996-07-08 1997-12-04 Mannesmann Ag Verfahren zur Herstellung von Präzisionsstahlrohren
JP3978807B2 (ja) 1997-04-09 2007-09-19 Ｊｆｅスチール株式会社 耐歪時効性に優れた塗装焼付硬化型冷延鋼板の製造方法
JPH10287954A (ja) 1997-04-15 1998-10-27 Kawasaki Steel Corp 耐歪時効特性に優れた塗装焼付硬化型冷延鋼板およびその製造方法
WO2001012864A1 (fr) * 1999-08-10 2001-02-22 Nkk Corporation Procede de production de feuillards d'acier lamines a froid
CN1190513C (zh) * 2000-06-20 2005-02-23 杰富意钢铁株式会社 薄钢板及其制造方法
JP3958921B2 (ja) * 2000-08-04 2007-08-15 新日本製鐵株式会社 塗装焼付硬化性能と耐常温時効性に優れた冷延鋼板及びその製造方法
JP2003041342A (ja) * 2002-05-29 2003-02-13 Nkk Corp 打ち抜き性に優れる冷延鋼板
EP1689901B1 (fr) * 2003-11-10 2018-03-21 Posco Tole d'acier laminee a froid possedant une resistance au vieillissement et une formabilite superieure, et son procede de fabrication
WO2005061748A1 (fr) * 2003-12-23 2005-07-07 Posco Acier lamine a froid durci au four presentant une excellente formabilite, et procede de fabrication dudit acier

Non-Patent Citations (1)

* Cited by examiner, † Cited by third party
Title
None *

Also Published As

Publication number Publication date
WO2006118425A1 (fr) 2006-11-09
EP1888800A1 (fr) 2008-02-20
EP1888800A4 (fr) 2011-01-26

Similar Documents

Publication Publication Date Title
US20090126837A1 (en) Cold rolled steel sheet having superior formability and high yield ratio, process for producing the same
EP2138596B1 (fr) Feuille d'acier pour une utilisation dans une boîte métallique, et son procédé de fabrication
EP3476965B1 (fr) Plaque principale étamée à haute résistance à allongement élevé et procédé associé de double réduction à froid
EP3144406B1 (fr) Feuille d'acier laminé à froid de résistance élévée présentant une excellente ductilité, feuille d'acier galvanisé à chaud et son procédé de fabrication
EP2554699B1 (fr) Tôle d'acier présentant une résistance à la traction élevée et une meilleure ductilité et procédé de fabrication de cette dernière
KR101676137B1 (ko) 굽힘가공성과 구멍확장성이 우수한 고강도 냉연강판, 용융아연도금강판과 그 제조방법
EP2290111B1 (fr) Feuille d'acier double phase et son procédé de fabrication
EP3705594B1 (fr) Plaque brute en acier étamé multiphase à résistance élevée et procédé de fabrication associé
EP4166685A1 (fr) Acier à ultra-haute résistance présentant une excellente plasticité et son procédé de fabrication
EP3885462A1 (fr) Acier à deux phases laminé à froid à rendement élevé et son procédé de fabrication
EP1888799B1 (fr) Feuille d acier laminee a froid ayant une formabilite superieure et son procede de production
EP3378958B1 (fr) Tôle d'acier plaquée et procédé de fabrication associé
JP7482231B2 (ja) 低炭素低コスト超高強度多相鋼板／鋼帯およびその製造方法
EP3296416A1 (fr) Tôle d'acier, laminée à chaud, de résistance ultra-élevée, présentant une excellente aptitude au façonnage par pliage et son procédé de fabrication
EP1888800B1 (fr) Feuille d acier laminee a froid ayant une formabilite superieure et un rapport de rendement eleve et son procede de production
EP3231886B1 (fr) Tôle d'acier a phases complexes ayant une excellente formabilité et son procédé de fabrication
JPH03277741A (ja) 加工性、常温非時効性及び焼付け硬化性に優れる複合組織冷延鋼板とその製造方法
EP1885899B1 (fr) Feuille d'acier laminée à froid ayant un rapport de rendement élevé et moins d'anisotropie et son procédé de production
KR102426248B1 (ko) 선영성이 우수한 고강도 아연계 도금강판 및 그 제조방법
CN113227427A (zh) 延展性和加工性优异的高强度钢板及其制造方法
CN115491597B (zh) 一种780MPa增强塑性双相钢及其制备方法
CN110997963A (zh) 表面质量、强度及延展性优异的镀覆钢板
KR102468036B1 (ko) 성형성이 우수한 고강도 아연계 도금강판 및 그 제조방법
EP4273291A1 (fr) Tôle d'acier laminée à froid résistante à l'enfoncement, ayant d'excellentes propriétés de résistance à l'enfoncement, tôle d'acier plaqué résistante à l'enfoncement et procédé de fabrication de ces dernières
WO2005061748A1 (fr) Acier lamine a froid durci au four presentant une excellente formabilite, et procede de fabrication dudit acier

Legal Events

Date Code Title Description
PUAI Public reference made under article 153(3) epc to a published international application that has entered the european phase

Free format text: ORIGINAL CODE: 0009012

17P Request for examination filed

Effective date: 20071203

AK Designated contracting states

Kind code of ref document: A1

Designated state(s): AT BE BG CH CY CZ DE DK EE ES FI FR GB GR HU IE IS IT LI LT LU LV MC NL PL PT RO SE SI SK TR

DAX Request for extension of the european patent (deleted)
A4 Supplementary search report drawn up and despatched

Effective date: 20101223

17Q First examination report despatched

Effective date: 20130211

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: EXAMINATION IS IN PROGRESS

GRAP Despatch of communication of intention to grant a patent

Free format text: ORIGINAL CODE: EPIDOSNIGR1

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: GRANT OF PATENT IS INTENDED

INTG Intention to grant announced

Effective date: 20180614

GRAS Grant fee paid

Free format text: ORIGINAL CODE: EPIDOSNIGR3

GRAA (expected) grant

Free format text: ORIGINAL CODE: 0009210

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: THE PATENT HAS BEEN GRANTED

AK Designated contracting states

Kind code of ref document: B1

Designated state(s): AT BE BG CH CY CZ DE DK EE ES FI FR GB GR HU IE IS IT LI LT LU LV MC NL PL PT RO SE SI SK TR

REG Reference to a national code

Ref country code: GB

Ref legal event code: FG4D

REG Reference to a national code

Ref country code: CH

Ref legal event code: EP

Ref country code: AT

Ref legal event code: REF

Ref document number: 1062135

Country of ref document: AT

Kind code of ref document: T

Effective date: 20181115

REG Reference to a national code

Ref country code: IE

Ref legal event code: FG4D

REG Reference to a national code

Ref country code: DE

Ref legal event code: R096

Ref document number: 602006056761

Country of ref document: DE

REG Reference to a national code

Ref country code: NL

Ref legal event code: MP

Effective date: 20181107

REG Reference to a national code

Ref country code: LT

Ref legal event code: MG4D

REG Reference to a national code

Ref country code: AT

Ref legal event code: MK05

Ref document number: 1062135

Country of ref document: AT

Kind code of ref document: T

Effective date: 20181107

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: BG

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20190207

Ref country code: AT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

Ref country code: LV

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

Ref country code: FI

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

Ref country code: IS

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20190307

Ref country code: ES

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

Ref country code: LT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: GR

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20190208

Ref country code: PT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20190307

Ref country code: NL

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

Ref country code: SE

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: PL

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

Ref country code: CZ

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

Ref country code: DK

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

Ref country code: IT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

REG Reference to a national code

Ref country code: DE

Ref legal event code: R097

Ref document number: 602006056761

Country of ref document: DE

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: RO

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

Ref country code: EE

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

Ref country code: SK

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

PLBE No opposition filed within time limit

Free format text: ORIGINAL CODE: 0009261

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: NO OPPOSITION FILED WITHIN TIME LIMIT

26N No opposition filed

Effective date: 20190808

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: SI

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

REG Reference to a national code

Ref country code: CH

Ref legal event code: PL

GBPC Gb: european patent ceased through non-payment of renewal fee

Effective date: 20190503

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: CH

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20190531

Ref country code: LI

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20190531

Ref country code: MC

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

REG Reference to a national code

Ref country code: BE

Ref legal event code: MM

Effective date: 20190531

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: LU

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20190503

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: TR

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: IE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20190503

Ref country code: GB

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20190503

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: BE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20190531

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: CY

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20181107

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: HU

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT; INVALID AB INITIO

Effective date: 20060503

REG Reference to a national code

Ref country code: DE

Ref legal event code: R081

Ref document number: 602006056761

Country of ref document: DE

Owner name: POSCO CO., LTD, POHANG-SI, KR

Free format text: FORMER OWNER: POSCO, POHANG, KYUNGSANGBOOK, KR

Ref country code: DE

Ref legal event code: R081

Ref document number: 602006056761

Country of ref document: DE

Owner name: POSCO CO., LTD, POHANG- SI, KR

Free format text: FORMER OWNER: POSCO, POHANG, KYUNGSANGBOOK, KR

Ref country code: DE

Ref legal event code: R081

Ref document number: 602006056761

Country of ref document: DE

Owner name: POSCO HOLDINGS INC., KR

Free format text: FORMER OWNER: POSCO, POHANG, KYUNGSANGBOOK, KR

REG Reference to a national code

Ref country code: DE

Ref legal event code: R081

Ref document number: 602006056761

Country of ref document: DE

Owner name: POSCO CO., LTD, POHANG-SI, KR

Free format text: FORMER OWNER: POSCO HOLDINGS INC., SEOUL, KR

Ref country code: DE

Ref legal event code: R081

Ref document number: 602006056761

Country of ref document: DE

Owner name: POSCO CO., LTD, POHANG- SI, KR

Free format text: FORMER OWNER: POSCO HOLDINGS INC., SEOUL, KR

PGFP Annual fee paid to national office [announced via postgrant information from national office to epo]

Ref country code: FR

Payment date: 20230321

Year of fee payment: 18

PGFP Annual fee paid to national office [announced via postgrant information from national office to epo]

Ref country code: DE

Payment date: 20230320

Year of fee payment: 18