[go: up one dir, main page]
More Web Proxy on the site http://driver.im/

EP2871254B1 - Heissgewalztes stahlblech und verfahren zu seiner herstellung - Google Patents

Heissgewalztes stahlblech und verfahren zu seiner herstellung Download PDF

Info

Publication number
EP2871254B1
EP2871254B1 EP13837646.2A EP13837646A EP2871254B1 EP 2871254 B1 EP2871254 B1 EP 2871254B1 EP 13837646 A EP13837646 A EP 13837646A EP 2871254 B1 EP2871254 B1 EP 2871254B1
Authority
EP
European Patent Office
Prior art keywords
cooling
hot
steel sheet
less
rolled steel
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
EP13837646.2A
Other languages
English (en)
French (fr)
Other versions
EP2871254A4 (de
EP2871254A1 (de
Inventor
Chikara Kami
Sota GOTO
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.)
JFE Steel Corp
Original Assignee
JFE Steel Corp
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by JFE Steel Corp filed Critical JFE Steel Corp
Publication of EP2871254A1 publication Critical patent/EP2871254A1/de
Publication of EP2871254A4 publication Critical patent/EP2871254A4/de
Application granted granted Critical
Publication of EP2871254B1 publication Critical patent/EP2871254B1/de
Active legal-status Critical Current
Anticipated expiration legal-status Critical

Links

Images

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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B3/00—Rolling materials of special alloys so far as the composition of the alloy requires or permits special rolling methods or sequences ; Rolling of aluminium, copper, zinc or other non-ferrous metals
    • B21B3/02—Rolling special iron alloys, e.g. stainless steel
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/004—Heat treatment of ferrous alloys containing Cr and Ni
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • 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
    • C21D8/12—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties
    • C21D8/1216—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties the working step(s) being of interest
    • C21D8/1222—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
    • C21D8/12—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties
    • C21D8/1244—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties the heat treatment(s) being of interest
    • C21D8/1261—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of articles with special electromagnetic properties the heat treatment(s) being of interest 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
    • 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
    • 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
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel with vanadium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel with boron
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • 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
    • C21D2211/008—Martensite

Definitions

  • Patent Literature 1 can be used to manufacture a hot-rolled steel sheet having a tensile strength of 60 kg/mm 2 or more (590 MPa or more), a yield ratio of 85% or less, and high toughness represented by a fracture transition temperature of -60°C or less.
  • high-strength refers to a yield strength of 480 MPa or more at an angle of 30 degrees with the rolling direction and a tensile strength of 600 MPa or more in the sheet width direction.
  • high low-temperature toughness refers to a fracture transition temperature vTrs of -80°C or less in a Charpy impact test.
  • low yield ratio refers to a case where a steel sheet has a continuous yielding type stress-strain curve and a yield ratio of 85% or less.
  • steel sheets includes steel sheets and steel strips.
  • the present inventors also found that the surface microstructure of steel sheets composed of a tempered martensite single phase or a mixed phase of tempered martensite and tempered bainite is effective in preventing an uneven increase in the surface hardness of the steel sheets and providing steel pipes having the desired pipe shape and uniform ductility after pipe manufacturing.
  • Mn can dissolve in steel, improve quenching hardenability, and promote the formation of martensite. Mn is also an element that can lower the bainitic ferrite transformation start temperature and contribute to improved toughness of steel sheets by decreasing the microstructure size. These effects require a Mn content of 1.4% or more. However, a Mn content of more than 2.2% results in a heat affected zone having low toughness. Thus, the Mn content is limited to the range of 1.4% to 2.2%. The Mn content preferably ranges from 1.6% to 2.0% in terms of stable formation of massive martensite.
  • P can dissolve in steel and contribute to increased strength of steel sheets, but lowers toughness.
  • P is preferably minimized as an impurity.
  • a P content of up to 0.025% is acceptable.
  • the P content is limited to 0.025% or less, preferably 0.015% or less. Since an excessively low P content results in high refining costs, the P content is approximately 0.001% or more.
  • Al can act as a deoxidizing agent.
  • Al is an element that is effective in fixing N, which is responsible for strain aging. These effects require an Al content of 0.005% or more.
  • an Al content of more than 0.10% results in a high oxide content of steel and low toughness of base materials and welds.
  • the Al content is limited to the range of 0.005% to 0.10%, preferably 0.08% or less.
  • Nb can dissolved in steel or precipitate as carbonitride, can suppress coarsening and recrystallization of austenite grains, and allows rolling of austenite in a un-recrystallization temperature range.
  • Nb is also an element that can form fine carbide or carbonitride precipitates and contribute to increased strength of steel sheets.
  • Nb can precipitate as carbide or carbonitride on dislocations introduced by hot rolling, act as a nucleus for ⁇ ⁇ ⁇ transformation, promote the formation of bainitic ferrite in grains, and contribute to the formation of fine massive untransformed austenite, which results in the formation of fine massive martensite.
  • Ti can fix N as nitride and contribute to the prevention of cracking of slabs. Furthermore, Ti can form fine carbide precipitates and increase the strength of steel sheets. These effects require a Ti content of 0.001% or more. However, a high Ti content of more than 0.030% results in an excessively high bainitic ferrite transformation point and low toughness of steel sheets. Thus, the Ti content is limited to the range of 0.001% to 0.030%, preferably 0.005% to 0.025%.
  • Mo can contribute to improved quenching hardenability and promote the formation of martensite by moving C from bainitic ferrite to untransformed austenite and thereby improving the hardenability of the untransformed austenite.
  • Mo is an element that can dissolve in steel and contribute to increased strength of steel sheets by solid-solution hardening. These effects require a Mo content of 0.01% or more. However, a Mo content of more than 0.50% results in the formation of an excessive amount of martensite and low toughness of steel sheets. Furthermore, a large amount of expensive Mo results in high material costs. Thus, the Mo content is limited to the range of 0.01% to 0.50%, preferably 0.10% to 0.40%.
  • Cr has the effects of delaying ⁇ ⁇ ⁇ transformation, contributing to improved quenching hardenability, and promoting the formation of martensite. These effects require a Cr content of 0.01% or more. However, a Cr content of more than 0.50% tends to result in a frequent occurrence of defects in welds. Thus, the Cr content is limited to the range of 0.01% to 0.50%, preferably 0.20% to 0.45%.
  • Ni can contribute to improved quenching hardenability and promote the formation of martensite. Furthermore, Ni is an element that can contribute to further improved toughness. These effects require a Ni content of 0.01% or more. However, such effects level off at a Ni content of more than 0.50% and are not expected to be proportional to the Ni content beyond this threshold. A Ni content of more than 0.50% is therefore economically disadvantageous. Thus, the Ni content is limited to the range of 0.01% to 0.50%, preferably 0.30% to 0.45%.
  • Moeq % Mo + 0.36 Cr + 0.77 Mn + 0.07 Ni (wherein Mn, Ni, Cr, and Mo denote the corresponding element contents (% by mass)).
  • Moeq is an indicator of the quenching hardenability of untransformed austenite that remains in a steel sheet after the cooling step. Moeq of less than 1.4% results in insufficient quenching hardenability of untransformed austenite, which results in transformation of untransformed austenite into pearlite or the like during the subsequent coiling step. Moeq of more than 2.2% results in the formation of an excessive amount of martensite and low toughness. Thus, Moeq is limited to the range of 1.4% to 2.2%. Moeq of 1.5% or more results in a low yield ratio and further improved ductility. Thus, Moeq is more preferably 1.5% or more.
  • a hot-rolled steel sheet according to the present invention may contain one or two or more selected from Cu: 0.50% or less, V: 0.10% or less, and B: 0.0005% or less, and/or Ca: 0.0005% to 0.0050%.
  • Cu, V, and B are elements that can contribute to reinforcement of steel sheets and can be used as required.
  • V and Cu can contribute to reinforcement of steel sheets by solid-solution hardening or precipitation hardening.
  • B can segregate at grain boundaries and contribute to reinforcement of steel sheets due to improved quenching hardenability.
  • Cu 0.01% or more
  • V 0.01% or more
  • B 0.0001% or more
  • steel sheets having a V content of more than 0.10% have low weldability.
  • Steel sheets having a B content of more than 0.0005% have low toughness.
  • Steel sheets having a Cu content of more than 0.50% have poor hot workability.
  • Cu: 0.50% or less, V: 0.10% or less, and/or B: 0.0005% or less are preferred.
  • the incidental impurities may be N: 0.005% or less, O: 0.005% or less, Mg: 0.003% or less, and/or Sn: 0.005% or less.
  • a high-strength hot-rolled steel sheet with a low yield ratio has a composition as described above and has different microstructures on an outer surface layer (hereinafter also referred to simply as an outer layer) in the thickness direction and on an inner surface layer (hereinafter also referred to simply as an inner layer) in the thickness direction.
  • Steel pipes formed of a steel sheet having such different microstructures at different positions in the thickness direction can have a low yield ratio and uniform ductility.
  • an inner surface layer (inner layer) in the thickness direction refers to a region having a depth of 1.5 mm or more from the front and back sides of a steel sheet in the thickness direction.
  • An uneven cooling history of a hot-rolled steel sheet for example, cooling of a hot-rolled steel sheet through a transition boiling region results in a local increase in hardness and uneven hardness.
  • These problems can be avoided when the outer layer has a single-phase microstructure composed of a tempered martensite phase or a mixed microstructure composed of a tempered martensite phase and a tempered bainite phase.
  • the mixture ratio of the tempered martensite phase to the tempered bainite phase of the mixed microstructure is not particularly limited. From the perspective of temper softening treatment, the area fraction of the tempered martensite phase preferably ranges from 60% to 100%, and the area fraction of the tempered bainite phase preferably ranges from 0% to 40%.
  • the microstructure can be formed under certain manufacturing conditions, in particular, at a cumulative rolling reduction of 50% or more at a temperature of 930°C or less in finish rolling, and by sequentially performing a first cooling, second cooling, third cooling, and fourth cooling in a cooling step after the completion of the finish rolling.
  • the first cooling includes cooling the hot-rolled steel sheet to a martensitic transformation start temperature (Ms point) or less at an average cooling rate of 100°C/s or more with respect to surface temperature.
  • the second cooling includes, after the completion of the first cooling, holding the hot-rolled steel sheet for 1 s or more at a surface temperature of 600°C or more.
  • the third cooling includes, after the completion of the second cooling, cooling the hot-rolled steel sheet to a cooling stop temperature in the range of 600°C to 450°C at an average cooling rate in the range of 5°C to 30°C/s with respect to the temperature at half the thickness of the hot-rolled steel sheet.
  • the fourth cooling includes cooling the hot-rolled steel sheet from the cooling stop temperature of the third cooling to a coiling temperature at an average cooling rate of 2°C/s or less with respect to the temperature at half the thickness of the hot-rolled steel sheet or alternatively holding the hot-rolled steel sheet at a temperature in the range of the cooling stop temperature of the third cooling to the coiling temperature for 20 s or more.
  • the microstructure and area fraction can be identified and calculated by observing and measuring using the methods described below in the examples.
  • the hardness of a steel sheet at a depth of 0.5 mm from a surface thereof in the thickness direction is preferably 95% or less of the maximum hardness in the thickness direction.
  • the fact that the hardness of a hot-rolled steel sheet at a depth of 0.5 mm from a surface thereof in the thickness direction is not equal to the maximum hardness in the thickness direction is important in ensuring the workability of the hot-rolled steel sheet and the desired pipe shape after pipe manufacturing.
  • the maximum hardness in the thickness direction preferably corresponds to a Vickers hardness HV 0.5 of 165 points or more and 300 points or less, more preferably 280 points or less.
  • This hardness can be achieved under certain manufacturing conditions, in particular, by performing a first cooling and a second cooling in a cooling step after the completion of finish rolling, the first cooling including cooling the hot-rolled steel sheet to a martensitic transformation start temperature (Ms point) or less at an average cooling rate of 100°C/s or more with respect to surface temperature, the second cooling including, after the completion of the first cooling, holding the hot-rolled steel sheet for 1 s or more at a surface temperature of 600°C or more.
  • Ms point martensitic transformation start temperature
  • the hardness can be measured using the method described below in the examples.
  • the inner surface layer (inner layer) in the thickness direction has a microstructure composed of a main phase and a second phase.
  • the main phase is a bainitic ferrite phase.
  • the second phase is formed of massive martensite having an aspect ratio of less than 5.0 dispersed in the main phase.
  • the main phase herein refers to a phase having an occupied area of 50% by area or more.
  • the bainitic ferrite preferably has an area fraction of 85% or more, more preferably 88.3% or more.
  • the bainitic ferrite main phase has a substructure having a high dislocation density and contains needle-shaped ferrite and acicular ferrite.
  • the bainitic ferrite does not include polygonal ferrite having a very low dislocation density or semi(quasi)-polygonal ferrite including a substructure, such as fine subgrains.
  • the bainitic ferrite main phase must contain fine carbonitride precipitates.
  • the bainitic ferrite main phase has an average grain size of 10 ⁇ m or less. An average grain size of more than 10 ⁇ m results in insufficient work hardening ability in a region having a low strain of less than 5% and a decrease in yield strength due to bending in spiral pipe manufacturing.
  • the desired low-temperature toughness can be achieved by decreasing the average grain size of the main phase even when the steel sheet contains much martensite.
  • the second phase in the inner layer has massive martensite having an area fraction in the range of 1.4% to 15% and an aspect ratio of less than 5.0.
  • Massive martensite in the present invention is martensite formed from untransformed austenite at prior ⁇ grain boundaries or within prior ⁇ grains in a cooling process after rolling. In the present invention, such massive martensite is dispersed at prior ⁇ grain boundaries or between bainitic ferrite grains of the main phase. Martensite is harder than the main phase, can introduce a large number of mobile dislocations into bainitic ferrite during processing, and allows yield behavior of a continuous yielding type. Since martensite, which has higher tensile strength than bainitic ferrite, a low yield ratio can be achieved.
  • the martensite When the martensite is massive martensite having an aspect ratio of less than 5.0, the martensite can introduce more mobile dislocations into adjacent bainitic ferrite and effectively improve ductility. Martensite having an aspect ratio of 5.0 or more becomes rod-like martensite (non-massive martensite) and cannot achieve the desired low yield ratio. Nevertheless, rod-like martensite having an area fraction of less than 30% of the total amount of martensite is allowable.
  • the massive martensite preferably has an area fraction of 70% or more of the total amount of martensite.
  • the aspect ratio can be measured using the method described below in the examples.
  • Such effects require dispersion of massive martensite having an area fraction of 1.4% or more. It is difficult to achieve the desired low yield ratio with massive martensite having an area fraction of less than 1.4%. When the massive martensite has an area fraction of more than 15%, the low-temperature toughness is significantly decreased. Thus, the area fraction of massive martensite is limited to the range of 1.4% to 15%, preferably 10% or less.
  • the second phase may contain bainite having an area fraction of approximately 7.0% or less.
  • the microstructure can be formed under certain manufacturing conditions, in particular, at a cumulative rolling reduction of 50% or more at a temperature of 930°C or less in finish rolling, and by sequentially performing a first cooling, second cooling, third cooling, and fourth cooling in a cooling step after the completion of the finish rolling.
  • the first cooling includes cooling the hot-rolled steel sheet to a martensitic transformation start temperature (Ms point) or less at an average cooling rate of 100°C/s or more with respect to surface temperature.
  • the second cooling includes, after the completion of the first cooling, holding the hot-rolled steel sheet for 1 s or more at a surface temperature of 600°C or more.
  • the third cooling includes, after the completion of the second cooling, cooling the hot-rolled steel sheet to a cooling stop temperature in the range of 600°C to 450°C at an average cooling rate in the range of 5°C to 30°C/s with respect to the temperature at half the thickness of the hot-rolled steel sheet.
  • the fourth cooling includes cooling the hot-rolled steel sheet from the cooling stop temperature of the third cooling to a coiling temperature at an average cooling rate of 2°C/s or less with respect to the temperature at half the thickness of the hot-rolled steel sheet or alternatively holding the hot-rolled steel sheet at a temperature in the range of the cooling stop temperature of the third cooling to the coiling temperature for 20 s or more.
  • the massive martensite preferably has a maximum size of 5.0 ⁇ m or less and an average size in the range of 0.5 to 3.0 ⁇ m.
  • Coarse massive martensite having an average size of more than 3.0 ⁇ m tends to act as a starting point of brittle fracture or promote crack propagation and lowers the low-temperature toughness.
  • Excessively fine massive martensite grains having an average size of less than 0.5 ⁇ m result in a decreased number of mobile dislocations introduced into adjacent bainitic ferrite.
  • Massive martensite having a maximum size of more than 5.0 ⁇ m results in low toughness.
  • the massive martensite preferably has a maximum size of 5.0 ⁇ m or less and an average size in the range of 0.5 to 3.0 ⁇ m.
  • the arithmetic mean of the "diameters" of grains is the "average" size of the massive martensite. At least 100 martensite grains are subjected to the measurement.
  • the microstructure can be formed under certain manufacturing conditions, in particular, at a cumulative rolling reduction of 50% or more at a temperature of 930°C or less in finish rolling, and by sequentially performing a first cooling, second cooling, third cooling, and fourth cooling in a cooling step after the completion of the finish rolling.
  • the first cooling includes cooling the hot-rolled steel sheet to a martensitic transformation start temperature (Ms point) or less at an average cooling rate of 100°C/s or more with respect to surface temperature.
  • the second cooling includes, after the completion of the first cooling, holding the hot-rolled steel sheet for 1 s or more at a surface temperature of 600°C or more.
  • the third cooling includes, after the completion of the second cooling, cooling the hot-rolled steel sheet to a cooling stop temperature in the range of 600°C to 450°C at an average cooling rate in the range of 5°C to 30°C/s with respect to the temperature at half the thickness of the hot-rolled steel sheet.
  • the fourth cooling includes cooling the hot-rolled steel sheet from the cooling stop temperature of the third cooling to a coiling temperature at an average cooling rate of 2°C/s or less with respect to the temperature at half the thickness of the hot-rolled steel sheet or alternatively holding the hot-rolled steel sheet at a temperature in the range of the cooling stop temperature of the third cooling to the coiling temperature for 20 s or more.
  • microstructure, area fraction, and average grain size can be identified and calculated by observing and measuring using the methods described below in the examples.
  • steel having a composition as described above is subjected to a hot-rolling step, a cooling step, and a coiling step to form a hot-rolled steel sheet.
  • the steel may be manufactured by any method.
  • molten steel having a composition as described above is smelted using a known melting method, such as using a converter or an electric furnace, and the molten steel is formed into steel, such as a slab, using a known casting method, such as a continuous casting process.
  • the steel is subjected to the hot-rolling step.
  • the hot-rolling step includes heating steel having a composition as described above to a heating temperature in the range of 1050°C to 1300°C, rough-rolling the heated steel to form a sheet bar, and finish-rolling the sheet bar such that the cumulative rolling reduction at a temperature of 930°C or less is 50% or more, thereby forming a hot-rolled steel sheet.
  • Heating temperature 1050°C to 1300°C
  • Steel used in the present invention essentially contains Nb and Ti, as described above.
  • coarse carbide and nitride must be once dissolved in steel and then finely precipitated.
  • the steel is heated to a heating temperature of 1050°C or more.
  • a heating temperature of more than 1300°C results in coarsening of crystal grains and steel sheets having low toughness.
  • the heating temperature for the steel is limited to the range of 1050°C to 1300°C.
  • the steel heated to the heating temperature is subjected to rough rolling to form a sheet bar.
  • the steel may be subjected to rough rolling under any conditions, provided that the sheet bar has the desired size and shape.
  • the sheet bar is then subjected to finish rolling to form a hot-rolled steel sheet having the desired size and shape.
  • finish rolling the cumulative rolling reduction at a temperature of 930°C or less is 50% or more.
  • the cumulative rolling reduction at a temperature of 930°C or less is 50% or more in order to decrease the size of bainitic ferrite and finely disperse massive martensite in the inner layer microstructure.
  • a cumulative rolling reduction of less than 50% at a temperature of 930°C or less results in an insufficient rolling reduction and a lack of a fine bainitic ferrite main phase in the inner layer microstructure. This also results in insufficient dislocations that act as precipitation sites for NbC and the like, which promotes nucleation in ⁇ ⁇ ⁇ transformation, and insufficient formation of bainitic ferrite in grains. It is therefore impossible to keep a large number of finely dispersed massive untransformed ⁇ grains for forming massive martensite.
  • the cumulative rolling reduction at a temperature of 930°C or less is limited to 50% or more.
  • the cumulative rolling reduction is preferably 80% or less. Such effects level off at a rolling reduction of more than 80%. Furthermore, a rolling reduction of more than 80% may result in a frequent occurrence of separation and low absorbed energy in a Charpy impact test.
  • the finishing temperature of the finish rolling preferably ranges from 850°C to 760°C in terms of steel sheet toughness, steel sheet strength, and rolling load.
  • the finishing temperature of the finish rolling is as high as more than 850°C, the rolling reduction per pass must be increased to achieve the cumulative rolling reduction of 50% or more at a temperature of 930°C or less, which sometimes results in increased rolling load.
  • the finishing temperature of the finish rolling is as low as less than 760°C, this sometimes results in the formation of ferrite during rolling, coarsening of the microstructure and precipitates, and decreases in low-temperature toughness and strength.
  • the hot-rolled steel sheet is then subjected to the cooling step.
  • the cooling step includes first cooling, second cooling, third cooling, and fourth cooling in this order.
  • the first cooling is started immediately after the completion of the finish rolling and including cooling the hot-rolled steel sheet to a martensitic transformation start temperature (Ms point) or less at an average cooling rate of 100°C/s or more with respect to surface temperature.
  • the second cooling includes, after the completion of the first cooling, holding the hot-rolled steel sheet for 1 s or more at a surface temperature of 600°C or more.
  • the third cooling includes, after the completion of the second cooling, cooling the hot-rolled steel sheet to a cooling stop temperature in the range of 600°C to 450°C at an average cooling rate in the range of 5°C to 30°C/s with respect to the temperature at half the thickness of the hot-rolled steel sheet.
  • the fourth cooling includes cooling the hot-rolled steel sheet from the cooling stop temperature of the third cooling to a coiling temperature at an average cooling rate of 2°C/s or less with respect to the temperature at half the thickness of the hot-rolled steel sheet or alternatively holding the hot-rolled steel sheet at a temperature in the range of the cooling stop temperature of the third cooling to the coiling temperature for 20 s or more.
  • the coiling step includes coiling the hot-rolled steel sheet at a surface temperature of 450°C or more.
  • Cooling is started immediately, within 15 s, after the completion of the finish rolling.
  • the holding time at the martensitic transformation start temperature (Ms point) or less with respect to surface temperature depends on the desired surface microstructure and is 10 s or less, preferably 7 s or less. Holding the hot-rolled steel sheet at a temperature of the Ms point or less for a long time results in an excessively high occupied area of a single phase formed of a martensite phase or a mixed microstructure composed of a martensite phase and a bainite phase, which results in a lower thickness percentage of the desired microstructure.
  • Test pieces were taken from the hot-rolled steel sheet and were subjected to microstructure observation, a tensile test, an impact test, and a hardness test.
  • test methods are as follows:
  • a test piece for microstructure observation was taken from the hot-rolled steel sheet such that a cross section thereof in the rolling direction (L cross section) served as an observation surface. After the test piece was polished and was etched with nital, the microstructure of the test piece was observed and photographed with an optical microscope (magnification ratio: 500) or an electron microscope (magnification ratio: 2000). The type of microstructure, the fraction (area fraction) of the microstructure of each phase, and the average grain size were determined from the photograph of the inner layer microstructure with an image analyzing apparatus. For the outer layer, only the type of microstructure was identified from the microstructure photograph.
  • the average grain size of the bainitic ferrite main phase in the inner layer microstructure was determined using an intercept method in accordance with JIS G 0552.
  • the aspect ratio of martensite grains was calculated as the ratio (the length along the major axis)/(the length along the minor axis) of the length of a grain in the longitudinal direction or in a direction of the maximum grain size (the length along the major axis) to the length of the grain in a direction perpendicular to the longitudinal direction (the length along the minor axis).
  • Martensite grains having an aspect ratio of less than 5.0 were defined as massive martensite.
  • Martensite grains having an aspect ratio of 5.0 or more were referred to as "rod-like" martensite.
  • the average size of massive martensite in the steel sheet was calculated by determining half the sum of the length along the major axis and the length along the minor axis of each massive martensite grain as the diameter thereof and calculating the arithmetic mean of the diameters.
  • the maximum diameter of each massive martensite grain was the maximum size of the massive martensite. At least 100 martensite grains were subjected to the measurement.
  • Test pieces for tensile test (full-thickness test pieces specified in API-5L, (width: 38.1 mm, GL: 50 mm)) were taken from the hot-rolled steel sheet such that the tensile direction was perpendicular to the rolling direction (sheet width direction) or at an angle of 30 degrees with the rolling direction.
  • a tensile test was performed in accordance with the ASTM A 370 specification to determine tensile properties (yield strength YS and tensile strength TS) .
  • V-notched test pieces were taken from the hot-rolled steel sheet such that the longitudinal direction of the test pieces was perpendicular to the rolling direction, and were subjected to a Charpy impact test in accordance with the ASTM A 370 specification to determine the fracture transition temperature vTrs (°C).
  • Test pieces for hardness measurement were taken from the hot-rolled steel sheet.
  • the cross section hardness of the test pieces was measured with a Vickers hardness tester (test force: 4.9 N) (load: 500 g).
  • the cross section hardness of each of the test pieces was continuously measured at intervals of 0.5 mm from a surface of the steel sheet in the thickness direction.
  • the hardness at a depth of 0.5 mm from the surface of the steel sheet in the thickness direction (depth direction) and the maximum hardness in the thickness direction were determined.
  • the hardness distribution was judged to be good when the maximum hardness in the thickness direction was 300 points or less, and the hardness at a depth of 0.5 mm from the surface was 95% or less of the maximum hardness in the thickness direction.
  • a spiral steel pipe (outer diameter: 1067 mm ⁇ ) was then manufactured by using a spiral pipe manufacturing process using the hot-rolled steel sheet as a material for pipes.
  • Test pieces for tensile test (test pieces specified in API) were taken from the steel pipe such that the tensile direction was the circumferential direction of the pipe, and were subjected to a tensile test in accordance with the ASTM A 370 specification to measure tensile properties (yield strength YS and tensile strength TS).
  • Cooling step Coiling step Note Heating Rough rolling Finish rolling Cooling start time (s) First cooling*2 Second cooling*2 Third cooling*3 Fourth cooling*3 Coiling temperature *11 (°C) Heating temperature (°C) Thickness of sheet bar (mm) Finish rolling temperature (°C) Rolling reduction *1 (%) Thickness (mm) Average cooling rate *4 (°C/s) Cooling stop temperature *5 (°C) Ms (°C) Final surface temperature *6 (°C) Holding time *7 (s) Average cooling rate *8 (°C/s) Cooling stop temperature (°C) Average cooling rate *9 (°C/s) Holding time *10 (s) 1 A 1059 51 768 74 8 2.4 111 373 406 608 1.4 18 551 1.5 - 538 Example 2 A 1091 55 759 55 25 7.6 145 372 406 613 2.7 28 558 0.5 - 536 Example 3 A 1099 51 777 61 16 4.8 122 372 406 603 2.0
  • All the examples provided high-strength high-toughness hot-rolled steel sheets having a low yield ratio without particular heat treatment.
  • These hot-rolled steel sheets had a yield strength of 480 MPa or more at an angle of 30 degrees with the rolling direction, a tensile strength of 600 MPa or more in the sheet width direction, high toughness represented by a fracture transition temperature vTrs of-80°C or less, and a yield ratio of 85% or less.
  • the comparative examples outside the scope of the present invention could not provide hot-rolled steel sheets having the desired characteristics because of low toughness or a high yield ratio.

Landscapes

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

Claims (6)

  1. Warmgewalztes Stahlblech, das eine Zusammensetzung aufweist, basierend auf Massenprozent, bestehend aus:
    C: 0,03% bis 0,10%, Si: 0,01% bis 0,50%, Mn: 1,4% bis 2,2%, P: 0,001% bis 0,025%, S: 0,0001% bis 0,005%, Al: 0,005% bis 0,10%, Nb: 0,02% bis 0,10%, Ti: 0,001% bis 0,030%, Mo: 0,01% bis 0,50%, Cr: 0,01% bis 0,50%, Ni: 0,01% bis 0,50%, gegebenenfalls einem oder zwei oder mehreren, ausgewählt aus Cu: 0,50% oder weniger, V: 0,10% oder weniger und B: 0,0005% oder weniger, gegebenenfalls Ca: 0,0005% bis 0,0050%, auf die Masse bezogen, und einem Rest aus Fe und unvermeidbaren Verunreinigungen, worin unvermeidbare Verunreinigungen N: 0,005% oder weniger, O: 0,005% oder weniger, Mg: 0,003% oder weniger und/oder Sn: 0,005% oder weniger umfassen,
    worin das warmgewalzte Stahlblech eine Innenschicht mit einer Tiefe von 1,5 mm oder mehr von der Vorder- und Rückseite des Stahlblechs in Dickenrichtung umfasst und eine Mikrostruktur aufweist, die eine Hauptphase mit einer eingenommenen Fläche von 50% der Fläche oder mehr und eine zweite Phase enthält, wobei die Hauptphase bainitischer Ferrit mit einer durchschnittlichen Korngröße von 10 µm oder weniger ist, wobei die zweite Phase Massivmartensit mit einem Flächenanteil im Bereich von 1,4% bis 15% und einem Aspektverhältnis von weniger als 5,0 aufweist,
    das warmgewalzte Stahlblech eine Außenschicht mit einer Mikrostruktur umfasst, die eine getemperte Martensitphase oder eine getemperte Martensitphase und eine getemperte Bainitphase enthält, worin der Flächenanteil der getemperten Martensitphase im Bereich von 60% bis 100% liegt und der Flächenanteil der getemperten Bainitphase im Bereich von 0% bis 40% liegt und worin die Zusammensetzung ein durch die nachfolgende Formel (1) definiertes Moeq im Bereich von 1,4% bis 2,2%, auf die Masse bezogen, aufweist: Moeq % = Mo + 0,36 Cr + 0,77 Mn + 0,07 Ni
    Figure imgb0005
     worin Mn, Ni, Cr und Mo die Gehalte des entsprechenden Elements in Massen-% anzeigen.
  2. Warmgewalztes Stahlblech gemäß Anspruch 1, umfassend Ca: 0,0005% bis 0,0050%, auf die Masse bezogen.
  3. Warmgewalztes Stahlblech gemäß Anspruch 1 oder 2, worin der Massivmartensit eine maximale Größe von 5,0 µm oder weniger und eine durchschnittliche Größe im Bereich von 0,5 bis 3,0 µm aufweist.
  4. Warmgewalztes Stahlblech gemäß mindestens einem der Ansprüche 1 bis 3, worin die Härte des warmgewalzten Stahlblechs bei einer Tiefe von 0,5 mm von der Oberfläche in Dickenrichtung 95% oder weniger der maximalen Härte in der Dickenrichtung beträgt.
  5. Verfahren zur Herstellung eines warmgewalzten Stahlblechs wie in Anspruch 1 definiert, das Verfahren umfassend:
    einen Schritt des Warmwalzens, einen Schritt des Abkühlens und einen Schritt des Aufrollens von Stahl, so dass das warmgewalzte Stahlblech gebildet wird,
    worin der Stahl, basierend auf Massenprozent, aus C: 0,03% bis 0,10%, Si: 0,01% bis 0,50%, Mn: 1,4% bis 2,2%, P: 0,001% bis 0,025%, S: 0,0001% bis 0,005%, Al: 0,005% bis 0,10%, Nb: 0,02% bis 0,10%, Ti: 0,001% bis 0,030%, Mo: 0,01% bis 0,50%, Cr: 0,01% bis 0,50%, Ni: 0,01% bis 0,50%, gegebenenfalls einem oder zwei oder mehreren, ausgewählt aus Cu: 0,50% oder weniger, V: 0,10% oder weniger und B: 0,0005% oder weniger, gegebenenfalls Ca: 0,0005% bis 0,0050%, auf die Masse bezogen, und einem Rest aus Fe und unvermeidbaren Verunreinigungen besteht, worin unvermeidbare Verunreinigungen N: 0,005% oder weniger, O: 0,005% oder weniger, Mg: 0,003% oder weniger und/oder Sn: 0,005% oder weniger umfassen,
    der Schritt des Warmwalzens das Erwärmen des Stahls auf eine Erwärmungstemperatur im Bereich von 1050°C bis 1300°C, das Vorwalzen des erwärmten Stahls, so dass ein Vorblech gebildet wird, und das Endwalzen des Vorblechs, so dass die kumulative Walzreduktion bei einer Temperatur von 930°C oder weniger 50% oder mehr beträgt, wodurch ein warmgewalztes Stahlblech gebildet wird, umfasst,
    der Schritt des Abkühlens ein erstes Abkühlen, ein zweites Abkühlen, ein drittes Abkühlen und ein viertes Abkühlen in dieser Reihenfolge umfasst, wobei das erste Abkühlen innerhalb von 15 s nach Beendigung des Endwalzens begonnen wird und das Abkühlen des warmgewalzten Stahlblechs auf eine martensitische Umwandlungsstarttemperatur oder weniger bei einer durchschnittlichen Abkühlgeschwindigkeit von 100°C/s oder mehr in Bezug auf die Oberflächentemperatur umfasst, das zweite Abkühlen, nach Beendigung des ersten Abkühlens, das Halten des warmgewalzten Stahlblechs für 1 s oder mehr bei einer Oberflächentemperatur von 600°C oder mehr umfasst, das dritte Abkühlen, nach Beendigung des zweiten Abkühlens, das Abkühlen des warmgewalzten Stahlblechs auf eine Abkühlstopptemperatur im Bereich von 600°C bis 450°C bei einer durchschnittlichen Abkühlgeschwindigkeit im Bereich von 5°C/s bis 30°C/s in Bezug auf die Temperatur bei der Hälfte der Dicke des warmgewalzten Stahlblechs umfasst, das vierte Abkühlen das Abkühlen des warmgewalzten Stahlblechs von der Abkühlstopptemperatur des dritten Abkühlens auf eine Aufrolltemperatur bei einer durchschnittlichen Abkühlgeschwindigkeit von 2°C/s oder weniger in Bezug auf die Temperatur bei der Hälfte der Dicke des warmgewalzten Stahlblechs oder alternativ das Halten des warmgewalzten Stahlblechs bei einer Temperatur im Bereich von der Abkühlstopptemperatur des dritten Abkühlens bis zur Aufrolltemperatur für 20 s oder mehr umfasst,
    der Schritt des Aufrollens das Aufrollen des warmgewalzten Stahlblechs bei einer Oberflächentemperatur von 450°C oder mehr umfasst, und worin der Stahl ein durch die nachfolgende Formel (1) definiertes Moeq im Bereich von 1,4% bis 2,2%, auf die Masse bezogen, aufweist: Moeq % = Mo + 0,36 Cr + 0,77 Mn + 0,07 Ni
    Figure imgb0006
     worin Mn, Ni, Cr und Mo die Gehalte des entsprechenden Elements in Massen-% anzeigen.
  6. Verfahren zur Herstellung eines warmgewalzten Stahlblechs gemäß Anspruch 5, worin der Stahl Ca: 0,0005% bis 0,0050%, auf die Masse bezogen, enthält.
EP13837646.2A 2012-09-13 2013-09-11 Heissgewalztes stahlblech und verfahren zu seiner herstellung Active EP2871254B1 (de)

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
JP2012201266 2012-09-13
PCT/JP2013/005388 WO2014041802A1 (ja) 2012-09-13 2013-09-11 熱延鋼板およびその製造方法

Publications (3)

Publication Number Publication Date
EP2871254A1 EP2871254A1 (de) 2015-05-13
EP2871254A4 EP2871254A4 (de) 2015-11-18
EP2871254B1 true EP2871254B1 (de) 2020-06-24

Family

ID=50277943

Family Applications (1)

Application Number Title Priority Date Filing Date
EP13837646.2A Active EP2871254B1 (de) 2012-09-13 2013-09-11 Heissgewalztes stahlblech und verfahren zu seiner herstellung

Country Status (7)

Country Link
US (1) US20150232970A1 (de)
EP (1) EP2871254B1 (de)
JP (1) JP5605527B2 (de)
KR (1) KR101702794B1 (de)
CN (1) CN104619876B (de)
IN (1) IN2015DN00772A (de)
WO (1) WO2014041802A1 (de)

Families Citing this family (16)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2014041801A1 (ja) 2012-09-13 2014-03-20 Ｊｆｅスチール株式会社 熱延鋼板およびその製造方法
JP2018532045A (ja) * 2015-09-22 2018-11-01 タタ、スティール、アイモイデン、ベスローテン、フェンノートシャップＴａｔａ Ｓｔｅｅｌ Ｉｊｍｕｉｄｅｎ Ｂｖ 優れた伸びフランジ成形性を有するロール成形可能な高強度熱間圧延鋼板、および該鋼の製造方法
WO2017130875A1 (ja) * 2016-01-27 2017-08-03 Ｊｆｅスチール株式会社 電縫鋼管用高強度熱延鋼板およびその製造方法
KR101830437B1 (ko) * 2016-04-25 2018-02-20 현대자동차주식회사 3층 구조를 갖는 고인성 열처리 강관 및 이의 제조방법
JP6624103B2 (ja) * 2017-02-06 2019-12-25 Ｊｆｅスチール株式会社 高強度熱延鋼板およびその製造方法
WO2018150955A1 (ja) * 2017-02-17 2018-08-23 Ｊｆｅスチール株式会社 高強度熱延鋼板およびその製造方法
MX2019012110A (es) * 2017-04-28 2019-11-28 Nippon Steel Corp Lamina de acero de alta resistencia y metodo de produccion de la misma.
KR102031450B1 (ko) * 2017-12-24 2019-10-11 주식회사 포스코 고강도 강판 및 그 제조방법
KR101999015B1 (ko) 2017-12-24 2019-07-10 주식회사 포스코 취성균열 전파 저항성이 우수한 구조용 강재 및 그 제조방법
KR102200224B1 (ko) * 2018-12-19 2021-01-08 주식회사 포스코 취성파괴 저항성이 우수한 구조용 강재 및 그 제조방법
CN113677816B (zh) * 2019-03-29 2022-11-22 杰富意钢铁株式会社 电阻焊钢管及其制造方法、以及钢管桩
JP7147960B2 (ja) * 2019-11-27 2022-10-05 Ｊｆｅスチール株式会社 鋼板およびその製造方法
WO2021123877A1 (en) * 2019-12-17 2021-06-24 Arcelormittal Hot rolled steel sheet and method of manufacturing thereof
WO2021125386A1 (ko) * 2019-12-18 2021-06-24 주식회사 포스코 타발성 및 재질 균일성이 우수한 고강도 열연강판 및 그 제조방법
CN114107612B (zh) * 2021-11-30 2023-04-18 马鞍山钢铁股份有限公司 H型钢回火热处理设计方法、抗震耐火建筑结构用热轧h型钢及其回火热处理方法
WO2023162522A1 (ja) * 2022-02-24 2023-08-31 Ｊｆｅスチール株式会社 鋼板およびその製造方法

Citations (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2012172256A (ja) * 2011-02-24 2012-09-10 Jfe Steel Corp 低温靭性に優れた低降伏比高強度熱延鋼板およびその製造方法

Family Cites Families (12)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2510187B2 (ja) 1987-03-17 1996-06-26 川崎製鉄株式会社 低温靭性に優れた低降伏比高張力ラインパイプ用熱延鋼板の製造方法
DE19612818C2 (de) * 1996-03-30 1998-04-09 Schloemann Siemag Ag Verfahren zur Kühlung walzwarmer Stahlprofile
JPH10176239A (ja) 1996-10-17 1998-06-30 Kobe Steel Ltd 高強度低降伏比パイプ用熱延鋼板及びその製造方法
JP5011773B2 (ja) 2005-03-24 2012-08-29 Ｊｆｅスチール株式会社 低温靭性に優れた低降伏比電縫鋼管の製造方法
JP5223379B2 (ja) * 2007-03-08 2013-06-26 新日鐵住金株式会社 低温靭性に優れるスパイラルパイプ用高強度熱延鋼板およびその製造方法
WO2009014238A1 (ja) * 2007-07-23 2009-01-29 Nippon Steel Corporation 変形特性に優れた鋼管及びその製造方法
GB0719457D0 (en) * 2007-10-04 2007-11-14 Skf Ab Heat-treatment process for a steel
JP5162382B2 (ja) 2008-09-03 2013-03-13 株式会社神戸製鋼所 低降伏比高靭性厚鋼板
EP2392682B1 (de) * 2009-01-30 2019-09-11 JFE Steel Corporation Dickes heizgewalztes stahlblech mit hoher bruchfestigkeit sowie hervorragender niedrigtemperaturbeständigkeit sowie herstellungsverfahren dafür
KR20160057492A (ko) * 2009-01-30 2016-05-23 제이에프이 스틸 가부시키가이샤 내 ｈｉｃ 성이 우수한 후육 고장력 열연강판 및 그 제조 방법
JP5499731B2 (ja) 2009-01-30 2014-05-21 Ｊｆｅスチール株式会社 耐ｈｉｃ性に優れた厚肉高張力熱延鋼板及びその製造方法
JP6006477B2 (ja) * 2011-06-24 2016-10-12 株式会社神戸製鋼所 低温靭性と強度のバランスに優れた高強度鋼板の製造方法、及びその制御方法

Patent Citations (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2012172256A (ja) * 2011-02-24 2012-09-10 Jfe Steel Corp 低温靭性に優れた低降伏比高強度熱延鋼板およびその製造方法

Also Published As

Publication number Publication date
JP5605527B2 (ja) 2014-10-15
US20150232970A1 (en) 2015-08-20
CN104619876A (zh) 2015-05-13
WO2014041802A1 (ja) 2014-03-20
EP2871254A4 (de) 2015-11-18
JPWO2014041802A1 (ja) 2016-08-12
BR112015005419A2 (pt) 2017-07-04
IN2015DN00772A (de) 2015-07-03
CN104619876B (zh) 2016-12-21
KR101702794B1 (ko) 2017-02-03
KR20150038747A (ko) 2015-04-08
EP2871254A1 (de) 2015-05-13

Similar Documents

Publication Publication Date Title
EP2871254B1 (de) Heissgewalztes stahlblech und verfahren zu seiner herstellung
EP2735622B1 (de) Hochfeste warmgewalzte stahlplatte mit niedrigem ertragsverhältnis und hervorragender niedrigtemperaturbeständigkeit sowie herstellungsverfahren dafür
US10900104B2 (en) Hot rolled steel sheet and method for manufacturing the same
EP2949772B1 (de) Heissgewalztes stahlblech und verfahren zur herstellung davon
EP2309014B1 (de) Kaltgewalzte dicke stahlbleche mit hoher bruchfestigkeit sowie hervorragender niedrigtemperaturbeständigkeit sowie herstellungsverfahren dafür
JP5776398B2 (ja) 低温靭性に優れた低降伏比高強度熱延鋼板およびその製造方法
EP2258886B1 (de) Hochfestes feuerverzinktes stahlblech mit hervorragender verarbeitbarkeit und herstellungsverfahren dafür
JP5679114B2 (ja) 低温靭性に優れた低降伏比高強度熱延鋼板およびその製造方法
EP3608434B1 (de) Widerstandsgeschweisstes stahlrohr im walzzustand für leitungsrohr und warmgewalztes stahlblech
EP2133441A1 (de) Hochfeste, heissgewalzte stahlplatte mit exzellenter niedrigtemperaturfestigkeit für ein spiralrohr und herstellungsverfahren dafür
KR100815717B1 (ko) 수소유기균열 저항성과 저온인성이 우수한 고강도 대구경라인파이프 강재 및 그 제조방법
EP3246427A1 (de) Hochfestes widerstandsgeschweisstes stahlrohr und herstellungsverfahren dafür
JP2015190026A (ja) ラインパイプ用厚肉高強度電縫鋼管およびその製造方法
JP6519024B2 (ja) 低温靭性に優れた低降伏比高強度熱延鋼板の製造方法
EP3428299A1 (de) Elektronahtgeschweisstes stahlrohr für leitungsrohre
JP6565890B2 (ja) 低温靭性に優れた低降伏比高強度熱延鋼板の製造方法
BR112015005419B1 (pt) Chapa de aço laminada a quente e método para fabricar a mesma

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: 20150204

AK Designated contracting states

Kind code of ref document: A1

Designated state(s): AL AT BE BG CH CY CZ DE DK EE ES FI FR GB GR HR HU IE IS IT LI LT LU LV MC MK MT NL NO PL PT RO RS SE SI SK SM TR

AX Request for extension of the european patent

Extension state: BA ME

RA4 Supplementary search report drawn up and despatched (corrected)

Effective date: 20151020

RIC1 Information provided on ipc code assigned before grant

Ipc: C21D 8/02 20060101ALI20151014BHEP

Ipc: C22C 38/00 20060101AFI20151014BHEP

Ipc: B21B 3/02 20060101ALI20151014BHEP

Ipc: C22C 38/58 20060101ALI20151014BHEP

DAX Request for extension of the european patent (deleted)
STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: EXAMINATION IS IN PROGRESS

17Q First examination report despatched

Effective date: 20170407

REG Reference to a national code

Ref country code: DE

Ref legal event code: R079

Ref document number: 602013070203

Country of ref document: DE

Free format text: PREVIOUS MAIN CLASS: C22C0038000000

Ipc: C21D0006000000

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

RIC1 Information provided on ipc code assigned before grant

Ipc: C22C 38/02 20060101ALI20191217BHEP

Ipc: C22C 38/50 20060101ALI20191217BHEP

Ipc: C22C 38/00 20060101ALI20191217BHEP

Ipc: C21D 9/46 20060101ALI20191217BHEP

Ipc: C22C 38/04 20060101ALI20191217BHEP

Ipc: C22C 38/44 20060101ALI20191217BHEP

Ipc: C21D 8/02 20060101ALI20191217BHEP

Ipc: C22C 38/48 20060101ALI20191217BHEP

Ipc: C22C 38/46 20060101ALI20191217BHEP

Ipc: C21D 8/12 20060101ALI20191217BHEP

Ipc: C22C 38/06 20060101ALI20191217BHEP

Ipc: C21D 6/00 20060101AFI20191217BHEP

Ipc: C22C 38/54 20060101ALI20191217BHEP

Ipc: C22C 38/42 20060101ALI20191217BHEP

Ipc: C22C 38/58 20060101ALI20191217BHEP

INTG Intention to grant announced

Effective date: 20200117

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): AL AT BE BG CH CY CZ DE DK EE ES FI FR GB GR HR HU IE IS IT LI LT LU LV MC MK MT NL NO PL PT RO RS SE SI SK SM 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

REG Reference to a national code

Ref country code: AT

Ref legal event code: REF

Ref document number: 1283957

Country of ref document: AT

Kind code of ref document: T

Effective date: 20200715

REG Reference to a national code

Ref country code: DE

Ref legal event code: R096

Ref document number: 602013070203

Country of ref document: DE

REG Reference to a national code

Ref country code: IE

Ref legal event code: FG4D

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

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: 20200624

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: 20200624

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: 20200624

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: 20200925

Ref country code: NO

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: 20200924

REG Reference to a national code

Ref country code: LT

Ref legal event code: MG4D

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: 20200924

Ref country code: HR

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: 20200624

Ref country code: RS

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: 20200624

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: 20200624

REG Reference to a national code

Ref country code: NL

Ref legal event code: MP

Effective date: 20200624

REG Reference to a national code

Ref country code: AT

Ref legal event code: MK05

Ref document number: 1283957

Country of ref document: AT

Kind code of ref document: T

Effective date: 20200624

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

Ref country code: AL

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: 20200624

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: 20200624

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

Ref country code: SM

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: 20200624

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: 20200624

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: 20200624

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: 20200624

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: 20200624

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: 20200624

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: 20201026

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: 20200624

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: 20200624

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: 20200624

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: 20201024

REG Reference to a national code

Ref country code: DE

Ref legal event code: R097

Ref document number: 602013070203

Country of ref document: DE

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

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: 20200624

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: 20200624

PLBE No opposition filed within time limit

Free format text: ORIGINAL CODE: 0009261

REG Reference to a national code

Ref country code: CH

Ref legal event code: PL

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

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

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

Effective date: 20200924

26N No opposition filed

Effective date: 20210325

REG Reference to a national code

Ref country code: BE

Ref legal event code: MM

Effective date: 20200930

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: 20200911

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: 20200930

Ref country code: CH

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

Effective date: 20200930

Ref country code: IE

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

Effective date: 20200911

Ref country code: GB

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

Effective date: 20200924

Ref country code: LI

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

Effective date: 20200930

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: 20200624

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

Ref country code: MT

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: 20200624

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: 20200624

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

Ref country code: MK

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: 20200624

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

Ref country code: TR

Payment date: 20230908

Year of fee payment: 11

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

Ref country code: DE

Payment date: 20240730

Year of fee payment: 12

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

Ref country code: FR

Payment date: 20240808

Year of fee payment: 12