EP2589674A1 - Tôle d'acier laminée à froid à ultrahaute résistance présentant une excellente ductilité et résistance à la rupture différée, et son procédé de production - Google Patents

Tôle d'acier laminée à froid à ultrahaute résistance présentant une excellente ductilité et résistance à la rupture différée, et son procédé de production Download PDF

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EP2589674A1
EP2589674A1 EP11800982.8A EP11800982A EP2589674A1 EP 2589674 A1 EP2589674 A1 EP 2589674A1 EP 11800982 A EP11800982 A EP 11800982A EP 2589674 A1 EP2589674 A1 EP 2589674A1
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Prior art keywords
steel sheet
temperature
rolled steel
strength
phase
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EP11800982.8A
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German (de)
English (en)
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EP2589674A4 (fr
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Masataka Yoshino
Kohei Hasegawa
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JFE Steel Corp
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JFE Steel Corp
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/005—Modifying the physical properties by deformation combined with, or followed by, heat treatment of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • 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
    • C21D8/0421—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 characterised by the working steps
    • C21D8/0436—Cold rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—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
    • C21D8/0447—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 characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—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
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • 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

  • the present invention relates to an ultra-high-strength cold-rolled steel sheet which has an excellent strength-ductility balance and excellent delayed fracture resistance and which is a material suitable for use principally in ultra-high-strength automobile structural parts such as center pillars and door impact beams for automobiles and also relates to a method for manufacturing the same.
  • Patent Literature 1 discloses an example in which a steel sheet assumed to have a tensile strength of 1350 MPa and a tempered martensite single-phase microstructure is obtained by quench and tempering although the percentages of phases are not described therein.
  • the total elongation of the steel sheet is small, 7%. Therefore, it is extremely difficult to manufacture automobile safety parts from the steel sheet by pressing.
  • the martensite single-phase microstructure is probably obtained by quenching and therefore the steel sheet probably has a seriously bad shape. This case needs a step of correcting the shape thereof after annealing and therefore is not preferable in terms of manufacture.
  • Patent Literature 2 discloses a TRIP (Transformation-Induced Plasticity) steel sheet which has high strength and ductility and which is obtained by making use of strain-induced transformation, that is, the transformation of retained austenite into martensite by strain during deformation.
  • this steel sheet contains 0.3% to 2% Al on a mass basis.
  • a large amount of Al causes a problem that casting defects are likely to be caused.
  • isothermal holding needs to be performed at a temperature not lower than the Ms transformation temperature in the course of cooling from the annealing temperature, which results in an increased number of manufacturing steps. Since the change in rate of cooling to the temperature of isothermal holding during operation causes a significant change in material quality, operating conditions needs to be strictly controlled in order to stably manufacture steel sheets with a certain level of quality, which is not preferable in terms of manufacture.
  • Non-Patent Literatures 1 and 2 will be described in Examples.
  • the present invention has been made in view of the foregoing circumstances and has an object to provide an ultra-high-strength cold-rolled steel sheet which has excellent delayed fracture resistance and a tensile strength of 1320 MPa or more and which does not excessively contain a transition metal element, such as V or Mo, causing a significant increase in alloying cost or Al, which may possibly cause casting defects, and to provide a method for manufacturing the ultra-high-strength cold-rolled steel sheet.
  • a transition metal element such as V or Mo
  • a microstructure In order to obtain a conventional ultra-high-strength cold-rolled steel sheet with a tensile strength of 1320 MPa or more, a microstructure needs to be transformed into a martensite single-phase microstructure by quenching. In the case where a microstructure is a martensite single-phase, sufficient ductility cannot be achieved. Even if an attempt is made to increase the ductility by tempering subsequent to quenching, the strength is reduced and the ductility is apt not to be increased so much because of the recovery of a dislocation microstructure in a martensite phase and the coarsening of a carbide, such as Fe 3 C, precipitated in the martensite phase.
  • a carbide such as Fe 3 C
  • TRIP steels have by ma use of the many TRIP steels have been invented by making use of the strain-induced transformation of a retained austenite.
  • a large amount of an alloying element needs to be used to increase the stability of austenite and isothermal holding needs to be precisely performed at a temperature not lower than the Ms transformation temperature in the course of cooling from the annealing temperature, which is not preferable in terms of manufacturing stability and manufacturing costs.
  • hydrogen-trapping sites which cause delayed fracture, are preferably diminished as much as possible. Martensite phases are preferably diminished as much as possible because a large number of dislocations serving as hydrogen-trapping sites are introduced into the martensite phases during crystallographic transformation from austenitic phases.
  • Retained austenite which contributes to an increase in ductility, is known to serve as a hydrogen-trapping site like a dislocation and is present on a grain boundary in the form of a film. Therefore, the penetration of hydrogen into retained austenite may possibly cause grain boundary fracture to reduce delayed fracture resistance. Thus, it is not preferred that a metal microstructure contains retained austenite.
  • the inventors have made intensive studies to solve the above problems. As a result, the inventors have elucidated that the balance between tensile strength and ductility can be controlled in such a manner that a microstructure is converted into a microstructure containing a tempered martensite phase and a ferrite phase and the volume fraction of the tempered martensite phase is varied.
  • the inventors have discovered a technique in which a steel sheet with ultra-high strength is obtained in such a manner that the hardness of the tempered martensite phase and that of the ferrite phase are increased by the addition of C and Si the volume fraction of an untempered martensite phase is reduced.
  • the inventors have found that an ultra-high-strength steel sheet with high ductility can be obtained.
  • the inventors have elucidated that the density of dislocations in a microstructure can be significantly reduced as compared with a martensite single-phase microstructure by precipitating a ferrite phase containing substantially no dislocation in the microstructure and the amount of hydrogen permeating through steel can be significantly reduced by diminishing hydrogen-trapping sites.
  • the inventors have found that delayed fracture resistance can be increased.
  • the inventors have found that in view of manufacturing steps, it is effective the annealing temperature and the course of cooling are appropriately controlled during annealing and cooling subsequent to cold rolling and tempering heat treatment is performed at a temperature of 100°C to 300°C.
  • the present invention is based on the above findings.
  • the scope of the present invention is as described below.
  • a cold-rolled steel sheet according to the present invention has extremely high tensile strength, high ductility, and therefore excellent workability. Parts formed from the cold-rolled steel sheet have resistance to delayed fracture due to hydrogen coming from surroundings, that is, excellent delayed fracture resistance. For example, a tensile strength of 1320 MPa or more, a total elongation of 12% or more, and such delayed fracture resistance that fracture does not occur for 100 hours in a 25°C hydrochloric acid environment with a pH of 3 can be readily achieved. Furthermore, a cold-rolled steel sheet having such excellent properties as described above can be stably manufactured by a method according to the present invention.
  • the following sheet can be stably manufactured: an ultra-high-strength cold-rolled steel sheet which has a tensile strength of 1320 MPa or more and which exhibits excellent workability during forming.
  • Parts formed from the cold-rolled steel sheet by press molding have resistance to delayed fracture due to hydrogen coming from surroundings, that is, excellent delayed fracture resistance.
  • Ultra-high-strength parts such as automobile safety parts including center pillars and impact beams, resistant to delayed fracture can be provided.
  • An ultra-high-strength cold-rolled steel sheet according to the present invention has a specific chemical composition and a microstructure as described below. The chemical composition of the cold-rolled steel sheet is first described.
  • C is an element which stabilizes austenite and which is necessary to ensure the strength of the steel sheet.
  • the content of C is less than 0.15% by mass, it is difficult for a microstructure having a tempered martensite phase and a ferrite phase to stably obtain a tensile strength of 1320 MPa or more.
  • the content of C is more than 0.25% by mass, welded portions and heat-affected zones affected by welding are significantly hardened and therefore weldability is reduced. Therefore, the content of C is preferably 0.15% to 0.25% by mass and more preferably 0.18% to 0.22% by mass.
  • Si is a substitutional solid solution hardening element effective in hardening the steel sheet.
  • the content of Si needs to be 1.0% by mass or more.
  • the content of Si is more than 3.0% by mass, scales are significantly formed during hot rolling and the failure rate of final products is increased, which is not economically preferred. Therefore, the content of Si is 1.0% to 3.0% by mass.
  • Mn is an element which stabilizes austenite and which is effective in hardening steel.
  • the content of Mn is less than 1.5% by mass, it is difficult to stably manufacture the steel sheet having a target strength because the hardenability of steel is insufficient, the production of a ferrite phase during cooling from the annealing temperature and the formation of pearlite and bainite begin early, and the strength is significantly reduced.
  • the content thereof is more than 2.5% by mass, segregation is serious, workability is deteriorated in some cases, and delayed fracture resistance is reduced. Therefore, the content of Mn is preferably 1.5% to 2.5% by mass and more preferably 1.5% to 2.0% by mass.
  • P is an element conductive to grain boundary fracture due to grain boundary segregation and therefore is preferably low.
  • the upper limit thereof is 0.05% by mass and is preferably 0.010% by mass. In view of an increase in weldability, the upper limit thereof is more preferably 0.008% by mass or less.
  • S forms an inclusion, such as MnS, causing a reduction in impact resistance and/or delayed fracture resistance and is preferably minimized.
  • the upper limit thereof is 0.02% by mass and preferably 0.002% by mass.
  • A1 is an element effective in deoxidization.
  • the content thereof needs to be 0.01% by mass or more.
  • the content of Al is 0.01% to 0.05% by mass.
  • the content of N is 0.005% by mass or more, the formation of nitrides causes a reduction in ductility at high temperature and low temperature. Therefore, the content of N is less than 0.005% by mass.
  • the steel sheet may further contain one or more of Nb, Ti, and B as required. The effect of the addition of these three elements and the preferred content thereof are described below.
  • Nb and Ti are elements which have a grain-refining effect and which are effective in increasing the strength of the steel sheet; hence, the content of is preferably 0.015% by mass or more. However, when the content of each of Nb and Ti is more than 0.1% by mass, the effect thereof is saturated, which is not economically preferred. Therefore, the content of each of Nb and Ti is 0.1% by mass or less.
  • B is an element effective in increasing the strength of the steel sheet.
  • the content of B is less than 5 ppm by mass, the strength-increasing effect of B cannot be expected.
  • the content of B is more than 30 ppm by mass, hot workability is reduced, which is not preferable in terms of manufacture. Therefore, the content of B is 5 ppm to 30 ppm by mass.
  • the remainders other than the above components are Fe and unavoidable impuritzes.
  • microstructure of the cold-rolled steel sheet is described below.
  • the inventors have made investigations to increase ductility affecting press moldability and investigations to obtain a steel sheet exhibiting excellent delayed fracture resistance after press molding.
  • the inventors have found that the appropriate control of a microstructure is important in exhibiting high ductility.
  • the microstructure contains 40% or more of a tempered martensite phase on a volume fraction basis after continuous annealing, the remainder being a ferrite phase.
  • the microstructure is obtained by quenching from the annealing temperature and tempering subsequent to quenching.
  • an ultra-high-strength cold-rolled steel sheet with high ductility can be obtained without excessively using a transition metal element, such as V or Mo, causing an increase in cost or an alloying element, such as Al, possibly causing casting defects.
  • a transition metal element such as V or Mo
  • an alloying element such as Al
  • An extremely large number of dislocations are introduced into the tempered martensite phase by the crystallographic transformation from an austenite phase to a martensite phase during quenching.
  • the microstructure contains an appropriate amount of the ferrite phase, the number of the dislocations, which serve as hydrogen-trapping sites causing delayed fracture, can be more significantly reduced as compared with a tempered martensite single-phase microstructure and therefore the amount of hydrogen permeating through can be reduced.
  • the tensile strength of steel with a microstructure containing a tempered martensite phase and a ferrite phase increases with an increase in volume fraction of the tempered martensite phase. This is because the hardness of the tempered martensite phase is higher than the hardness of the ferrite phase, the tempered martensite phase, which is a hard phase, exhibits resistance to deformation during tensile deformation, and the larger the volume fraction of the tempered martensite phase is, the more the tensile strength of the steel is close to the tensile strength of the tempered martensite single-phase microstructure.
  • a tensile strength of 1320 MPa or more is not achieved when the volume fraction of the tempered martensite phase is less than 40%. Since ductility decreases with an increase in volume fraction of the tempered martensite phase, a microstructure containing more than 85% of the tempered martensite phase on a volume fraction basis cannot ensure the volume fraction of the ferrite phase that is necessary to achieve a high ductility of 12% or more in terms of total elongation and necessary to increase the delayed fracture resistance. When the volume fraction of the ferrite phase is less than 15%, a high ductility of 12% or more in terms of total elongation is not achieved or an increase in delayed fracture resistance not sufficient. However, when the volume fraction thereof is more than 60%, the volume fraction of the tempered martensite phase that is necessary to achieve a predetermined strength cannot be ensured.
  • the volume fraction of the tempered martensite phase and that of the ferrite phase are 40% to 85% and 15% to 60%, respectively, and more preferably 60% to 85% and 15% to 40%, respectively.
  • the microstructure of the cold-rolled steel sheet according to the present invention may be a two-phase microstructure containing a tempered martensite phase and ferrite phase each having a desired volume fraction and may contain a constituent phase, such as a retained austenite phase, a bainite phase, or a pearlite phase, other than these two phases.
  • the microstructure contains large amounts of the bainite and pearlite phases.
  • the retained austenite phase is principally present at a grain boundary in the form of a film, serves as a hydrogen-trapping site, and therefore may possibly act as an origin of fracture due to hydrogen embrittlement; hence, the content thereof is preferably minimized. Therefore, in the present invention, the volume fraction of the constituent phase (the retained austenite phase, the bainite phase, or the pearlite phase) other than the tempered martensite phase and the ferrite phase is preferably 1% or less in total.
  • the tensile strength and ductility (total elongation as determined by a tensile test using a JIS No. 5 tensile specimen) intended by the present invention are 1320 MPa or more and 12% or more, respectively.
  • the total elongation corresponds to the minimum elongation capable of pressing automobile structural parts such as impact beams. In the present invention, such a strength level and elongation level can be readily achieved.
  • the delayed fracture resistance intended by the present invention is such a performance that fracture does not occur for 100 hours in a 25°C hydrochloric acid environment with a pH of 3. In the present invention, such a performance can be readily achieved.
  • the cold-rolled steel sheet according to the present invention is not particularly limited. Since the cold-rolled steel sheet has the above properties, the cold-rolled steel sheet is particularly suitable for ultra-high-strength automobile safety parts such as automobile door impact beams and center pillars.
  • Steel sheets intended by the present invention include steel strips.
  • the cold-rolled steel sheet according to the present invention may be subjected to surface treatment such as plating (electroplating or the like) or chemical conversion so as to be used as a surface-treated steel sheet.
  • steel with the above composition is produced and is then continuously cast into a cast slab (slab). After being heated to 1200°C or higher, the slab is hot-rolled at a finish rolling end temperature of 800°C or higher. Reasons for limiting hot rolling are described below.
  • the heating temperature of the slab When the heating temperature of the slab is lower than 1200°C, an increase in rolling load increases the risk of causing troubles during hot rolling.
  • the heating temperature of the slab is 1200°C or higher.
  • the heating temperature thereof When the heating temperature thereof is excessively high, an increase in oxidation causes an increase in scale loss.
  • the heating temperature of the slab is preferably 1300°C or lower.
  • the finish rolling end temperature is 800°C or higher, a uniform hot-rolled microstructure can be obtained.
  • the finish rolling end temperature is lower than 800°C, the microstructure of the steel sheet is nonuniform, the ductility thereof is, and the risk of causing various failures during molding is increased.
  • the finish rolling end temperature is 800°C or higher.
  • the upper limit of the finish rolling end temperature is not particularly limited and is preferably 1000°C or lower because rolling at excessively high temperature causes scale defects.
  • the hot-rolled steel sheet is coiled.
  • the coiling temperature thereof is not particularly limited. When the coiling temperature thereof is excessively high, the microstructure of the steel sheet is nonuniform and the ductility thereof is low, due to formation of coarse grains. When the coiling temperature thereof is excessively low, a deformed microstructure caused by hot rolling remains to increase the rolling load in cold rolling subsequent to hot rolling. Therefore, the coiling temperature thereof is preferably 600°C to 700°C. In particular, the coiling temperature thereof is preferably 600°C to 650°C.
  • the hot-rolled steel sheet is pickled, is cold-rolled, is continuously annealed, and is then tempered.
  • Pickling and cold rolling conditions are not particularly limited.
  • the steel sheet is continuously annealed in such a manner that the steel sheet is held at a temperature ranging from the Ac 1 transformation temperature to Ac 3 transformation temperature thereof for 30 s to 1200 s, is cooled to a temperature of 600°C to 800°C at an average cooling rate of 100 °C/s or less, and is then cooled to 100°C or lower at an average cooling rate of 100 °C/s to 1000 °C/s.
  • the steel sheet is subsequently tempered in such a manner that the steel sheet is reheated and is held at a temperature of 100°C to 300°C for 120 s to 1800 s.
  • Reasons for limiting continuous annealing and tempering conditions are described below.
  • the annealing temperature When the annealing temperature is lower than the Ac 1 transformation temperature, an austenite phase (transformed into a martensite phase after quenching) necessary to ensure a predetermined strength is not produced during annealing and therefore such a predetermined strength cannot be achieved even if quenching is performed subsequently to annealing. Even if the annealing temperature is higher than the Ac 3 transformation temperature, 40% or more of the martensite phase can be obtained on a volume fraction basis by controlling a ferrite phase precipitated during cooling from the annealing temperature. In the case of performing annealing at a temperature higher than the Ac 3 transformation temperature, a desired microstructure is unlikely to be obtained. Therefore, the annealing temperature ranges from the Ac 1 transformation temperature to the Ac 3 transformation temperature.
  • the holding time (annealing time) at the annealing temperature is excessively short, a microstructure is not sufficiently annealed, a nonuniform microstructure in which a deformed microstructure caused by hot rolling is present is caused, and the ductility is reduced.
  • the holding time is 30 seconds to 1200 seconds. In particular, the holding time is preferably 250 seconds to 600 seconds.
  • the steel sheet is cooled (the term “cool” is hereinafter referred to as “anneal” in some cases) to a temperature (annealing end temperature) of 600°C to 800°C from the annealing temperature at an average cooling rate of 100 °C/s or less.
  • annealing end temperature 600°C to 800°C from the annealing temperature at an average cooling rate of 100 °C/s or less.
  • the ferrite phase is precipitated during annealing from the annealing temperature and the strength-ductility balance can be thereby controlled.
  • the annealing end temperature is lower than 600°C, a large amount of pearlite is formed in the microstructure to cause a significant reduction in strength and therefore a tensile strength of 1320 MPa cannot be achieved.
  • the annealing end temperature is 600°C to 800°C.
  • the annealing end temperature is preferably 700°C to 750°C.
  • the average annealing rate during annealing is more than 100 °C/s, a sufficient amount of the ferrite phase is not precipitated and therefore predetermined ductility cannot be achieved.
  • the ductility of the microstructure which contains the tempered martensite phase and the ferrite phase as intended by the present invention, results from high work hardenability developed by the coexistence of the tempered martensite phase, which is hard, and the ferrite phase, which is soft.
  • the average annealing rate is more than 100 °C/s, the concentration of carbon in the austenite phase during annealing is insufficient and therefore a hard martensite phase cannot be obtained during quenching.
  • the average annealing rate during annealing is 100 °C/s or less.
  • the average annealing rate is preferably 5 °C/s or less.
  • the steel sheet is cooled (the term “cool” is hereinafter referred to as “quench” in some cases) to a temperature (cooling end temperature) of 100°C or lower at an average cooling rate of 100 °C/s to 1000 °C/s. Quenching subsequent to annealing is performed for the purpose of transforming the austenite phase into the martensite phase.
  • the average cooling rate is less than 100 °C/s, the austenite phase is transformed into the ferrite phase, a bainite phase, or a pearlite phase during cooling and therefore a predetermined strength cannot be achieved.
  • the average cooling rate during quenching is 100 °C/s to 1000 °C/s.
  • the steel sheet is preferably cooled by water quenching.
  • the cooling end temperature is preferably 100°C or lower.
  • the volume fraction of the martensite phase is reduced because of the insufficient transformation of austenite phase into martensite phase during quenching and a reduction in material strength is caused by the self-tempering of the martensite phase produced by quenching, which is not preferable in terms of manufacture.
  • the steel sheet is tempered for the purpose of tempering the martensite phase in such a manner that the steel sheet is reheated and is then held at a temperature of 100°C to 300°C for 120 seconds to 1800 seconds.
  • the tempering thereof softens the martensite phase to increase the workability.
  • the softening of martensite is insufficient and therefore the effect of increasing the workability cannot be expected.
  • Performing tempering at higher than 300°C increases manufacturing costs for reheating, causes a significant reduction in strength, and is incapable of achieving a useful effect.
  • the holding time is less than 120 s, martensite phase is not sufficiently softened at a holding temperature and therefore the effect of increasing the workability cannot be expected.
  • the holding time is more than 1800 s, the strength is significantly reduced because of the excessive softening of martensite phase and manufacturing costs are increased because of an increase in reheating time, which is not preferable.
  • the ultra-high-strength cold-rolled steel sheet according to the present invention can be manufactured through the above manufacturing steps. Since the ultra-high-strength cold-rolled steel sheet according to the present invention has excellent shapeability (flatness) after annealing, a step of correcting the shape of the steel sheet by rolling, leveling, or the like is not necessarily needed. In view of adjusting the quality and/or surface roughness thereof, the annealed steel sheet may be rolled with an elongation of several percent.
  • Test Steels A to M with compositions shown in Table 1 were produced in a vacuum and were then formed into slabs, which were hot-rolled under conditions shown in Table 2, whereby hot-rolled steel sheets with a thickness of 3.4 mm were prepared.
  • the hot-rolled steel sheets were surface-descaled by pickling and were then cold-rolled to a thickness of 1.4 mm.
  • the cold-rolled steel sheets were continuously annealed and tempered under conditions shown in Table 2.
  • the Ac 1 transformation temperature and Ac 3 transformation temperature of each steels is determined from relational equations (the following two equations) described in Non-Patent Literatures 1 and 2, the equations being involved in the dependence of transformation temperature on alloying components:
  • Ac 1 °C 723 - 10.7 ⁇ % by mass Mn + 29.1 ⁇ % by mass Si
  • Ac 3 °C 910 - 230 ⁇ % by mass C 1 / 2 + 29.1 ⁇ % by mass Si - 30 ⁇ % by mass Mn + 700 ⁇ % by mass P + 400 ⁇ % by mass Al + 400 ⁇ % by mass Ti
  • Specimens were taken from the obtained cold-rolled steel sheets. A surface of each specimen that was parallel to the rolling direction was mirror-polished and was etched with nital. The microstructure thereof was observed and photographed with an optical microscope or a scanning electron microscope, whereby the type of a constituent phase such as a tempered martensite phase or a ferrite phase was identify. A photograph of the microstructure was binarized, whereby the volume fraction of each of the tempered martensite phase and the ferrite phase was determined. Since there was a possibility that a retained austenite phase was present in the obtained cold-rolled steel sheets, attempts were made to measure examples of the present invention for retained austenite phase by X-ray (Mo-K ⁇ ) determination. However, the amount of the retained austenite phase present therein was substantially zero and therefore was not included in the remainder shown in Table 3.
  • JIS No. 5 tensile specimens were taken from the obtained cold-rolled steel sheets in a direction perpendicular to the rolling direction and were subjected to a tensile test according to JIS Z 2241, whereby the specimens were determined for tensile property (0.2% proof stress (YS)), tensile strength (TS), and total elongation (EL).
  • YS proof stress
  • TS tensile strength
  • EL total elongation
  • a specimen with a size of 30 mm ⁇ 100 mm was cut out of each of the obtained cold-rolled steel sheets such that the longitudinal direction of the specimen corresponded to the rolling direction of the cold-rolled steel sheets.
  • An end surface of the specimen was ground.
  • the specimen was bent to 180 degrees using a punch having a tip with a radius of curvature of 10 mm.
  • the springback caused in the bent specimen was retained with a bolt 2 such that the distance between inner portions of the specimen 1 was 20 mm.
  • the specimen 1 was immersed in hydrochloric acid with a pH of 3 at 25°C and was measured for up to 100 hours until the specimen 1 was broken. A specimen that was not broken within 100 hours was judged to be acceptable.
  • Tables 1 to 3 confirm that examples of the present invention meet requirements specified herein and have a tensile strength of 1320 MPa or more, a total elongation of 12% or more, a high strength-ductility balance, and excellent delayed fracture resistance because the examples were not broken for 100 hours in the delayed fracture characterization test.
  • Example Nos. 1 to 23 which meet the requirements specified herein, were not broken for 100 hours in the delayed fracture characterization test. This confirms that a cold-rolled steel sheet obtained in accordance with the present invention has sufficient delayed fracture resistance. However, Comparative Example Nos. 25 and 29, each of which the microstructure is a tempered martensite single-phase which is outside the scope of the present invention, were broken within 100 hours and therefore failed in the delayed fracture characterization test.
  • the present invention provides a thin steel sheet for quenching or tempering, the thin steel sheet being suitable for use principally in ultra-high-strength automobile structural parts such as door impact beams and center pillars for automobiles.
  • the composition, rolling conditions, and annealing conditions are appropriately controlled.
  • This allows the steel sheet to have a microstructure containing 40% to 85% of a tempered martensite phase and 15% to 60% of a ferrite phase on a volume fraction basis, a tensile strength of 1320 MPa or more, a total elongation of 12% or more, an excellent strength-ductility balance, and excellent delayed fracture resistance.
  • the use of an ultra-high-strength cold-rolled steel sheet according to the present invention enables the pressing of automobile safety parts such as impact beams.
  • the automobile safety parts exhibit excellent delayed fracture resistance.

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  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Mechanical Engineering (AREA)
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  • Metallurgy (AREA)
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  • Physics & Mathematics (AREA)
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  • Heat Treatment Of Sheet Steel (AREA)
EP11800982.8A 2010-06-30 2011-06-24 Tôle d'acier laminée à froid à ultrahaute résistance présentant une excellente ductilité et résistance à la rupture différée, et son procédé de production Withdrawn EP2589674A4 (fr)

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JP2010148531A JP5668337B2 (ja) 2010-06-30 2010-06-30 延性及び耐遅れ破壊特性に優れる超高強度冷延鋼板およびその製造方法
PCT/JP2011/065135 WO2012002520A1 (fr) 2010-06-30 2011-06-24 Tôle d'acier laminée à froid à ultrahaute résistance présentant une excellente ductilité et résistance à la rupture différée, et son procédé de production

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EP2857539A4 (fr) * 2012-05-31 2016-07-20 Kobe Steel Ltd Plaque d'acier laminé à froid à résistance élevée et son procédé de fabrication
EP3128026A1 (fr) * 2014-03-31 2017-02-08 JFE Steel Corporation Tôle d'acier laminée à froid à grande résistance mécanique présentant une excellente uniformité de la qualité du matériau, et son procédé de production
EP3088547A4 (fr) * 2013-12-27 2017-07-26 Nippon Steel & Sumitomo Metal Corporation Élément en tôle d'acier pressée à chaud, son procédé de production et tôle d'acier pressée à chaud
EP3272892A4 (fr) * 2015-03-18 2018-01-24 JFE Steel Corporation Tôle d'acier laminée à froid à haute résistance et son procédé de fabrication
EP3567132A4 (fr) * 2017-01-05 2019-11-13 JFE Steel Corporation Tôle en acier haute résistance laminée à froid
EP4012055A4 (fr) * 2019-08-06 2022-08-31 JFE Steel Corporation Feuille d'acier mince à haute résistance et son procédé de fabrication
EP4159886A4 (fr) * 2020-05-27 2024-04-17 Baoshan Iron & Steel Acier biphasé à ultra haute résistance et son procédé de fabrication

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CN103805840B (zh) * 2012-11-15 2016-12-21 宝山钢铁股份有限公司 一种高成形性热镀锌超高强度钢板及其制造方法
JP5632947B2 (ja) * 2012-12-12 2014-11-26 株式会社神戸製鋼所 加工性と低温靭性に優れた高強度鋼板およびその製造方法
CN103215505B (zh) * 2013-04-18 2015-08-26 首钢总公司 超高强热连轧带钢及其生产方法
JP6194526B2 (ja) * 2013-06-05 2017-09-13 高周波熱錬株式会社 板状ワークの加熱方法及び加熱装置並びにホットプレス成形方法
JP5888471B1 (ja) * 2014-03-31 2016-03-22 Ｊｆｅスチール株式会社 高降伏比高強度冷延鋼板及びその製造方法
CN106164313B (zh) * 2014-03-31 2018-06-08 杰富意钢铁株式会社 高屈强比高强度冷轧钢板及其制造方法
JP5991450B1 (ja) 2014-12-12 2016-09-14 Ｊｆｅスチール株式会社 高強度冷延鋼板及びその製造方法
CN107250406B (zh) * 2015-02-27 2019-11-15 杰富意钢铁株式会社 高强度冷轧钢板及其制造方法
JP6554397B2 (ja) * 2015-03-31 2019-07-31 株式会社神戸製鋼所 加工性および衝突特性に優れた引張強度が９８０ＭＰａ以上の高強度冷延鋼板、およびその製造方法
WO2016158160A1 (fr) * 2015-03-31 2016-10-06 株式会社神戸製鋼所 TÔLE D'ACIER LAMINÉE À FROID À HAUTE RÉSISTANCE PRÉSENTANT D'EXCELLENTES CARACTÉRISTIQUES D'APTITUDE AU FAÇONNAGE ET DE COLLISION ET PRÉSENTANT UNE RÉSISTANCE À LA TRACTION SUPÉRIEURE OU ÉGALE À 980 MPa, ET SON PROCÉDÉ DE PRODUCTION
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EP3187614A1 (fr) * 2012-05-31 2017-07-05 Kabushiki Kaisha Kobe Seiko Sho (Kobe Steel, Ltd.) Feuille d'acier haute résistance laminée à froid et son procédé de fabrication
EP2857539A4 (fr) * 2012-05-31 2016-07-20 Kobe Steel Ltd Plaque d'acier laminé à froid à résistance élevée et son procédé de fabrication
US10253387B2 (en) 2013-12-27 2019-04-09 Nippon Steel & Sumitomo Metal Corporation Hot-pressed steel sheet member, method of manufacturing the same, and steel sheet for hot pressing
EP3088547A4 (fr) * 2013-12-27 2017-07-26 Nippon Steel & Sumitomo Metal Corporation Élément en tôle d'acier pressée à chaud, son procédé de production et tôle d'acier pressée à chaud
US10711322B2 (en) 2013-12-27 2020-07-14 Nippon Steel Corporation Hot-pressed steel sheet member, method of manufacturing the same, and steel sheet for hot pressing
EP3128026A4 (fr) * 2014-03-31 2017-04-05 JFE Steel Corporation Tôle d'acier laminée à froid à grande résistance mécanique présentant une excellente uniformité de la qualité du matériau, et son procédé de production
EP3128026A1 (fr) * 2014-03-31 2017-02-08 JFE Steel Corporation Tôle d'acier laminée à froid à grande résistance mécanique présentant une excellente uniformité de la qualité du matériau, et son procédé de production
US10329636B2 (en) 2014-03-31 2019-06-25 Jfe Steel Corporation High-strength cold-rolled steel sheet with excellent material homogeneity and production method therefor
EP3272892A4 (fr) * 2015-03-18 2018-01-24 JFE Steel Corporation Tôle d'acier laminée à froid à haute résistance et son procédé de fabrication
EP3567132A4 (fr) * 2017-01-05 2019-11-13 JFE Steel Corporation Tôle en acier haute résistance laminée à froid
US11293103B2 (en) 2017-01-05 2022-04-05 Jfe Steel Corporation High-strength cold-rolled steel sheet
EP4012055A4 (fr) * 2019-08-06 2022-08-31 JFE Steel Corporation Feuille d'acier mince à haute résistance et son procédé de fabrication
EP4159886A4 (fr) * 2020-05-27 2024-04-17 Baoshan Iron & Steel Acier biphasé à ultra haute résistance et son procédé de fabrication

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JP2012012642A (ja) 2012-01-19
JP5668337B2 (ja) 2015-02-12
KR101540507B1 (ko) 2015-07-29
KR20130037208A (ko) 2013-04-15
CN102971442A (zh) 2013-03-13
US20130087257A1 (en) 2013-04-11
EP2589674A4 (fr) 2017-06-21
WO2012002520A1 (fr) 2012-01-05

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