WO2012002520A1 - 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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WO2012002520A1
WO2012002520A1 PCT/JP2011/065135 JP2011065135W WO2012002520A1 WO 2012002520 A1 WO2012002520 A1 WO 2012002520A1 JP 2011065135 W JP2011065135 W JP 2011065135W WO 2012002520 A1 WO2012002520 A1 WO 2012002520A1
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steel sheet
phase
fracture resistance
strength
less
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PCT/JP2011/065135
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English (en)
Japanese (ja)
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正崇 吉野
長谷川　浩平
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Ｊｆｅスチール株式会社
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Priority to US13/805,144 priority Critical patent/US20130087257A1/en
Priority to CN2011800326392A priority patent/CN102971442A/zh
Priority to KR1020127034013A priority patent/KR101540507B1/ko
Priority to EP11800982.8A priority patent/EP2589674A4/fr
Publication of WO2012002520A1 publication Critical patent/WO2012002520A1/fr

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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
    • 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
    • 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
    • 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 excellent in strength / ductility balance and delayed fracture resistance, which is suitable as a material for ultra-high-strength car body structural members such as automobile center pillars and door impact beams, and a method for producing the same. It is.
  • Patent Document 1 Although there is no description regarding the composition ratio of the metal structure as an example of the invention, there is a steel plate that has a tempered martensite single-phase structure having a tensile strength of 1350 MPa by a quenching / tempering method. It is disclosed. However, the elongation at break of the steel sheet is as low as 7%, and it is extremely difficult to produce an automobile security member by pressing. Furthermore, it is surmised that the shape of the steel sheet, which is presumed to have a martensite single-phase structure by rapid cooling, is significantly deteriorated. In this case, a shape correction step is required after annealing, which is not preferable in production.
  • Patent Document 2 discloses a TRIP type (Transformation Induced Plasticity) steel sheet having a high strength but high ductility, using a strain-induced transformation in which retained austenite transforms into martensite due to strain during processing.
  • Al is added in an amount of 0.3 to 2% by mass.
  • it is necessary to perform isothermal holding at a temperature equal to or higher than the Ms transformation point in the cooling process from the annealing temperature, which increases the number of manufacturing steps.
  • the cooling rate to the isothermal holding temperature fluctuates during operation, it will cause large material fluctuations, so in order to stably produce a steel plate of a certain quality, it is necessary to strictly manage the operating conditions, It is not preferable in production.
  • Non-patent document 1 and non-patent document 2 will be described in Examples.
  • the present invention has been made in view of such circumstances, and is a steel component that does not contain excessive amounts of transition metal elements such as V and Mo that significantly increase alloy costs and Al that may induce casting defects.
  • An object of the present invention is to provide an ultra-high strength cold-rolled steel sheet having excellent delayed fracture resistance and having a tensile strength of 1320 MPa or more and a method for producing the same.
  • the martensite phase is a dislocation that becomes a hydrogen trap site during crystal structure transformation from the austenite phase. Is introduced in large quantities, it is desirable to reduce it as much as possible.
  • residual austenite which contributes to improving ductility, is known to act as a hydrogen trap site, as is the case with dislocations, and since residual austenite exists in the form of a film on the grain boundary, hydrogen penetrates into the residual austenite. It is not preferable that residual austenite is included in the metal structure because it induces grain boundary fracture and may deteriorate the delayed fracture resistance.
  • the microstructure is a structure having a tempered martensite phase and a ferrite phase, and by changing the volume ratio of the tempered martensite phase, the tensile strength and It becomes clear that the balance of ductility can be controlled, and by adding C and Si, the hardness of the tempered martensite phase and the ferrite phase is increased, and the volume ratio of the tempered martensite phase is reduced, thereby reducing the volume ratio of the steel sheet.
  • a technique for achieving ultra-high strength has been found, and it has been found that it is possible to obtain a steel sheet having high ductility while having extremely high strength.
  • the dislocation density in the metal structure is greatly reduced compared to the martensite single phase structure, and by reducing the hydrogen trap sites, It was clarified that the amount of hydrogen penetrating into steel can be greatly reduced, and it was found that the delayed fracture resistance can be improved.
  • the annealing temperature and the subsequent cooling process are appropriately controlled in the annealing and cooling after the cold rolling, and then the tempering heat treatment is performed in the temperature range of 100 ° C. or more and 300 ° C. or less. The knowledge that it is effective was obtained.
  • the present invention is based on the above findings. That is, the gist configuration of the present invention is as follows.
  • the ultra-high-strength cold-rolled steel sheet of [1] further contains one or more of Nb: 0.1% or less, Ti: 0.1% or less, and B: 5 to 30 ppm by mass ratio.
  • Nb 0.1% or less
  • Ti 0.1% or less
  • B 5 to 30 ppm by mass ratio.
  • a steel slab having the chemical composition described in [1] or [2] above is heated to 1200 ° C. or higher and hot-rolled at a finish rolling exit temperature of 800 ° C. or higher, and then pickled and cold. rolled, then the time of continuous annealing, after 30 ⁇ 1200 sec maintained at a temperature in the range of a C1 transformation point ⁇ a C3 transformation point, cooling at an average cooling rate of 100 ° C.
  • the cold-rolled steel sheet of the present invention has an extremely high tensile strength, a high ductility and an excellent workability associated therewith. In addition, it has excellent delayed fracture resistance that hardly causes delayed fracture due to hydrogen entering from the environment even after being molded into the member. For example, it has a tensile strength of 1320 MPa or more, an elongation at break of 12% or more, and it can easily realize delayed fracture resistance that does not break for 100 hours in a hydrochloric acid environment at 25 ° C. and pH 3. Furthermore, according to the manufacturing method of the present invention, the cold-rolled steel sheet having the excellent performance as described above can be stably manufactured.
  • the material has excellent delayed fracture resistance that hardly causes delayed fracture due to hydrogen penetrating from the environment even after press molding to the member, and has a tensile strength of 1320 MPa or more that exhibits excellent workability during molding.
  • An ultra-high strength cold-rolled steel sheet can be stably produced, and an ultra-high-strength member that hardly causes delayed fracture, for example, an automobile security member such as a center pillar or an impact beam can be provided.
  • the ultra-high strength cold-rolled steel sheet of the present invention has a specific chemical component and metal structure as described below. First, chemical components of the cold rolled steel sheet will be described.
  • C 0.15 to 0.25% by mass>
  • C is an element that stabilizes austenite and is an element necessary to ensure the strength of the steel sheet. If the amount of C is less than 0.15% by mass, it is difficult to stably obtain a tensile strength of 1320 MPa or more in a structure having a tempered martensite phase and a ferrite phase. On the other hand, if the amount of C exceeds 0.25% by mass, the welded part and the heat-affected zone due to welding are markedly cured, and the weldability is lowered. For this reason, the C content is in the range of 0.15 to 0.25% by mass. More preferably, it is in the range of 0.18 to 0.22 mass%.
  • Si 1.0 to 3.0% by mass> Si is a substitutional solid solution strengthening element effective for hardening a steel plate. In order to exhibit this effect, it is necessary to contain 1.0 mass% or more. When the amount of Si exceeds 3.0% by mass, scale formation by hot rolling becomes remarkable, and the defect rate in the final product increases, which is economically undesirable. Therefore, the Si amount is set to 1.0 to 3.0% by mass.
  • Mn stabilizes austenite and is an element effective for strengthening steel.
  • Mn is less than 1.5% by mass, the hardenability of the steel is not sufficient, and the formation of ferrite phase and the formation of pearlite and bainite during cooling from the annealing temperature starts early, and the strength is significantly reduced. Therefore, it becomes difficult to stably manufacture a steel plate having the target strength.
  • Mn is 1.5 to 2.5 mass%, preferably 1.5 to 2.0 mass%.
  • P 0.05% by mass or less>
  • P is an element that promotes grain boundary fracture due to grain boundary segregation, so a lower value is desirable, and its upper limit is 0.05 mass%, preferably 0.010 mass%. Furthermore, 0.008 mass% or less is more preferable from a viewpoint of weldability improvement.
  • S 0.02 mass% or less> S is an inclusion such as MnS and induces deterioration of impact resistance and delayed fracture resistance. Therefore, the lower limit is desirable, and the upper limit is set to 0.02 mass%, preferably 0.002 mass%. .
  • Al 0.01 to 0.05% by mass>
  • Al is an element effective for deoxidation, and in order to obtain a useful deoxidation effect, it is necessary to make it 0.01% by mass or more. Inclusions increase and ductility decreases. Therefore, the Al amount is set to 0.01 to 0.05% by mass.
  • ⁇ N less than 0.005% by mass>
  • the content of N is 0.005% by mass or more, ductility at high and low temperatures due to the formation of nitrides decreases. Therefore, the N amount is less than 0.005% by mass.
  • the steel plate can further contain one or more of Nb, Ti, and B as required.
  • Nb, Ti, and B as required.
  • Nb and Ti have an effect of refining crystal grains and are effective elements for increasing the strength of the steel sheet. Therefore, it is preferable to add 0.015% by mass or more. However, even if Nb and Ti are contained in amounts exceeding 0.1% by mass, the effect is saturated, which is not economically preferable. For this reason, the addition amounts of Nb and Ti are each 0.1% by mass or less.
  • B is an element effective for increasing the strength of the steel sheet. If the amount of B is less than 5 ppm by mass, the effect of increasing the strength due to B cannot be expected. On the other hand, if the amount of B exceeds 30 ppm by mass, the hot workability is lowered, which is not preferable for production. For this reason, the addition amount of B is set to 5 to 30 ppm by mass.
  • the remainder other than the above is Fe and inevitable impurities.
  • the present inventors have studied the microstructure, It was found that proper control is important. Specifically, it has been found that it is important that the microstructure after continuous annealing contains a tempered martensite phase in a volume ratio of 40% or more and the balance has a ferrite phase. This structure is obtained by rapid cooling from the annealing temperature during annealing and tempering after the rapid cooling. According to this technique, there is a possibility of inducing transition metal elements such as V and Mo and casting defects that increase costs. It is possible to obtain an ultra-high strength cold-rolled steel sheet having high ductility without excessively adding an alloy element such as Al.
  • the delayed fracture resistance is better as the amount of hydrogen entering the steel is smaller.
  • an extremely large amount of dislocation is introduced by the crystal structure transformation from the austenite phase to the martensite phase during quenching, but it is said that delayed fracture is induced by including an appropriate amount of ferrite phase in the metal structure.
  • dislocations that serve as hydrogen trap sites can be greatly reduced, and the amount of hydrogen intrusion into the steel can be reduced.
  • the tensile strength of steel having a structure having a tempered martensite phase and a ferrite phase increases as the volume ratio of the tempered martensite phase increases. This is because the tempered martensite phase and the ferrite phase have higher hardness in the tempered martensite phase, and the deformation resistance during tensile deformation is borne by the tempered martensite phase, which is the hard phase, and the volume ratio of the tempered martensite phase. This is because the larger the value, the closer to the tensile strength of the tempered martensite single phase structure. In the steel component range of the present invention, if the tempered martensite volume fraction is less than 40%, a tensile strength of 1320 MPa or more cannot be obtained.
  • the ductility decreases, and in the structure where the volume fraction of the tempered martensite phase exceeds 85%, in order to improve the high ductility and delayed fracture resistance of 12% or more at break elongation.
  • the necessary ferrite phase cannot be secured.
  • the volume fraction of the ferrite phase is less than 15%, the ductility and delayed fracture resistance are not sufficiently improved by 12% or more in elongation at break, whereas when it exceeds 60%, tempering necessary for obtaining a predetermined strength is achieved.
  • the volume ratio of the martensite phase cannot be secured.
  • the metal structure of the cold-rolled steel sheet of the present invention has a tempered martensite phase volume ratio of 40 to 85% and a ferrite phase volume ratio of 15 to 60%. More preferably, it is a metal structure in which the volume ratio of the tempered martensite phase is 60 to 85% and the volume ratio of the ferrite phase is 15 to 40%.
  • the metal structure of the cold-rolled steel sheet of the present invention may be a two-phase structure composed of a tempered martensite phase and a ferrite phase having a desired volume ratio, and as a structure other than these two phases, a residual austenite phase, a bainite phase Further, a constituent phase such as a pearlite phase may be included.
  • the constituent phases (bainite phase, pearlite phase, residual austenite phase, etc.) other than the tempered martensite phase and the ferrite phase are preferably 1% or less in total in volume ratio.
  • the target tensile strength and ductility of the present invention are a tensile strength of 1320 MPa or more and a breaking elongation (breaking elongation in a tensile test using a JIS No. 5 tensile test piece) of 12% or more. Although this corresponds to the minimum ductility that can be pressed into a member, such strength and ductility levels can be easily realized in the present invention.
  • the delayed fracture resistance targeted by the present invention is a performance that does not cause a fracture for 100 hours in a hydrochloric acid environment at 25 ° C. and pH 3, such a performance can be easily realized in the present invention.
  • the cold-rolled steel sheet of the present invention has the above-described performance, and thus is particularly suitable for an ultra-high-strength body security member including an automobile door impact beam and a center pillar.
  • a steel strip is also included in the steel sheet targeted by the present invention, and the cold-rolled steel sheet of the present invention is subjected to surface treatment such as plating (electroplating) or chemical conversion treatment on the surface and used as a surface-treated steel sheet. You can also.
  • a steel having the above composition is melted to form a slab by continuous casting, and the slab is heated to 1200 ° C. or higher and hot rolled at a finish rolling exit temperature of 800 ° C. or higher.
  • the reason for limiting hot rolling will be described.
  • the slab heating temperature is set to 1200 ° C. or higher. If the heating temperature is too high, the slab heating temperature is desirably 1300 ° C. or less because it leads to an increase in scale loss accompanying an increase in oxidized weight.
  • ⁇ Finishing rolling exit temperature 800 ° C or higher> By setting the finish rolling exit temperature to 800 ° C. or higher, a uniform hot rolled matrix phase structure can be obtained. If the finish rolling exit temperature is lower than 800 ° C., the structure of the steel sheet becomes non-uniform, the ductility is lowered, and the risk of generating various problems during forming increases. Accordingly, the finish rolling exit temperature is set to 800 ° C. or higher.
  • the upper limit of the finish rolling outlet temperature is not particularly limited, but if it is rolled at an excessively high temperature, it causes scale wrinkles and the like, and is preferably about 1000 ° C. or less.
  • the coiling temperature is not particularly limited, but if the coiling temperature is too high, coarse grains are generated, the steel sheet structure becomes non-uniform, and the ductility decreases. In addition, when the coiling temperature is too low, the processed structure generated by hot rolling remains, and the rolling load of cold rolling, which is the next process, becomes large. Therefore, the winding temperature is preferably 600 to 700 ° C. A particularly preferable winding temperature is 600 to 650 ° C.
  • pickling and cold rolling are performed, followed by continuous annealing and tempering.
  • the conditions for pickling and cold rolling are not particularly limited. Continuous annealing, after 30 ⁇ 1200 sec maintained at a temperature in the range of A C1 transformation point ⁇ A C3 transformation point, 100 ° C. / sec is cooled to 600 ⁇ 800 ° C. or less of the average cooling rate subsequent 100 ⁇ 1000 ° C. / sec It cools to 100 degrees C or less with an average cooling rate, Then, it reheats and performs the tempering process hold
  • the reasons for limiting the conditions for continuous annealing and tempering will be described.
  • the annealing temperature is in the range of A C1 transformation point to A C3 transformation point. Further, from the viewpoint of stably ensuring that the equilibrium volume fraction of the austenite phase is 40% or more in this temperature range, it is preferably 760 ° C. or more, and more preferably 780 ° C. or more.
  • the holding time at the annealing temperature annealing time
  • the microstructure is not sufficiently annealed and becomes a non-uniform structure in which a cold-rolled processed structure exists, and ductility is lowered.
  • the holding time is 30 to 1200 seconds.
  • a particularly preferable holding time is in the range of 250 to 600 seconds.
  • the slow cooling stop temperature is set to 600 to 800 ° C.
  • the slow cooling stop temperature is preferably 700 to 750 ° C. from the viewpoint of suppressing material fluctuations accompanying the slow cooling stop temperature fluctuation in operation.
  • the average cooling rate of slow cooling exceeds 100 ° C./second, a sufficient amount of ferrite phase does not precipitate during slow cooling, so that a predetermined ductility cannot be obtained.
  • the ductility of the metal structure having a tempered martensite phase and a ferrite phase intended in the present invention is also attributed to the high work-hardening ability expressed by the mixture of a hard tempered martensite phase and a soft ferrite phase.
  • the cooling rate exceeds 100 ° C./second, carbon concentration in the austenite during slow cooling becomes insufficient, and a hard martensite phase cannot be obtained during rapid cooling. As a result, the work hardening ability of the final structure is lowered and sufficient ductility cannot be obtained.
  • the average cooling rate during slow cooling is 100 ° C./second or less. In order to sufficiently cause carbon concentration in the austenite phase, an average cooling rate of 5 ° C./second or less is preferable.
  • Rapid cooling is performed at an average cooling rate of 100 to 1000 ° C./second to a temperature of 100 ° C. or lower (cooling stop temperature) (in the following description, this cooling may be referred to as “rapid cooling”). Rapid cooling after slow cooling is performed to transform the austenite phase into the martensite phase, but if the average cooling rate is less than 100 ° C / second, the austenite phase transforms into the ferrite phase, bainite phase or pearlite phase during cooling. The predetermined strength cannot be obtained.
  • the average cooling rate during rapid cooling is set to 100 to 1000 ° C./second. This cooling is preferably rapid cooling by water quenching.
  • the cooling stop temperature is preferably 100 ° C. or lower. Manufactured to induce a decrease in the volume fraction of the martensite phase due to insufficient quenching of austenite during rapid cooling when the cooling stop temperature exceeds 100 ° C, and a decrease in material strength due to self-tempering of the martensite phase generated by rapid cooling. Not preferable.
  • ⁇ Tempering treatment held at 100 to 300 ° C. for 120 to 1800 seconds> Subsequent to the rapid cooling described above, a tempering treatment is performed in order to temper the martensite phase and hold it for 120 to 1800 seconds in a temperature range of 100 to 300 ° C. This tempering softens the martensite phase and improves workability. When tempering is performed at a temperature lower than 100 ° C., the martensite is not sufficiently softened and an effect of improving workability cannot be expected. In addition, performing tempering above 300 ° C. not only increases the manufacturing cost for reheating, but also causes a significant decrease in strength, and a useful effect cannot be obtained.
  • the holding time is less than 120 sec
  • the martensite is not sufficiently softened at the holding temperature, so that the workability improvement effect cannot be expected.
  • the holding time exceeds 1800 sec, the strength is remarkably lowered due to excessive softening of martensite, and the manufacturing cost is increased due to an increase in reheating time, which is not preferable.
  • the ultra-high strength cold-rolled steel sheet of the present invention can be manufactured.
  • the ultra-high-strength cold-rolled steel sheet of the present invention is excellent in plate shape (flatness) after annealing, a process for correcting the shape of the steel sheet, such as rolling and leveler processing, is not necessarily required. From the viewpoint of adjusting the material and surface roughness, there is no problem even if the annealed steel sheet is rolled at an elongation of about several percent.
  • Test steels A to M having the composition shown in Table 1 were melted in vacuum to form a slab, and then hot rolled under the conditions shown in Table 2 to obtain a hot-rolled steel sheet having a thickness of 3.4 mm.
  • the hot-rolled steel sheet was pickled to remove the surface scale, and then cold-rolled to a thickness of 1.4 mm.
  • continuous annealing and tempering treatments were performed under the conditions shown in Table 2.
  • the AC1 transformation point of each steel type is the value calculated
  • a C1 [° C.] 723-10.7 ⁇ (mass% Mn) + 29.1 ⁇ (mass% Si) (1)
  • a C3 [° C.] 910 ⁇ 203 ⁇ (mass% C) 1/2 + 29.1 ⁇ (mass% Si) ⁇ 30 ⁇ (mass% Mn) + 700 ⁇ (mass% P) + 400 ⁇ (mass% Al) +400 X (mass% Ti) (2)
  • Observation (measurement) and performance test of the metal structure were performed as follows.
  • (1) Observation of metal structure A specimen was taken from the obtained cold-rolled steel sheet, mirror-polished with respect to a cross section parallel to the rolling direction, etched with nital, and a fine structure was obtained using an optical microscope or a scanning electron microscope. Observe and photograph, identify the type of constituent phases such as tempered martensite phase and ferrite phase, and binarize the structure picture using an image analysis device, so that the volume ratio of tempered martensite phase and ferrite phase Asked.
  • the obtained cold-rolled steel sheet may have a retained austenite phase. Therefore, the inventive example tried to measure the retained austenite phase by the X-ray (Mo-K ⁇ ray) measurement method. Are almost zero and are not included in the remainder of Table 3.
  • the inventive examples meeting the conditions of the present invention have a high strength / ductility balance of a tensile strength of 1320 MPa or more and a breaking elongation of 12% or more, and fracture occurs for 100 hours in the delayed fracture property evaluation test. It was confirmed that it has excellent delayed fracture resistance.
  • the annealing time was 10 seconds, which was outside the scope of the present invention. 24, the pearlite structure generated after the hot rolling remains after the annealing process, and the effect of working strain accompanying the cold rolling is not sufficiently removed, so that the predetermined strength and ductility are obtained. Absent.
  • the annealing temperature was set to AC3 point or higher. Nos. 25 and 29 cannot cause precipitation of a ferrite phase during slow cooling, and have a martensite single-phase structure. Although a predetermined strength is obtained, a predetermined ductility is not obtained. The steel component is outside the scope of the present invention. Nos.
  • Invention Example No. adapted to the conditions of the present invention
  • the delayed fracture property evaluation test did not cause fracture for 100 hours, and it was confirmed that the cold-rolled steel sheet obtained by the present invention has sufficient delayed fracture resistance.
  • the metallographic structure is a single phase of tempered martensite, which is outside the scope of the present invention. In Nos. 25 and 29, cracks occurred within 100 hours, so the delayed fracture resistance test result failed.
  • the present invention is a thin steel plate for quenching and tempering suitable for the use of an ultra-high-strength car body safety member such as an automobile door impact beam and a center pillar, and an automotive part using such a steel plate.
  • an ultra-high-strength car body safety member such as an automobile door impact beam and a center pillar
  • an automotive part using such a steel plate In manufacturing, by appropriately controlling the steel composition, rolling conditions, and annealing conditions, a structure containing a tempered martensite phase of 40% to 85% by volume and a ferrite phase of 15% to 60% by volume. It has a tensile strength of 1320 MPa or more, an elongation at break of 12% or more and an excellent strength-ductility balance, and also has excellent delayed fracture resistance.
  • the ultra-high-strength cold-rolled steel sheet of the present invention it is possible to press an automobile safety member such as an impact beam, and this automobile safety member exhibits excellent delayed fracture resistance.

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  • Chemical & Material Sciences (AREA)
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  • Mechanical Engineering (AREA)
  • Materials Engineering (AREA)
  • Metallurgy (AREA)
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  • Heat Treatment Of Sheet Steel (AREA)

Abstract

L'invention concerne une tôle d'acier laminée à froid à ultrahaute résistance dont la composition ne contient aucun excès de métal de transition qui augmenterait considérablement le coût de l'alliage comme V et Mo, ni d'aluminium qui peut induire des défauts de coulée, et qui, malgré cette composition, présente une excellente résistance à la rupture différée et une résistance à la traction d'au moins 1320 MPa. L'invention concerne également un procédé de production de ladite tôle d'acier. Ladite tôle d'acier, qui présente une excellente ductilité et résistance à la rupture différée contient, en pourcentage massique, 0,15-0,25 % de C, 1,0-3,0 % de Si, 1,5-2,5 % de Mn, jusqu'à 0,05 % de P, jusqu'à 0,02 % de S, 0,01-0,05 % d'Al et moins de 0,005 % de N, le solde étant du fer et les inévitables impuretés. La tôle d'acier présente une structure métallographique qui comprend 40-85 % en volume de phase martensitique revenue et 15-60 % en volume de phase ferritique, et possède une résistance à la traction d'au moins 1320 MPa.
PCT/JP2011/065135 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 WO2012002520A1 (fr)

Priority Applications (4)

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US13/805,144 US20130087257A1 (en) 2010-06-30 2011-06-24 Ultra high strength cold rolled steel sheet having excellent ductility and delayed fracture resistance and method for manufacturing the same
CN2011800326392A CN102971442A (zh) 2010-06-30 2011-06-24 延展性以及耐延迟断裂特性优良的超高强度冷轧钢板及其制造方法
KR1020127034013A KR101540507B1 (ko) 2010-06-30 2011-06-24 연성 및 내지연 파괴 특성이 우수한 초고강도 냉연 강판 및 그 제조 방법
EP11800982.8A EP2589674A4 (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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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
JP2010-148531 2010-06-30

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WO2016147549A1 (fr) * 2015-03-18 2016-09-22 Ｊｆｅスチール株式会社 Tôle d'acier laminée à froid à haute résistance et son procédé de fabrication
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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JP6697728B1 (ja) * 2018-10-04 2020-05-27 日本製鉄株式会社 冷延鋼板
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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
CN104838027A (zh) * 2012-12-12 2015-08-12 株式会社神户制钢所 高强度钢板及其制造方法
WO2014092025A1 (fr) * 2012-12-12 2014-06-19 株式会社神戸製鋼所 Plaque d'acier à résistance élevée et son procédé de fabrication
US9322088B2 (en) 2012-12-12 2016-04-26 Kabushiki Kaisha Kobe Seiko Sho (Kobe Steel, Ltd.) High-strength steel sheet and method for producing the same
JP2015025208A (ja) * 2012-12-12 2015-02-05 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
JP2014133944A (ja) * 2012-12-12 2014-07-24 Kobe Steel Ltd 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
WO2016147549A1 (fr) * 2015-03-18 2016-09-22 Ｊｆｅスチール株式会社 Tôle d'acier laminée à froid à haute résistance et son procédé de fabrication
JPWO2016147549A1 (ja) * 2015-03-18 2017-07-13 Ｊｆｅスチール株式会社 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
JP2016194139A (ja) * 2015-03-31 2016-11-17 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
CN107429371A (zh) * 2015-03-31 2017-12-01 株式会社神户制钢所 加工性和碰撞特性优异且抗拉强度为980MPa以上的高强度冷轧钢板及其制造方法
CN107429371B (zh) * 2015-03-31 2020-04-21 株式会社神户制钢所 加工性和碰撞特性优异且抗拉强度为980MPa以上的高强度冷轧钢板及其制造方法

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KR101540507B1 (ko) 2015-07-29
EP2589674A1 (fr) 2013-05-08
EP2589674A4 (fr) 2017-06-21
CN102971442A (zh) 2013-03-13

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