High-strength steel sheet and its manufacturing method, high-strength plated steel sheet and its manufacturing method, member and its manufacturing method, and automobile frame structural part or automobile reinforcing part

Abstract

The present invention provides a high-strength steel sheet that is excellent in component strength, ductility, stretch flangeability, and delayed fracture resistance at a stress-loaded shear edge. The high-strength steel sheet of the present invention has a predetermined chemical composition containing one or both of Mo and V that satisfies the following formula (1), and a predetermined steel structure at the 1 / 4 position of the sheet thickness, and the number density of composite carbides containing Ti and one or both of Mo and V is 1.0 × 10 at the 1 / 4 position of the sheet thickness. 19 pieces / m 3 or more, and the average grain size of the prior austenite grains is 3.00 μm or less. 0.0010≦[%Mo] / 95.94+[%V] / 50.94≦0.0200 ···(1) Here, [%Mo] and [%V] are the contents of Mo and V, respectively.

Description

[Technical Field]

[0001] The present invention relates to a high-strength steel sheet and a method for manufacturing the same, a high-strength plated steel sheet and a method for manufacturing the same, a member and a method for manufacturing the same, and an automobile frame structural part or automobile reinforcing part. [Background technology]

[0002] With the aim of reducing CO2 emissions by reducing vehicle weight and improving crashworthiness at the same time, progress is being made in increasing the strength of automotive steel sheets, and new regulations are being introduced one after another. As a result, there has been an increase in the use of high-strength steel sheets with a tensile strength (TS) of 1180 MPa or more in the main structural parts that form the framework of automobile cabins.

[0003] High strength steel sheets used in automotive reinforcement and frame structural parts are often required to have high part strength. In this case, a high yield ratio (YR = yield strength YS / tensile strength TS × 100) is required, but parts such as side sills are processed with sheared edges. Therefore, the steel sheets used for these parts also require good ductility and stretch flangeability.

[0004] Here, a problem with high-strength steel sheets with a tensile strength of 1180 MPa or more is that hydrogen intrusion in the atmospheric corrosive environment in which a vehicle is traveling can cause delayed fracture, a sudden breakage of the component. Automotive steel sheets are subjected to stress during pressing and assembly, and there is a risk that hydrogen from the environment will subsequently infiltrate the steel sheet. Because delayed fracture is likely to occur at the sheared edge, there is a need to improve the delayed fracture resistance of the stressed sheared edge. In order to increase the application rate of high-strength steel sheets in automotive parts, it is necessary to comprehensively satisfy these properties.

[0005] In response to these demands, for example, Patent Document 1 provides a high-strength steel sheet that has excellent delayed fracture resistance and a tensile strength of 1180 MPa or more by increasing the proportion of martensite blocks in which metastable carbides exist. [Prior art documents] [Patent documents]

[0006] [Patent Document 1] International Publication No. 2023 / 026819 Summary of the Invention [Problem to be solved by the invention]

[0007] However, in order to increase the application rate of high-strength steel sheets to automotive parts more than ever before, it is necessary to improve the delayed fracture resistance of the stress-loaded sheared edge portion, rather than the machine-ground edge portion. In Patent Document 1, the delayed fracture resistance of the machine-ground edge portion was evaluated, and the inventors' investigation revealed that there was room for improvement.

[0008] In view of the above problems, the present invention aims to provide a high-strength steel sheet and a high-strength plated steel sheet that are excellent in part strength, ductility, stretch flangeability, and delayed fracture resistance at a stressed shear edge, as well as advantageous methods for producing them. It is also an object of the present invention to provide a member including the high-strength steel sheet or the high-strength plated steel sheet, a method for producing the same, and an automobile frame structural part or automobile reinforcing part that includes the member.

[0009] Here, "high strength" means that the tensile strength (TS) determined by the tensile test described below is 1180 MPa or more.

[0010] "Excellent part strength" means that the yield ratio (YR) determined by the tensile test described below is 70% or more.

[0011] "Excellent ductility" means that the total elongation (El) determined by the tensile test described below is 7.0% or more.

[0012] "Excellent stretch flangeability" means that the limiting hole expansion ratio (λ) determined by the hole expansion test described below is 25% or more.

[0013] The phrase "excellent delayed fracture resistance at the stress-loaded sheared end face" means that a delayed fracture test specimen having an end face sheared at a shear angle of 0.5° and a clearance of 10%, as described below, is subjected to stress loading and immersed for 96 hours in a 0.1% by mass aqueous solution of ammonium thiocyanate at 25°C, the pH of which has been adjusted to 7.0 using McIlvaine buffer solution, and the fracture threshold stress after immersion is 900 MPa or more. [Means for solving the problem]

[0014] As a result of extensive research, the present inventors have found that the above object can be achieved by employing the following configuration, and have completed the present invention.

[0015] That is, the gist and configuration of the present invention are as follows. [1] In mass%, C: 0.070% or more and 0.390% or less, Si: 0.01% or more and 2.00% or less, Mn: 1.70% or more and 4.00% or less, P: 0.100% or less, S: 0.0200% or less, Al: 1.000% or less, N: 0.0100% or less, O: 0.0100% or less, Ti: 0.015% or more and 0.200% or less, and One or both of Mo and V satisfying the following formula (1): and the remainder being Fe and unavoidable impurities; At 1 / 4 of the plate thickness, The total area ratio of martensite and bainite is 70% or more, The area ratio of ferrite is 20% or less, a steel structure in which the area ratio of retained austenite is 15% or less; and At 1 / 4 of the plate thickness, The number density of the composite carbide containing Ti and further containing one or both of Mo and V is 1.0 × 10 19 pieces / m3 That's all, A high-strength steel plate having an average prior austenite grain size of 3.00 μm or less. 0.0010≦[%Mo] / 95.94+[%V] / 50.94≦0.0200 ···(1) Here, [%Mo] and [%V] are the contents of Mo and V, respectively.

[0016] [2] The component composition further includes, in mass%, Nb: 0.200% or less, Ta: 0.10% or less, W: 0.10% or less, B: 0.0100% or less, Cr: 1.00% or less, Ni: 1.00% or less, Co: 1.00% or less, Cu: 1.00% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.100% or less, Te: 0.100% or less, Hf: 0.100% or less, and Bi:0.200% or less The high-strength steel plate according to the above [1], containing at least one element selected from the group consisting of:

[0017] [3] A high-strength plated steel sheet comprising the high-strength steel sheet according to [1] or [2] above and a plating layer located on the surface of the high-strength steel sheet.

[0018] [4] A member comprising the high-strength steel plate according to [1] or [2] above.

[0019] [5] A member comprising the high-strength plated steel sheet according to [3] above.

[0020] [6] An automobile frame structural part or an automobile reinforcement part, comprising the member described in [4] above.

[0021] [7] An automobile frame structural part or an automobile reinforcement part, comprising the member described in [5] above.

[0022] [8] A heating step of heating a steel slab having the composition described in [1] or [2] above to 1150°C or higher; A hot rolling process in which the steel slab is subjected to rough rolling with a rolling reduction of 13% or more in a temperature range of 1050 ° C. or higher and the number of passes is 4 or more, and then subjected to finish rolling to obtain a hot-rolled sheet; Next, a cold rolling step in which the hot-rolled sheet is subjected to pickling and cold rolling to obtain a cold-rolled sheet; Next, an annealing step of annealing the cold-rolled sheet; Next, a cooling step of cooling the cold-rolled sheet to room temperature to obtain a high-strength steel sheet; and In the annealing step, The average heating rate v1 in the temperature range T1 of 600°C or higher and Ac3°C or lower is 1.5°C / s or higher, The maximum temperature of the cold-rolled sheet is in a temperature range T2 of Ac3 ° C or more and (Ac3 + 200) ° C or less, The retention condition from the time when the temperature of the cold-rolled sheet reaches Ac3 ° C. to the time when it reaches 750 ° C. via the maximum temperature reached satisfies the following formula (2), In the cooling step, The average cooling rate v3 in the temperature range T3 of 550°C or more and 750°C or less is 4°C / s or more, A method for producing high-strength steel plate, wherein the residence time t4 in the temperature range T4 of 350°C or higher and 550°C or lower is 600 seconds or shorter. P≦30000 (2) where:

number

number

[0023] [9] A method for producing a high-strength steel plate according to [8] above; In the cooling step, a plating step of performing a plating treatment on the cold-rolled sheet; A method for producing a high-strength plated steel sheet having the above-mentioned steps.

[0024]

[10] A method for manufacturing a component, comprising subjecting the high-strength steel plate according to [1] or [2] above to one or both of forming and joining processes to form the component.

[0025]

[11] A method for manufacturing a component, comprising subjecting the high-strength plated steel sheet according to [3] above to one or both of forming and joining processes to form the component. [Effects of the Invention]

[0026] According to the present invention, it is possible to provide a high-strength steel sheet and a high-strength plated steel sheet which are excellent in part strength, ductility, stretch flangeability, and delayed fracture resistance at a stressed shear edge, as well as an advantageous method for manufacturing them. Furthermore, it is possible to provide a member including a high-strength steel sheet or a high-strength plated steel sheet, a method for manufacturing the same, and an automobile frame structural part or automobile reinforcing part including the member. [Brief explanation of the drawings]

[0027] [Figure 1] 1 is a graph showing the relationship between temperature and time in a method for producing a high-strength steel plate according to an embodiment of the present invention. DETAILED DESCRIPTION OF THE INVENTION

[0028] Hereinafter, embodiments of the high-strength steel plate and the like according to the present invention will be described. Note that the embodiment described below is an example of the present invention, and the configuration of the present invention is not limited to this specific example.

[0029] (High strength steel plate) First, a high-strength steel plate according to one embodiment of the present invention will be described. The high-strength steel plate has a predetermined chemical composition and steel structure. Furthermore, the high-strength steel plate contains Ti and one or both of Mo and V at a quarter-thickness position, and the number density of composite carbides is 1.0 × 10 19 pieces / m 3 The average grain size of prior austenite grains is 3.00 μm or less.

[0030] <Component composition> First, the appropriate range of the chemical composition of the high-strength steel plate and the reasons for limiting it will be described. In the following description, "%" representing the content of the chemical elements of the high-strength steel plate means "% by mass" unless otherwise specified.

[0031] [C: 0.070% or more and 0.390% or less] Carbon (C) is one of the important basic components of high-strength steel sheets and affects the area fractions of ferrite, martensite, and bainite, as well as delayed fracture resistance. If the C content is too low, the area fraction of ferrite increases and the area fractions of martensite and bainite decrease, making it difficult to achieve a TS of 1180 MPa or more. Furthermore, the YR decreases. Therefore, the C content is set to 0.070% or more, preferably 0.085% or more, and more preferably 0.100% or more. On the other hand, if the C content is too high, the strength of martensite increases significantly, promoting crack propagation during delayed fracture testing and reducing the delayed fracture resistance of the stress-loaded shear edge. Therefore, the C content is set to 0.390% or less, preferably 0.375% or less, and more preferably 0.360% or less.

[0032] [Si:0.01% or more and 2.00% or less] Si is an element that increases the strength of steel sheets by suppressing cementite precipitation in martensite and by solid solution strengthening. To achieve this effect, the Si content is set to 0.01% or more, preferably 0.05% or more, and more preferably 0.10% or more. On the other hand, if the Si content is too high, carbide precipitation during bainite transformation and martensitic transformation is significantly suppressed, retained austenite increases excessively, and the hardness of martensite formed from the retained austenite during shearing increases significantly. As a result, the hardness difference between microstructures increases, resulting in a decrease in λ. Therefore, the Si content is set to 2.00% or less, preferably 1.75% or less, and more preferably 1.50% or less.

[0033] [Mn:1.70% or more and 4.00% or less] Mn is an element that improves hardenability and affects the area fractions of ferrite, martensite, and bainite. It also embrittles grain boundaries, affecting delayed fracture resistance. If the Mn content is too low, the area fraction of ferrite increases and the area fractions of martensite and bainite decrease, making it difficult to achieve a TS of 1180 MPa or higher. Furthermore, the YR decreases. Therefore, the Mn content is set to 1.70% or more, preferably 1.90% or more, and more preferably 2.10% or more. On the other hand, if the Mn content is too high, austenite is stabilized, retained austenite increases excessively, and the hardness of martensite formed from the retained austenite during shearing increases significantly. As a result, the hardness difference between microstructures increases, resulting in a decrease in λ. Furthermore, grain boundaries are embrittled, resulting in a decrease in delayed fracture resistance at the stress-loaded sheared edge. Therefore, the Mn content is set to 4.00% or less, preferably 3.80% or less, and more preferably 3.60% or less.

[0034] [P:0.100% or less] Since P is an element that segregates at prior austenite grain boundaries and embrittles the grain boundaries, excessive P content reduces the delayed fracture resistance of the stressed shear edge. Therefore, the P content is set to 0.100% or less, and preferably 0.070% or less. On the other hand, although there is no particular lower limit for the P content, since P is a solid solution strengthening element and can increase the strength of high-strength steel sheet, the P content is preferably set to 0.001% or more.

[0035] [S:0.0200% or less] S exists as sulfides and can become the starting point of delayed fracture, so if S is contained in excess, the delayed fracture resistance of the stressed sheared edge portion will decrease. Therefore, the S content is set to 0.0200% or less, and preferably 0.0050% or less. On the other hand, there is no particular lower limit for the S content, but due to constraints on production technology, the S content is preferably set to 0.0001% or more.

[0036] [Al: 1.000% or less] Al is an element that provides sufficient deoxidation and reduces inclusions in steel. If the Al content is too high, the area fraction of ferrite increases and the area fractions of martensite and bainite decrease, making it difficult to achieve a TS of 1180 MPa or more. Furthermore, the YR decreases. Therefore, the Al content is set to 1.000% or less, preferably 0.500% or less, and more preferably 0.100% or less. On the other hand, to ensure stable deoxidation, the Al content is preferably 0.010% or more, more preferably 0.015% or more, and even more preferably 0.020% or more.

[0037] [N:0.0100% or less] Since N exists as coarse nitrides and can become the starting point of delayed fracture, excessive N content reduces the delayed fracture resistance of the stressed sheared edge. Therefore, the N content is set to 0.0100% or less, and preferably 0.0070% or less. On the other hand, there is no particular lower limit for the N content, but due to production technology constraints, the N content is preferably set to 0.0001% or more.

[0038] [O:0.0100% or less] O exists as an oxide and can become the initiation point of delayed fracture, so if O is contained in excess, the delayed fracture resistance of the stressed sheared edge portion will decrease. Therefore, the O content is set to 0.0100% or less, and preferably 0.0050% or less. On the other hand, there is no particular lower limit for the O content, but due to constraints on production technology, the O content is preferably set to 0.0001% or more.

[0039] [Ti: 0.015% or more and 0.200% or less] In the present invention, Ti is an important element that forms complex carbides with V and Mo, and improves delayed fracture resistance by refining prior austenite grains and suppressing dislocation slip through the austenite pinning effect. If the Ti content is too low, the number density of the complex carbides decreases, the prior austenite grains become coarse, and their average grain size increases, resulting in a decrease in delayed fracture resistance at the stress-loaded sheared edge. Therefore, the Ti content is set to 0.015% or more, preferably 0.030% or more, and more preferably 0.045% or more. On the other hand, if the Ti content is too high, coarse inclusions form, which become the starting point for delayed fracture, thereby decreasing the delayed fracture resistance at the stress-loaded sheared edge. Therefore, the Ti content is set to 0.200% or less, preferably 0.180% or less, and more preferably 0.160% or less.

[0040] [One or both of Mo and V: Satisfies formula (1)] In the present invention, Mo and V are important elements that form complex carbides with Ti and improve delayed fracture resistance by refining prior austenite grains and suppressing dislocation slippage through the austenite pinning effect. To achieve the pinning effect during annealing in the austenite region, fine precipitates must be present in the austenite region. However, conventionally, Mo or V has not been utilized for the pinning effect because it dissolves in the matrix in the austenite region. After extensive research, the inventors discovered that the combined addition of Ti and one or both of Mo and V precipitates a large amount of complex carbides containing Ti and one or both of Mo and V in the austenite region. This results in the formation of fine complex carbides and significant refinement of prior austenite grains, dramatically improving delayed fracture resistance.

[0041] One or both of Mo and V must satisfy the following formula (1). If the value of [%Mo] / 95.94+[%V] / 50.94 is too small, the number density of complex carbides decreases, prior austenite grains become coarse, their average grain size increases, and the delayed fracture resistance of the stress-loaded sheared edge deteriorates. Therefore, this value should be 0.0010% or more, preferably 0.0030% or more, and more preferably 0.0050% or more. On the other hand, if the value of [%Mo] / 95.94+[%V] / 50.94 is too large, coarse inclusions form, which become the initiation points of delayed fracture, thereby deteriorating the delayed fracture resistance of the stress-loaded sheared edge. Therefore, this value should be 0.0200% or less, preferably 0.0180% or less, and more preferably 0.0160% or less. 0.0010≦[%Mo] / 95.94+[%V] / 50.94≦0.0200 ···(1) Here, [%Mo] and [%V] are the contents of Mo and V, respectively.

[0042] The Mo and V contents are each preferably 1.00% or less, while the Mo and V contents are each preferably 0.05% or more.

[0043] The high-strength steel sheet has a chemical composition containing the above elements, with the balance being Fe and unavoidable impurities. Here, it is preferable that the high-strength steel sheet according to one embodiment of the present invention contains only the above basic elements and the balance, with the balance being Fe (iron) and unavoidable impurities. Examples of unavoidable impurities include Zn, Pb, As, Ge, Sr, and Cs. A total content of 0.100% or less of these impurities is permitted. Unavoidable impurities may be mixed in from scrap, etc., and their inclusion is permitted as long as it does not impair the objectives of the present invention.

[0044] In addition to the above-described chemical composition, the high-strength steel sheet may further contain, by mass%, at least one element selected from the group consisting of Nb: 0.200% or less, Ta: 0.10% or less, W: 0.10% or less, B: 0.0100% or less, Cr: 1.00% or less, Ni: 1.00% or less, Co: 0.010% or less, Cu: 1.00% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.100% or less, Te: 0.100% or less, Hf: 0.100% or less, and Bi: 0.200% or less.

[0045] [Nb:0.200% or less] If the Nb content is 0.200% or less, large amounts of coarse precipitates and inclusions are not formed, and they do not become the starting point of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge portion is not reduced. Therefore, when Nb is contained, the Nb content is set to 0.200% or less, and preferably 0.100% or less. On the other hand, although there is no particular lower limit for the Nb content, the inclusion of Nb can increase the strength of the steel sheet by forming fine carbides, nitrides, or carbonitrides during hot rolling or continuous annealing. Therefore, the Nb content is preferably set to 0.001% or more.

[0046] [Ta:0.10% or less] [W: 0.10% or less] If the Ta or W content is 0.10% or less, large amounts of coarse precipitates or inclusions are not formed and do not become the initiation point of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge is not reduced. Therefore, when Ta or W is contained, the Ta or W content is set to 0.10% or less, respectively, and preferably 0.08% or less, respectively. On the other hand, although there is no particular lower limit for the Ta or W content, the inclusion of Ta or W can increase the strength of the steel sheet by forming fine carbides, nitrides, or carbonitrides during hot rolling or continuous annealing. Therefore, the Ta or W content is preferably set to 0.01% or more, respectively.

[0047] [B:0.0100% or less] If the B content is 0.0100% or less, large amounts of coarse precipitates and inclusions are not formed, and they do not become the starting point of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge is not reduced. Therefore, when B is contained, the B content is set to 0.0100% or less, and preferably 0.0080% or less. On the other hand, although there is no particular lower limit for the B content, since B is an element that segregates to austenite grain boundaries during annealing and improves hardenability, the B content is preferably set to 0.0003% or more.

[0048] [Cr:1.00% or less] [Ni: 1.00% or less] If the Cr or Ni content is 1.00% or less, large amounts of coarse precipitates or inclusions are not formed, and they do not become the starting point of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge is not reduced. Therefore, when Cr or Ni is contained, the Cr or Ni content is set to 1.00% or less, and preferably 0.80% or less. On the other hand, although there is no particular lower limit for the Cr or Ni content, since these elements improve hardenability, the Cr or Ni content is preferably set to 0.01% or more.

[0049] [Co:1.00% or less] If the Co content is 1.00% or less, large amounts of coarse precipitates and inclusions are not formed, and they do not become the starting point of delayed fracture, so the delayed fracture resistance of the stressed sheared edge is not reduced. Therefore, when Co is contained, the Co content is set to 1.00% or less, and preferably 0.80% or less. On the other hand, although there is no particular lower limit for the Co content, since Co is an element that improves hardenability, the Co content is preferably set to 0.001% or more.

[0050] [Cu:1.00% or less] If the Cu content is 1.00% or less, large amounts of coarse precipitates and inclusions are not formed, and they do not become the starting point of delayed fracture, so the delayed fracture resistance of the stressed sheared edge is not reduced. Therefore, when Cu is contained, the Cu content is set to 1.00% or less, and preferably 0.80% or less. On the other hand, although there is no particular lower limit for the Cu content, since Cu is an element that improves hardenability, the Cu content is preferably set to 0.01% or more.

[0051] [Sn:0.200% or less] If the Sn content is 0.200% or less, cracks will not form inside the steel sheet during casting or hot rolling, and the Sn content will not become the initiation point of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge will not be reduced. Therefore, when Sn is contained, the Sn content is set to 0.200% or less, and preferably 0.100% or less. On the other hand, although there is no particular lower limit for the Sn content, since Sn is an element that improves hardenability, the Sn content is preferably set to 0.001% or more.

[0052] [Sb:0.200% or less] If the Sb content is 0.200% or less, large amounts of coarse precipitates and inclusions are not formed, and they do not become the starting point of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge is not reduced. Therefore, when Sb is contained, the Sb content is set to 0.200% or less, and preferably 0.100% or less. On the other hand, although there is no particular lower limit for the Sb content, since Sb is an element that controls the softened surface thickness and enables strength adjustment, the Sb content is preferably set to 0.001% or more.

[0053] [Ca: 0.0100% or less] [Mg:0.0100% or less] [REM:0.0100% or less] If the Ca, Mg, or REM content is 0.0100% or less, large amounts of coarse precipitates or inclusions are not formed and do not become the initiation point of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge is not reduced. Therefore, when Ca, Mg, or REM is contained, the Ca, Mg, or REM content is 0.0100% or less, and preferably 0.0050% or less. On the other hand, although there is no particular lower limit for the Ca, Mg, or REM content, since these elements spheroidize the shape of nitrides and sulfides and improve the ultimate deformability of the steel sheet, the Ca, Mg, or REM content is preferably 0.0005% or more.

[0054] [Zr:0.100% or less] [Te:0.100% or less] If the Zr or Te content is 0.100% or less, large amounts of coarse precipitates or inclusions are not formed and do not become the initiation points of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge is not reduced. Therefore, when Zr or Te is contained, the Zr or Te content is preferably 0.100% or less, and 0.080% or less, respectively. On the other hand, although there are no particular lower limits for the Zr and Te contents, since these elements spheroidize the shape of nitrides and sulfides and improve the ultimate deformability of the steel sheet, the Zr or Te content is preferably 0.001% or more, respectively.

[0055] [Hf:0.100% or less] If the Hf content is 0.100% or less, large amounts of coarse precipitates and inclusions are not formed, and they do not become the initiation points of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge is not reduced. Therefore, when Hf is contained, the Hf content is set to 0.100% or less, and preferably 0.080% or less. On the other hand, although there is no particular lower limit for the Hf content, since Hf is an element that spheroidizes the shape of nitrides and sulfides and improves the ultimate deformability of the steel sheet, the Hf content is preferably set to 0.010% or more.

[0056] [Bi:0.200% or less] If the Bi content is 0.200% or less, large amounts of coarse precipitates and inclusions are not formed, and they do not become the starting point of delayed fracture, so the delayed fracture resistance of the stressed sheared edge is not reduced. Therefore, when Bi is contained, the Bi content is set to 0.200% or less, and preferably 0.100% or less. On the other hand, although there is no particular lower limit for the Bi content, since Bi is an element that reduces segregation, the Bi content is preferably set to 0.001% or more.

[0057] In addition, when the content of each of the above-mentioned Nb, Ta, W, B, Cr, Ni, Co, Cu, Sn, Sb, Ca, Mg, REM, Zr, Te, Hf, and Bi is less than the preferable lower limit, the effect of the present invention is not impaired, and therefore they are treated as unavoidable impurities.

[0058] <Steel structure> Next, the steel structure of the high-strength steel plate will be described. Note that the steel structure of the high-strength steel plate described below is that at a 1 / 4 position in the plate thickness direction of the high-strength steel plate (a position corresponding to 1 / 4 of the plate thickness in the depth direction from the surface of the high-strength steel plate).

[0059] [Total area ratio of martensite and bainite: 70% or more] Martensite is a transformation phase that forms below the Ms point, and in the present invention, tempering is not an issue. Bainite is a mixed structure region that forms above the Ms point and is composed of angular bainitic ferrite, iron-based carbides, and retained austenite, and in the present invention, includes upper bainite and lower bainite. By including martensite and bainite, a TS of 1180 MPa or higher can be achieved. Therefore, the total area fraction of martensite and bainite is set to 70% or more, preferably 80% or more, and more preferably 90% or more. Meanwhile, there is no particular upper limit to the area fraction of martensite and bainite, and the above-mentioned effects can be obtained even if the total area fraction of martensite and bainite is 100%.

[0060] [Ferrite area ratio: 20% or less] Ferrite is a soft BCC iron formed at high temperatures, and in the present invention, it includes allotriomorph ferrite and idiomorph ferrite. Because ferrite is soft, if the area fraction of ferrite is too high, TS decreases. Also, because ferrite becomes the starting point of plastic deformation, YR decreases. Furthermore, the difference in hardness between structures increases, and λ decreases. Therefore, the area fraction of ferrite is set to 20% or less, preferably 13% or less, and more preferably 10% or less. On the other hand, there is no particular lower limit for the area fraction of ferrite, and the effects of the present invention can be obtained even if the area fraction of ferrite is 0%.

[0061] The area ratios of martensite, bainite, and ferrite can be measured as follows. First, a sample is cut out from a steel sheet so that the plate thickness cross section (L cross section) parallel to the rolling direction serves as the observation surface. The observation surface of the sample is mirror-polished using diamond paste, then finish-polished using colloidal silica, and further etched using 1 volume % nital to reveal the structure. Next, a portion of the observation surface of the sample corresponding to a position 1 / 4 of the plate thickness of the steel sheet is observed at 3000x magnification using a scanning electron microscope (SEM) at an acceleration voltage of 10 kV, and SEM images of three fields of view are obtained. From the obtained SEM images, the area of ​​each structure is calculated using Adobe Photoshop (manufactured by Adobe Systems). The area of ​​each structure is divided by the measured area to obtain a value, and the average value for the three fields of view is calculated to determine the area ratio of each structure.

[0062] In SEM images, martensite is a structure with a hierarchical structure with fine internal irregularities. Ferrite is a gray, flat structure region without carbides. Bainite is a structure region composed of gray, angular bainitic ferrite and white, contrasting iron-based carbides. Therefore, martensite, ferrite, and bainite can be distinguished from each other.

[0063] [Area ratio of retained austenite: 15% or less] If the amount of retained austenite is too large, a large amount of hard martensite is generated from the retained austenite during shearing. As a result, the difference in hardness between the structures increases, and λ decreases. Therefore, the area fraction of retained austenite is set to 15% or less, and preferably 10% or less. On the other hand, there is no particular lower limit for the area fraction of retained austenite, and the effects of the present invention can be obtained even if the area fraction of retained austenite is 0%.

[0064] The area ratio of retained austenite can be measured as follows. First, the steel plate is ground so that the position at 1 / 4 of its plate thickness becomes the measurement surface, and then it is further polished by 0.1 mm by chemical polishing to obtain a sample. For the measurement surface of the obtained sample, using a Co Kα ray source with an X-ray diffractometer, the integrated reflection intensities of the (200), (220), and (311) planes of fcc iron (austenite), and the (200), (211), and (220) planes of bcc iron are measured respectively. The intensity ratio of the integrated reflection intensity of each plane of fcc iron to the integrated reflection intensity of each plane of bcc iron measured is obtained, and a total of nine intensity ratios are obtained. The average value of the nine obtained intensity ratios is taken as the volume ratio of retained austenite. In the present invention, such a volume ratio of retained austenite is regarded as the area ratio of retained austenite.

[0065] [Remaining structure] The steel structure may have a structure other than the above-described martensite, ferrite, bainite, and retained austenite (remaining structure). However, the area ratio of the remaining structure is preferably 5% or less for the reason that the effects of the present invention are not impaired. Examples of the remaining structure include other structures known as steel plate structures such as pearlite and MnS inclusions.

[0066] The area ratio of the remaining structure can be obtained by subtracting the area ratios of the above-described martensite, bainite, ferrite, and retained austenite from 100.

[0067] <The number density of composite carbides containing Ti and further containing one or both of Mo and V: 1.0×10 19 pieces / m 3 or more> The composite carbide containing Ti and one or both of Mo and V improves delayed fracture resistance, and its number density is an important constituent element of the present invention. First, the composite carbide formed in the austenite region pins austenite grains and significantly refines the prior austenite grains. This suppresses hydrogen accumulation at the prior austenite grain boundaries, which are the propagation path of delayed fracture cracks. Furthermore, the composite carbide itself is hard, and in an environment where delayed fracture is likely to occur (under hydrogen charging), it suppresses dislocation slip at the tip of an embedded crack and suppresses the generation of voids. These two actions suppress delayed fracture crack propagation and improve the delayed fracture resistance of the stress-loaded shear edge.

[0068] If the number density of the composite carbide containing Ti and either or both of Mo and V is too low, the prior austenite grains become coarse and their average grain size becomes large. This reduces the delayed fracture suppression effect of the composite carbide itself, and the delayed fracture resistance of the stress-loaded sheared edge portion decreases. Therefore, the number density should be less than 1.0 × 10 19 pieces / m 3 Above 5.0 x 10 19 pieces / m 3 More than 1.0×10 is preferable. 20 pieces / m 3 On the other hand, the upper limit of the number density is not particularly limited, and the number density is generally 1.0 × 10 22 pieces / m 3 The following is the result.

[0069] The number density of composite carbides containing Ti and either or both of Mo and V can be measured as follows. First, a steel sheet is ground so that the observation surface is at 1 / 4 of the sheet thickness, and then electrolytically polished to prepare a sample. A 200 nm × 200 nm region of the observation surface of the prepared sample is observed using a transmission electron microscope (TEM) or a scanning transmission electron microscope (STEM) at an acceleration voltage of 200 kV. Composition analysis of the precipitates is performed using an energy dispersive X-ray analyzer attached to the TEM or STEM, and the number of composite carbides containing Ti and either or both of Mo and V is counted. The above measurement is performed in five fields of view, and the number of composite carbides is divided by the measurement volume (200 nm × 200 nm × 100 nm) for each of the five fields of view to obtain the value. Note that the measurement depth using the TEM or STEM is assumed to be 100 nm. The average of the determined values ​​is taken as the number density of the composite carbide containing Ti and further containing one or both of Mo and V.

[0070] <Average grain size of prior austenite grains: 3.00 μm or less> The average grain size of prior austenite grains is an important constituent element of the present invention. Significantly refining the prior austenite grains suppresses hydrogen accumulation at the prior austenite grain boundaries, which are the propagation paths of delayed fracture cracks. These fine prior austenite grains are achieved by the pinning effect of the composite carbides described above. If the average grain size of the prior austenite grains is too large, the interfacial area per unit volume decreases, hydrogen tends to accumulate at the prior austenite grain boundaries, and the delayed fracture resistance of the stress-loaded shear edge deteriorates. Therefore, the average grain size of the prior austenite grains is set to 3.00 μm or less, preferably 2.00 μm or less, and more preferably 1.00 μm or less. Meanwhile, there is no particular lower limit for the average grain size of the prior austenite grains, and the average grain size is generally 0.10 μm or more.

[0071] The average grain size of prior austenite grains can be measured as follows. First, a sample is cut out from a steel sheet so that the observation surface is the thickness cross section (L cross section) parallel to the rolling direction. The observation surface of the sample is mirror-polished using diamond paste, and then finish-polished using colloidal silica. Next, a location on the observation surface of the sample corresponding to 1 / 4 of the steel sheet thickness is measured using electron backscatter diffraction (EBSD) attached to a scanning electron microscope (SEM) at an accelerating voltage of 15 kV under conditions of 10 μm × 10 μm and a step size of 20 nm for three fields of view. Using OIM Analysis (manufactured by EDAX), matrix reconstruction is performed to obtain prior austenite grain maps for the three fields of view. A grid-like test line is drawn on the obtained prior austenite grain map, and the number of intersections between the test line and the prior austenite grain boundaries is counted. The average grain size of the prior austenite grains is calculated by dividing the length of the test line by the number of intersections.

[0072] <Characteristics of high-strength steel plates> [Tensile strength: 1180 MPa or more] The tensile strength (TS) of the high-strength steel sheet obtained by the present invention is 1180 MPa or more, preferably 1470 MPa or more, while the tensile strength of the high-strength steel sheet is generally 2000 MPa or less.

[0073] [Yield ratio: 70% or more] The yield ratio (YR) of the high strength steel sheet obtained by the present invention is 70% or more, preferably 72% or more, and more preferably 79% or more, while the yield ratio of the high strength steel sheet is generally 90% or less.

[0074] [Total elongation: 7.0% or more] The total elongation (El) of the high strength steel sheet obtained by the present invention is 7.0% or more, preferably 7.5% or more, while the total elongation of the high strength steel sheet is generally 13.0% or less.

[0075] The tensile strength, yield ratio, and total elongation can be measured by the tensile test described below. The tensile test is performed in accordance with JIS Z 2241:2021. First, a JIS No. 5 test piece is taken from the steel plate so that the longitudinal direction is perpendicular to the rolling direction of the steel plate. Using the taken test piece, a crosshead speed of 1.67 × 10 -1 A tensile test is carried out under the condition of 100 mm / s to measure the yield strength (YS), tensile strength (TS), and total elongation (El). From the measurement results, the yield ratio (YR) (= 100 × YS / TS) is calculated.

[0076] [Limiting hole expansion ratio: 25% or more] The high strength steel sheet obtained by the present invention has a limiting hole expansion ratio (λ) of 25% or more, preferably 28% or more, while the limiting hole expansion ratio of the high strength steel sheet is generally 60% or less.

[0077] The critical hole expansion ratio can be measured by the hole expansion test described below. The hole expansion test is performed in accordance with JIS Z 2256. First, a steel plate is sheared to obtain a test piece measuring 100 mm x 100 mm. A hole with a diameter of 10 mm is punched into the obtained test piece with a clearance of 12.5%. Then, using a die with an inner diameter of 75 mm and a blank holder force of 9 ton (88.26 kN), a conical punch with an apex angle of 60° is pressed into the hole, and the hole diameter D at the crack initiation limit is measured. f Using the initial hole diameter D0 [mm], calculate the limit hole expansion ratio (λ) [%] using the following formula (5). λ={(D f -D0) / D0}×100 (5)

[0078] [Excellent resistance to delayed fracture at stress-loaded shear end faces] The high-strength steel sheet obtained by the present invention has excellent delayed fracture resistance at the stress-loaded sheared edge. That is, in the delayed fracture test described below, the fracture limit stress of the high-strength steel sheet is 900 MPa or more, preferably 1050 MPa or more. On the other hand, the fracture limit stress of the high-strength steel sheet is generally 1400 MPa or less.

[0079] The critical fracture stress can be measured as follows. A 70 mm wide, 16 mm long rectangular test piece is obtained from a steel plate so that the surface perpendicular to the rolling direction of the steel plate is the sheared end face at a shear angle of 0.5° and a clearance of 10%. The sheared end face on the removed side is removed by mechanical grinding. Multiple test pieces are taken and each is subjected to bending deformation by four-point bending so that a stress equivalent to 600 MPa to 1500 MPa is applied to the center. The bent-deformed test pieces are then immersed for 96 hours in a 0.1% by mass aqueous solution of ammonium thiocyanate at 25°C, the pH of which has been adjusted to 7.0 using McIlvaine buffer solution. After immersion, the test pieces are removed from the solution and checked for the presence or absence of cracks. The maximum stress applied to a test piece that did not develop cracks is taken as the critical fracture stress.

[0080] (High-strength plated steel sheet) A high-strength plated steel sheet according to one embodiment of the present invention comprises the above-described high-strength steel sheet and a plating layer located on the surface of the high-strength steel sheet. The plating layer is formed by a plating process described below. The type of plating process is not particularly limited, and examples thereof include hot-dip plating and electroplating. Examples of the plating layer include a zinc plating layer (Zn plating layer) and an Al plating layer. A zinc plating layer is preferred as the plating layer. The zinc plating layer may contain elements such as Al and Mg. The plating layer may also be an alloyed plating layer (alloyed plating layer).

[0081] The composition of the coating layer is not particularly limited and may be a general composition. For example, when the coating layer is a hot-dip galvanized layer or a galvannealed hot-dip galvanized layer, the coating layer may contain 20 mass% or less of Fe, 0.001 to 1.0 mass% of Al, and a total of 0 mass% to 3.5 mass% of at least one element selected from the group consisting of Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr, Co, Ca, Cu, Li, Ti, Be, Bi, and REM, with the remainder being Zn and unavoidable impurities.

[0082] If the coating layer is a hot-dip galvanized layer, the coating weight per side of the coating layer is 20 g / m 2 On the other hand, the coating weight of the hot-dip galvanized layer per side is preferably 80 g / m or more. 2 On the other hand, when the plated layer is a hot-dip galvanized layer, the Fe content in the plated layer is preferably less than 7 mass %.

[0083] The plating layer may be a galvannealed layer obtained by alloying a hot-dip galvanized layer. When the plating layer is a galvannealed layer, the Fe content in the plating layer is preferably 7% by mass or more. On the other hand, when the plating layer is a galvannealed layer, the Fe content in the plating layer is preferably 20% by mass or less, and more preferably 15% by mass or less.

[0084] (components) A member according to one embodiment of the present invention includes the above-described high-strength steel sheet or high-strength plated steel sheet. The high-strength steel sheet included in the member is excellent in all of part strength, ductility, stretch flangeability, and delayed fracture resistance at stressed shear end faces, and therefore the member can be suitably used particularly for automotive frame structural parts or automotive reinforcing parts.

[0085] (parts) An automobile structural part or automobile reinforcing part according to one embodiment of the present invention includes the above-mentioned member. The high-strength steel sheet included in the member included in the part is excellent in all of part strength, ductility, stretch flangeability, and delayed fracture resistance of a stress-loaded shear edge, and is therefore suitable for an automobile structural part or an automobile reinforcing part.

[0086] (Manufacturing method for high-strength steel plates) Next, a method for manufacturing a high-strength steel plate according to one embodiment of the present invention will be described.

[0087] First, a steel material having the above-described composition is melted to produce a steel slab. The method for melting the molten steel to become the steel slab is not particularly limited, and known melting methods using a converter, an electric furnace, or the like can be adopted. The steel slab is preferably produced by a continuous casting method to prevent macrosegregation, but it can also be produced by other methods such as an ingot casting method or a thin slab casting method.

[0088] <Heating process> [Slab heating temperature: 1150℃ or higher] Next, the steel slab is heated to 1150°C or higher (heating step). If the temperature of the steel slab in the heating step (slab heating temperature) is too low, Ti, Mo, or V will not dissolve in austenite during rough rolling and will remain as coarse inclusions. This reduces the number density of composite carbides containing Ti and one or both of Mo and V in the final structure. Furthermore, prior austenite grains in the final structure also become coarse, increasing their average grain size. Therefore, the slab heating temperature is set to 1150°C or higher, preferably 1180°C or higher, and more preferably 1200°C or higher. On the other hand, to prevent an increase in scale loss, the slab heating temperature is preferably 1300°C or lower. The slab heating temperature refers to the surface temperature of the steel slab.

[0089] <Hot rolling process> Next, the steel slab is subjected to rough rolling and finish rolling to produce a hot-rolled sheet (hot rolling process). In one example, the steel slab is cooled to room temperature, then reheated and hot-rolled (rough rolling and finish rolling). In another example, the produced steel slab may be charged into a heating furnace as a hot slab without being cooled to room temperature, or may be briefly kept at room temperature and then immediately rough-rolled. Alternatively, the steel slab subjected to rough rolling is called a rough-rolled sheet.

[0090] [4 or more passes with a rolling reduction of 13% or more in the temperature range of 1050°C or higher] In rough rolling, if the number of passes at a reduction rate of 13% or more in a temperature range of 1050°C or higher is small, Ti, Mo, or V will be unevenly distributed within the steel sheet, resulting in a decrease in the number density of complex carbides in the final structure. Furthermore, prior austenite grains in the final structure will also become coarse, increasing their average grain size. This is thought to be because the diffusion of substitutional atoms such as Ti, Mo, or V is promoted during dynamic recrystallization of austenite grains. Therefore, the number of passes at a reduction rate of 13% or more in a temperature range of 1050°C or higher should be 4 passes or more, and preferably 5 passes or more. While there is no particular upper limit to the number of passes, the number of passes is generally 10 passes or less.

[0091] Next, the rough-rolled sheet is subjected to finish rolling to obtain a hot-rolled sheet. When the slab heating temperature is low or when the temperature of the rough-rolled sheet is low, it is preferable to heat the rough-rolled sheet using a bar heater or the like before finish rolling in order to prevent problems during hot rolling. The temperature when performing finish rolling (finish rolling temperature) is preferably 700°C or higher. This reduces the rolling load. Furthermore, the rolling reduction rate in the unrecrystallized austenite state is reduced, suppressing the development of abnormal structures elongated in the rolling direction, and improving workability.

[0092] In order to reduce the rolling load, part or all of the finish rolling may be performed as lubricated rolling. Lubricated rolling is also preferred from the viewpoint of uniforming the shape and material properties of the steel sheet. The friction coefficient during lubricated rolling is preferably 0.10 or more. Furthermore, the friction coefficient during lubricated rolling is preferably 0.25 or less. The coiling temperature after hot rolling is preferably 300°C or more from the viewpoint of improving the sheet passing properties during cold rolling and annealing, which will be described later. Furthermore, the coiling temperature after hot rolling is preferably 700°C or less.

[0093] Finish rolling may be performed continuously by joining together the rough rolled sheets. Before performing finish rolling, the rough rolled sheet may be temporarily wound up. Furthermore, the hot rolled sheet after finish rolling is appropriately subjected to a winding process.

[0094] <Cold rolling process> Next, the hot-rolled sheet is subjected to pickling and cold rolling to obtain a cold-rolled sheet (cold-rolling process). Pickling removes oxides from the surface of the hot-rolled sheet, resulting in a high-strength steel sheet as a final product with excellent chemical conversion treatability and coating layer quality. Pickling may be performed once or multiple times.

[0095] The pickled hot-rolled sheet is optionally subjected to a softening heat treatment and then cold-rolled. The cold-rolling conditions are not particularly limited, but the cumulative reduction in cold rolling is preferably 20 to 75%. The number of rolling passes and the reduction in each pass are not particularly limited.

[0096] <Annealing process> Next, the cold-rolled sheet is annealed (annealing step). FIG. 1 shows the relationship between temperature and time in the annealing step and the cooling step. In the present invention, the annealing step is performed by heating the cold-rolled sheet until the maximum temperature is reached and the cold-rolled sheet reaches 750°C. When heating the cold-rolled sheet, the temperature range of 600°C or higher and Ac3°C or lower is designated as T1, and the temperature range of Ac3°C or higher and (Ac3+200)°C or lower is designated as T2. Here, the maximum temperature of the cold-rolled sheet is Ac3°C or higher and (Ac3+200)°C or lower. Note that FIG. 1 shows an example in which isothermal holding is performed at the maximum temperature, but is not particularly limited thereto.

[0097] [Average heating rate v1 in the temperature range T1 between 600°C and Ac3°C: 1.5°C / s or more] In the annealing process, if the average heating rate v1 in the temperature range T1 of 600°C or higher and Ac3°C or lower is too small, composite carbides containing Ti and one or both of Mo and V will precipitate non-uniformly in the unrecrystallized structure and then coarsen during the subsequent residence in the austenite region. This will result in a decrease in the number density of the composite carbides, and the prior austenite grains will also coarsen, increasing their average grain size. Therefore, the average heating rate v1 in the temperature range T1 is set to 1.5°C / s or higher, and preferably 2.0°C / s or higher. On the other hand, although there is no particular upper limit for the average heating rate v1, from the viewpoints of operability and damage to the furnace body, the average heating rate v1 is preferably 100°C / s or lower.

[0098] The Ac3 point can be calculated using the following formula (6). Ac3 point (℃)=910-203×[%C]1 / 2+44.7×[%Si]-30×[%Mn]+700×[%P]+400×[%Al]+40 0×[%Ti]+104×[%V]+13.1×[%W]-11×[%Cr]+31.5×[%Mo]-15.2×[%Ni]-20×[%Cu] ...(6) Here, [%X] indicates the content (mass%) of component element X in the steel, and is set to 0 if no component element X is contained.

[0099] [Maximum temperature of cold rolled sheet: Ac3℃ or higher (Ac3+200)℃ or lower] The maximum temperature is set to Ac3°C or higher in order to promote the reverse transformation from ferrite to austenite and obtain sufficient amounts of martensite and bainite. On the other hand, if the maximum temperature of cold-rolled sheet exceeds (Ac3 + 200)°C, the number density of complex carbides decreases and prior austenite grains also coarsen, so the maximum temperature is set to (Ac3 + 200)°C or lower.

[0100] [Retention condition from the time when Ac reaches 3°C through the maximum temperature reached to the time when it reaches 750°C: Formula (2) is satisfied] The retention condition from the time when the temperature of the cold-rolled sheet reaches Ac3°C through the maximum temperature to the time when it reaches 750°C satisfies the following formula (2). P≦30000 (2) where:

number

number

[0101] In steel sheets that have reached the Ac3 point, i.e., the austenite single-phase region, during the annealing process, prior austenite grains are kept fine due to the pinning effect of complex carbides containing Ti and one or both of Mo and V. On the other hand, the constituent atoms of the complex carbides (especially Ti) diffuse, causing the complex carbides to coarsen, which in turn causes the prior austenite grains to coarsen. Based on these phenomena, the inventors defined a parameter P based on the diffusion phenomenon of Ti in the austenite region. Parameter P is calculated from the temperature and time during the annealing process, and by appropriately controlling the value of parameter P, the number density of the complex carbides and the average grain size of the prior austenite grains can be controlled.

[0102] If the amount of P is too large, the complex carbides become coarse and the number density of the complex carbides decreases. In addition, the pinning effect becomes insufficient, and the average grain size of the prior austenite grains also becomes coarse. Therefore, the amount of P is 30,000 nm or less. 2 Below 25,000 nm 2 Preferably below 20000nm 2 The temperature history in the annealing step is not particularly limited as long as the parameter P is within the above range. On the other hand, although there is no particular lower limit for P, from the viewpoint of controllability, P is preferably 1000 nm or less. 2 More than 1500nm is preferable. 2 More preferably, 2000 nm or more 2 The above is even more preferable.

[0103] <Cooling process> Next, the cold-rolled sheet is cooled to produce a high-strength steel sheet (cooling step). As shown in Figure 1, in the cooling step, the temperature range of 550°C or higher and 750°C or lower is designated as T3, and the temperature range of 350°C or higher and 550°C or lower is designated as T4.

[0104] [Average cooling rate v3 in the temperature range T3 between 550°C and 750°C: 4°C / s or more] In the cooling process, the temperature range T3 of 550°C or higher and 750°C or lower is the temperature range in which ferrite transformation occurs. If the average cooling rate v3 in the temperature range T3 is too low, excessive ferrite transformation occurs, resulting in a high ferrite area ratio. Therefore, the average cooling rate v3 is set to 4°C / s or higher, and preferably 6°C / s or higher. On the other hand, although there is no particular upper limit for the average cooling rate v3, from the viewpoint of controllability, the average cooling rate v3 is preferably 1500°C / s or lower.

[0105] [Dwell time t4 in the temperature range T4 of 350°C to 550°C: 600 seconds or less] In the cooling process, the temperature range T4 of 350°C to 550°C is a temperature range in which carbon enrichment in untransformed austenite due to bainite transformation progresses, resulting in the formation of retained austenite. If the residence time t4 in the temperature range T4 is too long, the amount of retained austenite increases. Therefore, the residence time t4 is set to 600 seconds or less, and preferably 500 seconds or less. On the other hand, although there is no particular lower limit for the residence time t4, from the viewpoint of controllability, the residence time t4 is preferably 0.5 seconds or more.

[0106] <Reheating treatment (optimal conditions)> In the cooling step, the cold-rolled sheet cooled to 350°C or less may be subjected to a reheating treatment in which, in order to further increase the YR, the cold-rolled sheet is cooled to a cooling stop temperature of 350°C or less, then heated from the cooling stop temperature to a reheating temperature, and then cooled to room temperature. The cooling stop temperature can be a temperature of room temperature or higher and 350°C or lower. The reheating temperature is preferably a temperature higher than the cooling stop temperature and is 150°C or higher. On the other hand, the reheating temperature is preferably 400°C or lower. Furthermore, the reheating time is preferably 10 seconds or longer. On the other hand, the reheating time is preferably 1000 seconds or shorter. If the reheating treatment is not performed, the cold-rolled sheet is cooled to room temperature in the cooling step.

[0107] (Method of manufacturing high-strength plated steel sheets) Next, a method for producing a high-strength plated steel sheet according to one embodiment of the present invention will be described. The method for producing a high-strength plated steel sheet includes the above-described method for producing a high-strength steel sheet and, in the cooling step, a plating step of applying a plating treatment to the cold-rolled sheet.

[0108] Examples of the plating treatment include a hot-dip galvanizing treatment (a treatment for forming a hot-dip galvanized layer), a hot-dip galvannealing treatment (a treatment for forming a hot-dip galvannealed layer by performing an alloying treatment after the hot-dip galvanizing treatment), etc. Alternatively, an electroplating layer may be formed by an electroplating treatment.

[0109] When hot-dip galvanizing is performed, it is preferable to immerse the cold-rolled sheet in a galvanizing bath and then adjust the coating weight of the coating layer by gas wiping or the like. The bath temperature of the galvanizing bath is not particularly limited, but is preferably 440°C or higher. The bath temperature of the galvanizing bath is preferably 500°C or lower. The Al content of the galvanizing bath is preferably 0.10 mass% or higher. The Al content of the galvanizing bath is preferably 0.23 mass% or lower. The galvanizing treatment is preferably performed after residence in a temperature range T3 of 550°C or higher and 750°C or lower in the above-mentioned cooling step.

[0110] The alloying temperature is preferably 470°C or higher to improve the Zn-Fe alloying rate and productivity, and is preferably 600°C or lower, more preferably 560°C or lower, to effectively prevent untransformed austenite from transforming into pearlite and improve TS.

[0111] The high-strength steel sheet after annealing and cooling may be subjected to skin-pass rolling. The reduction ratio of the skin-pass rolling is preferably 0.01% or more from the viewpoint of stabilizing the shape. On the other hand, although there is no particular upper limit to the reduction ratio of the skin-pass rolling, from the viewpoint of productivity, the reduction ratio is preferably 1.50% or less. Skin-pass rolling may be performed online or offline. Skin-pass rolling with the desired reduction ratio may be performed all at once, or may be performed in several steps.

[0112] From the viewpoint of productivity, the above-mentioned series of treatments such as annealing and plating treatments are preferably carried out in a CAL (Continuous Annealing Line) or a CGL (Continuous Galvanizing Line).

[0113] (Manufacturing method for components) Next, a method for manufacturing a component according to one embodiment of the present invention will be described. The high-strength steel sheet or high-strength plated steel sheet described above is subjected to one or both of forming and joining processes to form a component. The forming and joining processes can be carried out by conventional methods.

[0114] Although the present invention has been described above with reference to an embodiment, the present invention is not limited to the description of the present embodiment, which is a part of the disclosure of the present invention. In other words, other embodiments, examples, and operational techniques made by those skilled in the art based on the present embodiment are all included in the scope of the present invention. For example, in the series of heat treatments in the above-described manufacturing method, as long as the thermal history is satisfied, other conditions are not particularly limited, and the equipment for performing the heat treatment is not particularly limited.

[0115] For steps and conditions not described in this specification, conventional methods can be used. [Example]

[0116] The present invention will be specifically described below with reference to examples, but the present invention is not limited to the examples described below.

[0117] Molten steel having the chemical composition shown in Table 1 (the balance being Fe and unavoidable impurities) was melted in a converter and then subjected to continuous casting to obtain steel slabs.

[0118] [Table 1] TIFF0007806983000006.tif233103

[0119] The obtained steel slab was subjected to hot rolling to obtain a hot-rolled sheet. Table 2 shows the slab heating temperature and the number of passes with a rolling reduction of 13% or more in a temperature range of 1050°C or higher. Next, finish rolling was performed at a finish rolling temperature of 900°C, followed by coiling at 500°C and cooling to room temperature to obtain a hot-rolled sheet. The obtained hot-rolled sheet was pickled, then subjected to a softening heat treatment at 500°C, and then subjected to cold rolling at a rolling reduction of 50%. In this way, a cold-rolled sheet with a sheet thickness of 1.4 mm was obtained. The obtained cold-rolled sheet was subjected to annealing, cooling, and, if necessary, reheating under the conditions shown in Table 2, and in some cases, further subjected to plating to obtain a high-strength steel sheet (cold-rolled steel sheet) or a high-strength plated steel sheet.

[0120] In some examples, hot-dip galvanizing treatment was performed while the steel sheet was held in the temperature range T4 of 350°C or higher and 550°C or lower, forming a coating layer (hot-dip galvanized layer) on both sides to produce a high-strength coated steel sheet. That is, a hot-dip galvanized steel sheet (GI) was obtained. For the hot-dip galvanizing treatment, a hot-dip galvanizing bath (bath temperature: 470°C) containing 0.20 mass% Al, with the balance being Zn and unavoidable impurities, was used. The coating weight of the hot-dip galvanized layer per side was 45 to 72 g / m 2 The composition of the formed hot-dip galvanized layer contained 0.1 to 1.0 mass % of Fe, 0.2 to 1.0 mass % of Al, and the remainder being Fe and unavoidable impurities.

[0121] In some other examples, a galvannealed hot-dip galvanizing treatment was performed while the steel sheet was held in the temperature range T4 of 350°C or higher and 550°C or lower, forming a coating layer (galvannealed hot-dip galvanized layer) on both sides to produce a high-strength coated steel sheet. That is, a galvannealed hot-dip galvanized steel sheet (GA) was obtained. For the galvannealed hot-dip galvanizing treatment, a galvannealed bath (bath temperature: 470°C) containing 0.14 mass% Al, with the balance being Zn and unavoidable impurities, was used. The alloying treatment was performed at 530°C. The coating weight of the galvannealed hot-dip galvanized layer per side was 45 g / m 2The composition of the formed galvannealed layer contained 7 to 15 mass % of Fe, 0.1 to 1.0 mass % of Al, and the remainder being Fe and unavoidable impurities.

[0122] In Table 2, the "Plating Type" column indicates "GI" if a hot-dip galvanized layer was formed, "GA" if an alloyed hot-dip galvanized layer was formed, and "CR" if no plating layer was formed.

[0123] [Table 2] TIFF0007806983000008.tif233170

[0124] The area ratios of martensite, bainite, ferrite, retained austenite, and the remaining structure were determined at the 1 / 4 thickness position of the obtained steel plate by the methods described above. Furthermore, the number density of composite carbides containing Ti and either or both of Mo and V, and the average grain size of prior austenite grains were measured at the 1 / 4 thickness position of the obtained steel plate by the methods described above. Furthermore, various properties of the obtained steel plate were evaluated by the methods described above. The results are shown in Table 3.

[0125] [Table 3] TIFF0007806983000010.tif233165

[0126] As shown in Table 3, the examples of the present invention have high strength and are excellent in part strength, ductility, stretch flangeability, and delayed fracture resistance at the stressed sheared edge. On the other hand, the comparative examples are inferior in one or more of part strength, ductility, stretch flangeability, and delayed fracture resistance at the stressed sheared edge. [Industrial Applicability]

[0127] According to the present invention, it is possible to provide a high-strength steel sheet and a high-strength plated steel sheet which are excellent in part strength, ductility, stretch flangeability, and delayed fracture resistance at stressed sheared edges. By applying the high-strength steel sheet of the present invention to, for example, automobile structural members, it is possible to reduce the weight of the automobile body and thereby improve fuel economy, and the industrial value of the steel sheet is extremely great.

Claims

1. In mass%, C: 0.070% or more and 0.390% or less, Si: 0.01% or more and 2.00% or less, Mn: 1.70% or more and 4.00% or less, P: 0.100% or less, S: 0.0200% or less, Al: 1.000% or less, N: 0.0100% or less, O: 0.0100% or less, Ti: 0.015% or more and 0.200% or less, and One or both of Mo and V satisfying the following formula (1): and the balance being Fe and unavoidable impurities; At the 1 / 4 plate thickness position, The total area ratio of martensite and bainite is 70% or more, The area ratio of ferrite is 20% or less, a steel structure in which the area ratio of retained austenite is 15% or less; and At the 1 / 4 plate thickness position, The number density of the composite carbide containing Ti and further containing one or both of Mo and V is 1.0 × 10 １９ pieces / m ３ That's all, A high-strength steel plate having an average grain size of prior austenite grains of 3.00 μm or less. 0.0010≦[%Mo] / 95.94+[%V] / 50.94≦0.0200...(1) Here, [%Mo] and [%V] are the contents of Mo and V, respectively.

2. The component composition further includes, in mass %, Nb: 0.200% or less, Ta: 0.10% or less, W: 0.10% or less, B: 0.0100% or less, Cr: 1.00% or less, Ni: 1.00% or less, Co: 1.00% or less, Cu: 1.00% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.100% or less, Te: 0.100% or less, Hf: 0.100% or less, and Bi: 0.200% or less The high-strength steel plate according to claim 1, further comprising at least one element selected from the group consisting of:

3. A high-strength plated steel sheet comprising the high-strength steel sheet according to claim 1 or 2 and a plating layer located on a surface of the high-strength steel sheet.

4. A member comprising the high-strength steel plate according to claim 1 or 2.

5. A member comprising the high-strength plated steel sheet according to claim 3.

6. A frame structural part of an automobile or a reinforcing part of an automobile, comprising the member according to claim 4.

7. A frame structural part of an automobile or a reinforcing part of an automobile, comprising the member according to claim 5.

8. A method for manufacturing a high-strength steel plate according to claim 1 or 2, comprising: A heating step of heating a steel slab having the component composition according to claim 1 or 2 to 1150°C or higher; A hot rolling process in which the steel slab is subjected to rough rolling with a rolling reduction of 13% or more in a temperature range of 1050 ° C. or higher and the number of passes is 4 or more, and then subjected to finish rolling to obtain a hot-rolled sheet; Next, a cold rolling step in which the hot-rolled sheet is subjected to pickling and cold rolling to obtain a cold-rolled sheet; Next, an annealing step of annealing the cold-rolled sheet; Next, a cooling step of cooling the cold-rolled sheet to room temperature to obtain a high-strength steel sheet; and In the annealing step, Ac above 600℃ ３ The average heating rate v1 in the temperature range T1 of 1.5 ° C. / s or less is 1.5 ° C. / s or more, The maximum temperature of the cold-rolled sheet is Ac ３ ℃ or more (Ac ３ +200) ° C or less in the temperature range T2, The temperature of the cold-rolled sheet is Ac ３ The retention condition from the time when the temperature reaches 750°C through the maximum temperature to the time when the temperature reaches 750°C satisfies the following formula (2), In the cooling step, The average cooling rate v3 in the temperature range T3 of 550°C or higher and 750°C or lower is 4°C / s or higher, A method for manufacturing a high-strength steel plate, wherein the residence time t4 in a temperature range T4 of 350°C or higher and 550°C or lower is 600 seconds or shorter. P≦30000 (2) where: [Equation 1] [Equation 2] In the annealing step, the temperature of the cold-rolled sheet is Ac ３ The time when the temperature of the cold-rolled sheet reaches 750°C for the first time is defined as t=0 seconds, the time when the annealing process is completed and the temperature of the cold-rolled sheet reaches 750°C is defined as t=E seconds, the integer part of E is defined as EI, and T t (°C) is the average temperature of the cold-rolled sheet from (t-1) seconds to t seconds.

9. The method for producing a high-strength steel plate according to claim 8; In the cooling step, a plating step of performing a plating treatment on the cold-rolled sheet; A method for producing a high-strength plated steel sheet having the above-mentioned steps.

10. A method for manufacturing a component, comprising subjecting the high-strength steel plate according to claim 1 or 2 to one or both of forming and joining to form a component.

11. A method for manufacturing a component, comprising subjecting the high-strength plated steel sheet according to claim 3 to one or both of forming and joining processes to form the component.

Images