High-strength steel sheet and method for producing same, high-strength plated steel sheet and method for producing same, member and method for producing same, and automobile skeleton structure component or automobile reinforcement component

Abstract

The present invention provides a high-strength steel sheet which has excellent component strength, ductility, stretch flangeability, and delayed fracture resistance of a stress-loaded shear end surface part. A high-strength steel sheet according to the present invention has a specific component composition that contains one or both of Mo and V satisfying formula (1), and a specific steel structure at 1 / 4 the sheet thickness. This high-strength steel sheet contains Ti at 1 / 4 the sheet thickness, the number density of the composite carbides that contain one or both of Mo and V is 1.0 × 1019 or more per m3, and the average grain size of the prior austenite grains is 3.00 µm or less.　(1): 0.0010 ≤ [%Mo] / 95.94 + [%V] / 50.94 ≤ 0.0200 In the formula, [%Mo] and [%V] represent the contents of Mo and V, respectively.

Description

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

[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.

[0002] CO2 emissions from lighter vehicles ２ To achieve both reduced emissions and improved crashworthiness, the strength of automotive steel sheets is being increased, and new regulations are being introduced one after another. As a result, the use of high-strength steel sheets with a tensile strength (TS) of 1180 MPa or more is increasing in the main structural parts that form the framework of the automobile cabin.

[0003] High strength steel sheets used in automotive reinforcing parts 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 are also required to have 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 delayed fracture, in which components suddenly break due to hydrogen penetration in the atmospheric corrosive environment in which an automobile is traveling, can occur. Automotive steel sheets are subjected to stress during pressing and assembly, and there is a risk that hydrogen will subsequently penetrate into the steel sheet from the environment. Because delayed fracture is likely to occur at sheared edges, there is a need to improve the delayed fracture resistance of stress-loaded sheared edges. In order to increase the application rate of high-strength steel sheets to 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 having 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.

[0006] International Publication No. 2023 / 026819

[0007] However, in order to increase the application rate of high-strength steel sheets to automobile 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 is 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 in delayed fracture resistance of 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 is 900 MPa or more.

[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] A steel sheet having a component composition containing, by 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), with the balance consisting of Fe and unavoidable impurities; and a steel structure having, at a 1 / 4 position in the plate thickness direction, a total area fraction of martensite and bainite of 70% or more, an area fraction of ferrite of 20% or less, and an area fraction of retained austenite of 15% or less, wherein, at the 1 / 4 position in the plate thickness direction, 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 A high-strength steel plate having an average grain size of prior austenite grains of 3.00 μm or less, wherein 0.0010≦[%Mo] / 95.94+[%V] / 50.94≦0.0200 (1), where [%Mo] and [%V] are the contents of Mo and V, respectively.

[0016] [2] The 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. The high-strength steel plate according to the above [1], containing at least one element selected from the group consisting of: 0.200% or less of Si.

[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 reinforcing part, comprising the member described in [4] above.

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

[0022] [8] A method comprising the steps of: a heating step of heating a steel slab having the chemical composition according to the above [1] or [2] to 1150°C or higher; a hot rolling step of subjecting the steel slab to rough rolling with a rolling reduction of 13% or more in a temperature range of 1050°C or higher and four or more passes, followed by finish rolling, to obtain a hot-rolled sheet; a cold rolling step of subjecting the hot-rolled sheet to pickling and cold rolling to obtain a cold-rolled sheet; an annealing step of annealing the cold-rolled sheet; and a cooling step of cooling the cold-rolled sheet to room temperature to obtain a high-strength steel sheet, wherein in the annealing step, 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, and the maximum temperature of the cold-rolled sheet is Ac ３ ℃ or more (Ac ３ +200) ° C. or less, and the temperature of the cold-rolled sheet is in a temperature range T2 ３ a retention condition from the time when the temperature reaches 550°C through the maximum temperature reached to the time when the temperature reaches 750°C satisfies the following formula (2), and in the cooling step, an average cooling rate v3 in a temperature range T3 of 550°C or higher and 750°C or lower is 4°C / s or higher, and a retention time t4 in a temperature range T4 of 350°C or higher and 550°C or lower is 600 seconds or shorter. In the annealing step, the temperature of the cold-rolled sheet is Ac ３The 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.

[0023] [9] A method for producing a high-strength plated steel sheet, comprising: the method for producing a high-strength steel sheet according to the above [8]; and a plating step of performing a plating treatment on the cold-rolled sheet in the cooling step.

[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 a 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.

[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.

[0027] 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.

[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 at a quarter-thickness position, and the number density of composite carbides containing either or both of Mo and V is 1.0 × 10 １９ pieces / m ３The steel is characterized in that the average grain size of prior austenite grains is 3.00 μm or less.

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

[0031] [C: 0.070% or more and 0.390% or less] C is one of the important basic components of high-strength steel plates 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 significantly increases, 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 ratios of ferrite, martensite, and bainite. It also embrittles grain boundaries, affecting delayed fracture resistance. If the Mn content is too low, the area ratio of ferrite increases and the area ratios of martensite and bainite decrease, making it difficult to achieve a TS of 1180 MPa or more. 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] P is an element that segregates at prior austenite grain boundaries and embrittles the grain boundaries. Therefore, if P is contained in excess, the delayed fracture resistance of the stress-loaded sheared edge portion decreases. 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. Therefore, if S is contained in excess, the delayed fracture resistance of the stressed sheared edge portion decreases. Therefore, the S content is set to 0.0200% or less, and preferably 0.0050% or less. On the other hand, although there is no particular lower limit for the S content, 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 performs sufficient deoxidation and reduces inclusions in steel. If the Al content is too high, the area ratio of ferrite increases and the area ratios 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, in order to perform 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] N exists as coarse nitrides and becomes the starting point of delayed fracture. Therefore, if N is contained in excess, the delayed fracture resistance of the stressed sheared edge portion decreases. 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 constraints on production technology, the N content is preferably set to 0.0001% or more.

[0038] [O: 0.0100% or less] O exists as an oxide and becomes the starting 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 composite carbides with V and Mo, and improves delayed fracture resistance by refining prior austenite grains and suppressing dislocation slippage through the austenite pinning effect. If the Ti content is too low, the number density of the composite carbides decreases, the prior austenite grains become coarse, and their average grain size increases, resulting in a decrease in the delayed fracture resistance of 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 initiation points of delayed fracture, thereby decreasing the delayed fracture resistance of 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: Satisfying Formula (1)] In the present invention, Mo and V are important elements that form composite carbides with Ti and improve delayed fracture resistance by refining prior austenite grains and suppressing dislocation slippage through the austenite pinning effect. To generate 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 results in the precipitation of a large amount of composite carbides containing Ti and one or both of Mo and V in the austenite region. This results in the formation of fine composite 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 starting point 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) where [%Mo] and [%V] are the contents of Mo and V, respectively.

[0042] The content of Mo and V is preferably 1.00% or less, and the content of Mo and V is preferably 0.05% or more.

[0043] The high-strength steel sheet has a 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 plate may further contain, in 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 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 0.001% or more.

[0046] [Ta: 0.10% or less] [W: 0.10% or less] When 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 starting point of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge portion 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 portion is not reduced. Therefore, when B is contained, the B content is 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 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 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 preferably 1.00% or less, and 0.80% or less, respectively. 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 0.01% or more, respectively.

[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 stress-loaded sheared edge portion is not reduced. Therefore, when Co is contained, the Co content is 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 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 stress-loaded sheared edge portion is not reduced. Therefore, when Cu is contained, the Cu content is 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 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 will not become the starting point of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge portion 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 portion is not reduced. Therefore, when Sb is contained, the Sb content is 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 0.001% or more.

[0053] [Ca: 0.0100% or less] [Mg: 0.0100% or less] [REM: 0.0100% or less] When 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 initiation sites for 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] When 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 point of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge portion is not reduced. Therefore, when Zr or Te is contained, the Zr or Te content is set to 0.100% or less, and preferably 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 set to 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 starting point of delayed fracture, so the delayed fracture resistance of the stress-loaded sheared edge portion is not reduced. Therefore, when Hf is contained, the Hf content is 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 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 stress-loaded sheared edge portion is not reduced. Therefore, when Bi is contained, the Bi content is 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 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 preferred 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 Fraction 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 formed above the Ms point and composed of angular bainitic ferrite, iron-based carbides, and retained austenite. In the present invention, it includes upper bainite and lower bainite. By including martensite and bainite, a TS of 1180 MPa or more 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. On the other hand, 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] [Area Fraction of Ferrite: 20% or Less] Ferrite is 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. In addition, it becomes the starting point of plastic deformation, so 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 limitation on the lower limit of 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 the 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 then etched using 1 volume % nital to reveal the structure. Next, a portion of the observation surface of the sample corresponding to 1 / 4 of the plate thickness of the steel sheet is observed at a magnification of 3000 times using a scanning electron microscope (SEM) under 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 tissue is divided by the measured area to obtain a value, and the average value for the three fields of view is calculated to obtain the area ratio of each tissue.

[0062] In the SEM image, martensite is a structure with a hierarchical structure with fine internal irregularities. Ferrite is a gray, flat structure region that does not contain carbides. Bainite is a structure region composed of gray, angular bainitic ferrite and iron-based carbides that exhibit white contrast. Therefore, martensite, ferrite, and bainite can be distinguished from each other.

[0063] [Area Fraction 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 fraction of retained austenite can be measured as follows. First, a steel sheet is ground so that the 1 / 4 position of the sheet thickness becomes the measurement surface, and then the steel sheet is chemically polished by an additional 0.1 mm to obtain a sample. The measurement surface of the obtained sample is measured using an X-ray diffractometer with a Co Kα radiation source to measure 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. The intensity ratios of the integrated reflection intensities of each plane of fcc iron to the measured integrated reflection intensities of each plane of bcc iron are calculated, yielding a total of nine intensity ratios. The average of the nine intensity ratios is taken as the volume fraction of retained austenite. In the present invention, the volume fraction of retained austenite is considered to be the area fraction of retained austenite.

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

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

[0067] <Number density of composite carbide containing Ti and further containing one or both of Mo and V: 1.0 × 10 １９ pieces / m ３The composite carbide containing Ti and one or both of Mo and V improves delayed fracture resistance, and therefore 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 one or both of Mo and V is too low, the prior austenite grains become coarse and the 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 or more, 5.0 × 10 １９ pieces / m ３ More preferably, 1.0 × 10 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 one or both of Mo and V can be measured as follows. First, a steel plate is ground so that the observation surface is at 1 / 4 of the plate 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 one 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 finer prior austenite grains suppress hydrogen accumulation at 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. For this reason, 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 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, and then finish-polished using colloidal silica. Next, a location on the observation surface of the sample corresponding to a position 1 / 4 of the plate thickness of the steel sheet is measured over three fields of view using an electron backscatter diffraction (EBSD) method attached to a scanning electron microscope (SEM) under conditions of an acceleration voltage of 15 kV and a step size of 20 nm with a resolution of 10 μm × 10 μm. Using OIM Analysis (manufactured by EDAX), matrix reconstruction is performed to obtain three fields of view of prior austenite grain maps. 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 number of prior austenite grain boundaries is counted. The average grain size of the prior austenite grains is determined by dividing the length of the test line by the number of intersections.

[0072] <Characteristics of High-Strength Steel Plate> [Tensile Strength: 1180 MPa or More] The tensile strength (TS) of the high-strength steel plate obtained by the present invention is 1180 MPa or more, preferably 1470 MPa or more. On the other hand, the tensile strength of the high-strength steel plate is generally 2000 MPa or less.

[0073] [Yield ratio: 70% or more] The yield ratio (YR) of the high-strength steel plate obtained by the present invention is 70% or more, preferably 72% or more, and more preferably 79% or more. On the other hand, the yield ratio of the high-strength steel plate 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. On the other hand, 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 －１A 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] [Critical hole expansion ratio: 25% or more] The critical hole expansion ratio (λ) of the high-strength steel sheet obtained by the present invention is 25% or more, preferably 28% or more. On the other hand, the critical hole expansion ratio of the high-strength steel sheet is generally 60% or less.

[0077] The critical hole expansion ratio can be measured by a 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 having 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 The initial hole diameter is measured as D [mm]. ０ [mm], the limit hole expansion ratio (λ) [%] is calculated from the following formula (5): λ = {(D ｆ -D ０ ) / D ０} × 100 ... (5)

[0078] [Excellent delayed fracture resistance at stress-loaded sheared edge] The high-strength steel plate obtained by the present invention has excellent delayed fracture resistance at stress-loaded sheared edge. That is, in the delayed fracture test described below, the fracture limit stress of the high-strength steel plate is 900 MPa or more, preferably 1050 MPa or more. On the other hand, the fracture limit stress of the high-strength steel plate 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 under conditions of 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 the presence or absence of cracks is confirmed. 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 has 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 plating layer is not particularly limited and may be a general composition. For example, when the plating layer is a hot-dip galvanized layer or a galvannealed hot-dip galvanized layer, the plating layer may contain 20 mass % or less of Fe, 0.001 to 1.0 mass % of Al, and further contain 0 mass % or more and 3.5 mass % or less 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 balance being Zn and unavoidable impurities.

[0082] When the plating layer is a hot-dip galvanized layer, the coating weight of the plating layer per side is 20 g / m ２On 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] (Member) 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 of a stressed shear end face, and therefore the member can be suitably used particularly for automotive frame structural parts or automotive reinforcing parts.

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

[0086] (Method for manufacturing high-strength steel plate) 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 Step> [Slab Heating Temperature: 1150°C 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. As a result, the number density of composite carbides containing Ti and one or both of Mo and V in the final structure decreases. 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 is 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 once cooled to room temperature, and then heated again 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. In addition, the steel slab subjected to rough rolling is called a rough-rolled sheet.

[0090] [Number of passes with a rolling reduction of 13% or more in a temperature range of 1050°C or higher is 4 or more] If the number of passes with a rolling reduction of 13% or more in a temperature range of 1050°C or higher during rough rolling is small, Ti, Mo, or V will be unevenly distributed within the steel sheet, resulting in a decrease in the number density of composite 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 with a rolling reduction of 13% or more in a temperature range of 1050°C or higher is 4 or more, and preferably 5 or more. While there is no particular upper limit to the number of passes, the number of passes is generally 10 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 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 in 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 state of austenite is reduced, the development of abnormal structures elongated in the rolling direction is suppressed, and workability is excellent.

[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 passability 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 Step> The hot-rolled sheet is then subjected to pickling and cold rolling to obtain a cold-rolled sheet (cold rolling step). 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 hot-rolled sheet after pickling 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 step> Next, the cold-rolled sheet is annealed (annealing step). Figure 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 to a maximum temperature, and then cooling the cold-rolled sheet to 750°C. When heating the cold-rolled sheet, the Ac ３ The temperature range below ℃ is T1, Ac ３ ℃ or more (Ac ３ +200) ° C. or less is defined as T2. Here, the maximum temperature of the cold-rolled sheet is Ac ３ ℃ or more (Ac ３ +200° C. or less. Although FIG. 1 shows an example in which the temperature is maintained at the maximum temperature, the present invention is not limited to this.

[0097] [Ac above 600℃ ３ Average heating rate v1 in the temperature range T1 below 600 ° C.: 1.5 ° C. / s or more] In the annealing process, ３ If the average heating rate v1 in the temperature region T1 below 100°C is too small, composite carbides containing Ti and one or both of Mo and V will precipitate non-uniformly in the unrecrystallized structure and will then coarsen during retention 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 region T1 should be 1.5°C / s or more, and preferably 2.0°C / s or more. On the other hand, there is no particular upper limit to the average heating rate v1, but from the viewpoints of operability and damage to the furnace body, the average heating rate v1 is preferably 100°C / s or less.

[0098] In addition, Ac ３ The point can be calculated using the following formula (6): Ac ３ Point (°C) = 910 - 203 x [%C] 1 / 2 + 44.7 x [%Si] - 30 x [%Mn] + 700 x [%P] + 400 x [%Al] + 400 x [%Ti] + 104 x [%V] + 13.1 x [%W] - 11 x [%Cr] + 31.5 x [%Mo] - 15.2 x [%Ni] - 20 x [%Cu] (6) Here, [%X] represents the content (mass%) of component element X in the steel, and is set to 0 if not contained.

[0099] [Maximum temperature of cold-rolled sheet: Ac ３ ℃ or more (Ac ３ +200) ° C or less] In order to advance the reverse transformation from ferrite to austenite and obtain a sufficient amount of martensite and bainite, the maximum temperature is Ac ３ ° C. or more. On the other hand, the maximum temperature of the cold-rolled sheet is (Ac ３ +200) ° C., the number density of the complex carbides decreases and the prior austenite grains also become coarse. ３ +200)°C or less.

[0100] [Ac ３ ° C. through the maximum temperature reached to 750 ° C.: retention condition: formula (2) is satisfied] 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): P≦30000 (2) where 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.

[0101] Ac in the annealing process ３ In steel sheets that have reached the austenite single-phase region, prior austenite grains are kept fine due to the pinning effect of composite carbides containing Ti and one or both of Mo and V. On the other hand, the constituent atoms of the composite carbides (particularly Ti) diffuse, causing the composite 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 Ti diffusion phenomenon in the austenite region. The parameter P is calculated from the temperature and time during the annealing process, and by appropriately controlling the value of the parameter P, the number density of the composite carbides and the average grain size of the prior austenite grains can be controlled.

[0102] If the amount of P is too large, the composite carbides become coarse and the number density of the composite 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 set to 30,000 nm. ２ or less, 25,000 nm ２ Preferably, 20,000 nm or less 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 1500 nm is preferable. ２ More preferably, 2000 nm or more ２ The above is even more preferable.

[0103] <Cooling Step> Next, the cold-rolled sheet is cooled to obtain a high-strength steel sheet (cooling step). As shown in Fig. 1 , in the cooling step, a temperature range of 550°C or higher and 750°C or lower is designated as T3, and a 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 of 550°C or higher and 750°C or lower: 4°C / s or higher] In the cooling step, 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] [Residence time t4 in temperature region T4 of 350°C or higher and 550°C or lower: 600 seconds or less] In the cooling step, the temperature region T4 of 350°C or higher and 550°C or lower is a temperature region in which carbon enrichment in untransformed austenite due to bainite transformation progresses, and retained austenite is generated. If the residence time t4 in the temperature region 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 (Preferred 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 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. When the reheating treatment is not performed, the cold-rolled sheet is cooled to room temperature in the cooling step.

[0107] (Method for manufacturing high-strength plated steel sheet) Next, a method for manufacturing a high-strength plated steel sheet according to one embodiment of the present invention will be described. The method for manufacturing a high-strength plated steel sheet includes the above-mentioned method for manufacturing 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 a 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. In addition, the alloying temperature 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, the upper limit of the reduction ratio of the skin-pass rolling is not particularly limited, but from the viewpoint of productivity, the reduction ratio is preferably 1.50% or less. The skin-pass rolling may be performed online or offline. The skin-pass rolling may be performed at a desired reduction ratio in one go, or may be performed in several steps.

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

[0113] (Method for manufacturing a member) Next, a method for manufacturing a member 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 to form a member. The forming and joining can be performed 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.

[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 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]

[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 having a sheet thickness of 1.4 mm was obtained. The obtained cold-rolled sheet was subjected to an annealing treatment, a cooling treatment, and, if necessary, a reheating treatment under the conditions shown in Table 2, and further subjected to a plating treatment in some cases 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 obtain a high-strength plated 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 balance 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 coating layer) on both sides to obtain 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 coating layer per side was 45 g / m ２ The composition of the formed alloyed hot-dip galvanized layer contained 7 to 15 mass % of Fe, 0.1 to 1.0 mass % of Al, and the balance being Fe and unavoidable impurities.

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

[0123]

[0124] The area ratios of martensite, bainite, ferrite, retained austenite, and the remaining structure were determined at a quarter-thickness position of the obtained steel plate by the methods described above. Furthermore, the number density of composite carbides containing Ti and one or both of Mo and V, and the average grain size of prior austenite grains were measured at a quarter-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]

[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.

[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. A steel sheet having a composition containing, by 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), with the balance consisting of Fe and unavoidable impurities; and a steel structure having, at a 1 / 4 position in the plate thickness direction, a total area fraction of martensite and bainite of 70% or more, an area fraction of ferrite of 20% or less, and an area fraction of retained austenite of 15% or less, wherein, at the 1 / 4 position in the plate thickness direction, 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 A high-strength steel plate having an average grain size of prior austenite grains of 3.00 μm or less, wherein 0.0010≦[%Mo] / 95.94+[%V] / 50.94≦0.0200 (1), where [%Mo] and [%V] are the contents of Mo and V, respectively.

2. The chemical composition further contains, 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.

2. The high-strength steel plate according to claim 1, further comprising at least one element selected from the group consisting of: Cr, Mn, and Cu; ... and Cr, Mn, Mn, and Cu, and the like.

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 the 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. An automobile frame structural part or an automobile reinforcing part, comprising the member according to claim 4.

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

8. A method for producing a high-strength steel sheet, comprising: a heating step of heating a steel slab having the chemical composition according to claim 1 or 2 to 1150°C or higher; a hot rolling step of subjecting the steel slab to rough rolling with a rolling reduction of 13% or more in a temperature range of 1050°C or higher and four or more passes, followed by finish rolling, to produce a hot-rolled sheet; a cold rolling step of subjecting the hot-rolled sheet to pickling and cold rolling to produce a cold-rolled sheet; an annealing step of annealing the cold-rolled sheet; and a cooling step of cooling the cold-rolled sheet to room temperature to produce a high-strength steel sheet, wherein in the annealing step, the steel slab is subjected to Ac at 600°C or higher. 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, and the maximum temperature of the cold-rolled sheet is Ac ３ ℃ or more (Ac ３ +200) ° C. or less, and the temperature of the cold-rolled sheet is in a temperature range T2 ３ a retention condition from the time when the temperature reaches 550°C through the maximum temperature reached to the time when the temperature reaches 750°C satisfies the following formula (2), and in the cooling step, an average cooling rate v3 in a temperature range T3 of 550°C or higher and 750°C or lower is 4°C / s or higher, and a retention time t4 in a temperature range T4 of 350°C or higher and 550°C or lower is 600 seconds or shorter. 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. A method for producing a high-strength plated steel sheet, comprising: the method for producing a high-strength steel sheet according to claim 8; and a plating step of plating the cold-rolled sheet during the cooling step.

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 processes to form the 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.

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