Steel sheet, member, and method for manufacturing the same

Abstract

The present invention provides a steel sheet having high strength and excellent delayed fracture resistance and a method for manufacturing the same. The steel sheet has a specific composition and has a structure in which the area fraction of martensite is 95 to 100%, the balance is composed of one or more of bainite, ferrite and residual austenite, the average grain size of prior austenite grains is 18 μm or less, 90% or more of the total content of Nb and Ti is present in the form of a carbonitride having an equivalent circle diameter of 100 nm or more, and the number of Nb carbonitrides and Ti carbonitrides having an equivalent circle diameter of 1.0 μm or more is 800 or more per mm 2 The steel sheet has a tensile strength of 1310 MPa or more.

Description

Technical Field

[0001] This invention relates to high-strength steel sheets and components for cold pressing used in automobiles and the like, and methods for manufacturing them. Background Technology

[0002] In recent years, with the aim of reducing the weight of automobiles and improving collision safety, there has been a push to use steel plates with a tensile strength (TS) of 1310 MPa or higher for automotive frame components. Research is underway to use steel plates with a tensile strength (TS) of 1.8 GPa or higher for components such as bumpers and anti-collision beams.

[0003] Previously, for steel plates with a tensile strength (TS) of 1310 MPa or higher, high-strength steel plates obtained by hot pressing have been used. However, from the perspective of cost and productivity, research is also being conducted on the application of high-strength steel plates obtained by cold pressing.

[0004] When high-strength steel plates with a tensile strength (TS) of 1310 MPa or higher are cold-pressed and manufactured into components, delayed fracture may occur due to the increase in residual stress within the components and the deterioration of the delayed fracture resistance characteristics caused by the steel plate itself.

[0005] Here, delayed fracture refers to the following phenomenon: when a component is placed in a hydrogen-infiltrating environment under high stress, hydrogen penetrates into the steel plate constituting the component, reducing the interatomic bonding force and causing local deformation, thereby generating microcracks, which progress and eventually lead to fracture.

[0006] As a technique to improve such delayed fracture resistance, for example, the precipitation of fine carbides that become hydrogen trapping points can improve the delayed fracture resistance. Based on this result, Patent Document 1 discloses a high-strength cold-rolled steel sheet with excellent hydrogen embrittlement resistance and workability, characterized by having a composition containing C: 0.05% to 0.30%, Si: 0% to 2.0%, Mn: greater than 0.1% and less than 2.8%, P: less than 0.1%, S: less than 0.005%, N: less than 0.01%, Al: less than 0.01% to 0.50%, and so on. The composition consists of a total content of 0.01% or more of Nb, Ti, and Zr, satisfying the formula [%C]-[%Nb] / 92.9×12-[%Ti] / 47.9×12-[%Zr] / 91.2×12>0.03, with the balance being iron and unavoidable impurities. It also has a microstructure containing 50% or more (including 100%) tempered martensite by area ratio, with the balance being ferrite. Regarding the distribution of precipitates in the tempered martensite, precipitates with an equivalent circle diameter of 1–10 nm per 1 μm... 2Precipitates containing more than 20 tempered martensite particles and an equivalent circle diameter of more than 20 nm, i.e., precipitates containing one or more of Nb, Ti, and Zr per 1 μm 2 The tempered martensite has fewer than 10 grains, and the average grain size of the ferrite surrounded by large-angle grain boundaries with a crystal orientation difference of more than 15° is less than 5 μm.

[0007] In addition, Patent Document 2 discloses a high-strength quenched and tempered steel with excellent resistance to delayed fracture. Its characteristic is that, in a steel containing C: 0.1–0.5%, Si: 0.10–2%, Mn: 0.44–3%, N ≤ 0.008%, Al: 0.005–0.1%, it contains one or more of V: ​​0.05–2.82%, Mo: 0.1% or more but less than 3.0%, Ti: 0.03–1.24%, and Nb: 0.05–0.95%, with a mass percentage ratio of V to C of 0.5 ≤ (0.18V + 0.06Mo + 0.25Ti + 0.13Nb) / C, the balance being Fe and unavoidable impurities, and the tensile strength is 1200–1600 MPa.

[0008] Existing technical documents

[0009] Patent documents

[0010] Patent Document 1: Japanese Patent No. 4712882

[0011] Patent Document 2: Japanese Patent No. 4427010 Summary of the Invention

[0012] The problem that the invention aims to solve

[0013] However, in the existing technology, it cannot be said that there is sufficient technology that can ensure high strength while having excellent resistance to delayed fracture.

[0014] The present invention was made to solve such problems, and its purpose is to provide steel plates, components and methods thereof with tensile strength of 1310 MPa or above (TS≥1310 MPa) and excellent resistance to delayed fracture.

[0015] Excellent resistance to delayed fracture refers to the following: Conventional Strain Rate Test (CSRT) specimens (test pieces with 3mm radius semicircular notches at both ends of the parallel section of a tensile test piece with a parallel section width of 12.5mm and a parallel section length of 25mm) are cut from the 1 / 4 position of the width of the steel plate at a rolling right angle to the length direction. A 10% by mass ammonium thiocyanate aqueous solution is mixed with McIlvaine buffer solution at pH 3 at a 1:1 volume ratio. The CSRT specimen is then subjected to a test at a distance of 1cm from the test piece. 2The sample was immersed in a solution (pH 3) at 20°C for 24 hours, with the liquid volume adjusted to reach 20 ml to allow hydrogen to penetrate and diffuse into the test piece. After 24 hours, a tensile test was immediately performed at a crosshead speed of 1 mm / min, and the fracture stress was measured. The fracture stress before immersion was set as σ0, and the fracture stress after hydrogen penetration and diffusion through immersion was set as σ1. When the TS was 1310 MPa or higher and less than 1500 MPa, σ1 / σ0 was ≥0.80; when the TS was 1500 MPa or higher and less than 1800 MPa, σ1 / σ0 was ≥0.50; and when the TS was ≥1800 MPa, σ1 / σ0 was ≥0.35.

[0016] Methods for solving problems

[0017] In order to solve the above problems, the inventors have conducted in-depth research and found that the delayed fracture resistance can be greatly improved by satisfying all of the following conditions.

[0018] i) The area ratio of martensite is over 95%.

[0019] ii) The average grain size of the original austenite grains (original γ grain size) is less than 18 μm.

[0020] iii) More than 90% by mass of the Nb and Ti contained therein exist as precipitates with an equivalent circle diameter of more than 100 nm.

[0021] iv) The number of Nb carbonitrides and Ti carbonitrides with an equivalent circle diameter of 1.0 μm or more is 800 per mm. 2 the following.

[0022] This invention was completed based on the above insights and further research, and its main points are as follows.

[0023] [1] A type of steel plate,

[0024] It contains, by mass%, C: 0.12% or more and 0.40% or less, Si: 1.5% or less, Mn: 1.8% or more and 4.0% or less, P: 0.03% or less, S: less than 0.0023%, sol.Al: 0.20% or less, N: 0.005% or less, B: 0.0100% or less, and one or more of Nb and Ti: totaling 0.005% or more and 0.080% or less, with the balance being Fe and unavoidable impurities.

[0025] Furthermore, it has a microstructure in which the area fraction of martensite relative to the overall microstructure is 95% or more but less than 100%, and the balance is composed of one or more of bainite, ferrite, and retained austenite.

[0026] The average grain size of the original austenite grains is less than 18 μm.

[0027] More than 90% by mass of the total Nb and Ti content exists in the form of carbonitrides with an equivalent circle diameter of more than 100 nm.

[0028] Nb carbonitrides and Ti carbonitrides with an equivalent circle diameter of 1.0 μm or more, totaling 800 per mm. 2 The following exists,

[0029] The tensile strength of the steel plate is above 1310 MPa.

[0030] [2] According to the steel plate described in [1], the average grain size of the aforementioned original austenite grains is less than 10 μm.

[0031] [3] The steel plate according to [1] or [2] has a fracture stress σ0 before immersion in a solution containing 10% by mass of ammonium thiocyanate aqueous solution and McIlvaine buffer solution at pH 3, a fracture stress σ1 after immersion in the above solution and the above tensile strength satisfying the following (A), (B) or (C).

[0032] (A) Tensile strength is 1310 MPa or more and less than 1500 MPa, and σ1 / σ0 is 0.80 or more;

[0033] (B) Tensile strength is above 1500MPa and below 1800MPa, and σ1 / σ0 is above 0.50;

[0034] (C) Tensile strength is above 1800MPa, and σ1 / σ0 is above 0.35.

[0035] [4] The steel plate according to any one of [1] to [3], wherein, as a component of the above-mentioned composition, it contains less than 0.0010% S by mass.

[0036] [5] The steel plate according to any one of [1] to [4], wherein, as a component of the above composition, it further contains, by mass %, one or both selected from Cu: less than 1.0% and Ni: less than 1.0%.

[0037] [6] The steel plate according to any one of [1] to [5], wherein, as a component of the above composition, it further contains, by mass %, one or more of the following: Cr: less than 1.0%, Mo: less than 0.3%, V: less than 0.5%, Zr: less than 0.2% and W: less than 0.2%.

[0038] [7] The steel plate according to any one of [1] to [6], wherein, as a component of the above composition, it further contains, by mass %, one or more of the following: Ca: less than 0.0030%, Ce: less than 0.0030%, La: less than 0.0030%, REM (excluding Ce and La): less than 0.0030%, and Mg: less than 0.0030%.

[0039] [8] The steel plate according to any one of [1] to [7], wherein, as a component of the above composition, it further contains, by mass %, one or both selected from Sb: less than 0.1% and Sn: less than 0.1%.

[0040] [9] The steel plate according to any one of [1] to [8], wherein the steel plate has a coating on its surface.

[0041]

[10] A component which is formed by forming and welding of a steel plate as described in any one of [1] to [9].

[0042]

[11] A method for manufacturing a steel plate, wherein,

[0043] A steel billet having any one of the compositions described in [1], [4] to [8] is heated from 1000°C to a holding temperature of 1250°C or higher using a steel billet surface thermometer at an average heating rate of 10°C / min or less, and held at the holding temperature for at least 30 minutes.

[0044] Perform hot finishing rolling with the finishing temperature set above Ar3.

[0045] The cooling process will be carried out at an average cooling rate of 40°C / second or higher within the range from the aforementioned finishing rolling temperature to 650°C.

[0046] Then, it is cooled and wound at a winding temperature below 600°C to produce hot-rolled steel sheet.

[0047] The hot-rolled steel sheet is cold-rolled to produce a cold-rolled steel sheet by cold rolling with a reduction rate of more than 40%.

[0048] Perform the following continuous annealing:

[0049] The annealing temperature was set to 800–950°C, and the cold-rolled steel sheet was heated from 700°C to the above-mentioned annealing temperature at an average heating rate of 0.4°C / second or higher.

[0050] Hold at the above annealing temperature for less than 600 seconds.

[0051] Cool from the above annealing temperature to 420°C at a first average cooling rate of 2°C / second or higher.

[0052] Cooling from 420°C to a cooling stop temperature below 280°C at a second average cooling rate of 10°C / second or higher.

[0053] Then, maintain at a holding temperature of 120–260°C for 20–1500 seconds.

[0054]

[12] According to the steel plate manufacturing method described in

[11] , the surface of the steel plate is plated after the continuous annealing described above.

[0055]

[13] A method for manufacturing a component, comprising at least one of forming and welding a steel plate manufactured by the steel plate manufacturing method described in

[11] or

[12] .

[0056] Invention Effects

[0057] According to the present invention, steel plates, components and methods thereof with high strength and excellent resistance to delayed fracture can be provided. Detailed Implementation

[0058] The embodiments of the present invention will be described below.

[0059] The steel plate of the present invention comprises, by mass percent, C: 0.12% or more and 0.40% or less, Si: 1.5% or less, Mn: 1.8% or more and 4.0% or less, P: 0.03% or less, S: less than 0.0023%, sol.Al: 0.20% or less, N: 0.005% or less, B: 0.0100% or less, and one or more of Nb and Ti: totaling 0.005% or more and 0.080% or less, with the balance being Fe and unavoidable impurities, and has the following composition: The microstructure has a martensite area fraction of 95% to 100% relative to the overall microstructure, with the remainder consisting of one or more of bainite, ferrite, and retained austenite. The average grain size (original γ grain size) of the original austenite grains (hereinafter also referred to as original γ grains) is 18 μm or less. More than 90% by mass of the total Nb and Ti content exists in the form of carbonitrides with an equivalent circle diameter of 100 nm or more, and the total number of Nb carbonitrides and Ti carbonitrides with an equivalent circle diameter of 1.0 μm or more is 800 per mm. 2 The steel plate has a tensile strength of 1310 MPa or higher.

[0060] Composition

[0061] The reasons for limiting the range of the composition of the steel plate of the present invention will be explained below. It should be noted that the % of the composition content refers to "mass %".

[0062] C: Above 0.12% and below 0.40%

[0063] Carbon (C) is added to improve hardenability, thereby obtaining a martensitic steel structure, and to ensure a tensile strength of 1310 MPa or higher (hereinafter also referred to as TS≥1310 MPa) by increasing the strength of martensite. Excessive addition of C can lead to the formation of iron carbides and grain boundary segregation, contributing to the deterioration of delayed fracture resistance. Therefore, the C content is limited to a range of 0.12% or more and 0.40% or less, which is necessary to obtain the required strength of the steel. The C content is preferably 0.37% or less, more preferably 0.34% or less.

[0064] Si: below 1.5%

[0065] Si is included as a strengthening element based on solid solution strengthening, and from the viewpoint of improving resistance to delayed fracture by suppressing the formation of film-like carbides during tempering at temperatures above 200°C. Additionally, Si is included to suppress MnS formation by reducing Mn segregation in the central part of the plate thickness. Furthermore, Si is included to suppress decarburization and debonding caused by surface oxidation during annealing on a continuous annealing line (CAL). While no lower limit is specified for the Si content, from the viewpoint of achieving the aforementioned effects, Si is preferably 0.02% or more. The Si content is preferably 0.10% or more, more preferably 0.20% or more. On the other hand, excessive Si content leads to increased segregation and deterioration of resistance to delayed fracture. It also results in a significant increase in rolling load and a decrease in toughness during hot and cold rolling. Therefore, the Si content is set to 1.5% or less (including 0%). The Si content is preferably 1.2% or less, more preferably 1.0% or less.

[0066] Mn: 1.8% or more and 4.0% or less

[0067] Mn is included to improve the hardenability of steel and obtain the desired strength by ensuring that the martensite area ratio is within a specified range. To suppress ferrite formation, Mn content is 1.8% or more. On the other hand, Mn is particularly conducive to the formation and coarsening of MnS in the central part of the plate thickness. When the Mn content exceeds 4.0%, even if S is reduced to a limit, the number and size of large MnS in the central part of the plate thickness increase, significantly deteriorating the resistance to delayed fracture. Therefore, Mn content is set to 4.0% or less. From the viewpoint of further reducing coarse MnS and improving resistance to delayed fracture, the Mn content is preferably set to 3.2% or less. More preferably, the Mn content is 2.8% or less.

[0068] P: below 0.03%

[0069] Phosphorus (P) is an element that strengthens steel; however, high P content significantly deteriorates its resistance to delayed fracture and its weldability. Therefore, the P content is set to be below 0.03%. Based on this viewpoint, the P content is preferably set to below 0.004%. While no lower limit for P content is specified, 0.002% is currently set as an industrially feasible lower limit.

[0070] S: Less than 0.0023%

[0071] S forms MnS, which significantly reduces the resistance to delayed fracture at the shear end face. Therefore, to reduce MnS, the S content needs to be set to at least less than 0.0023%. From this point of view, S is preferably set to less than 0.0010%. From the point of view of improving resistance to delayed fracture, S is more preferably set to less than 0.0004%. No lower limit is specified, but as a lower limit currently feasible in industry, it is set to 0.0002%.

[0072] sol.Al: 0.20% or less

[0073] Al is included to ensure sufficient deoxidation and reduce inclusions in the steel. There is no specific lower limit for sol.Al, but for stable deoxidation, it is preferable to set sol.Al to 0.005% or more. Furthermore, it is more preferable to set sol.Al to 0.01% or more. Sol.Al is preferably 0.02% or more. On the other hand, when sol.Al exceeds 0.20%, the cementite generated during coiling is difficult to dissolve during annealing, resulting in deterioration of the delayed fracture resistance. Therefore, sol.Al is set to 0.20% or less. Sol.Al is preferably 0.10% or less, and more preferably 0.05% or less.

[0074] N: less than 0.005%

[0075] Nitrogen (N) is an element that forms inclusions in steel such as TiN, (Nb,Ti)(C,N), and AlN, as well as carbonitride systems. Their formation deteriorates the resistance to delayed fracture. They hinder the adjustment of the steel microstructure required by this invention and adversely affect the resistance to delayed fracture at the shear face. To reduce this adverse effect, the N content is set to 0.005% or less. The N content is preferably 0.0040% or less. While no lower limit is specified, it is set to 0.0006% as a currently industrially feasible lower limit.

[0076] B: Below 0.0100%

[0077] Boron (B) is an element that improves the hardenability of steel and has the advantage of generating martensite with a specified area ratio even with low Mn content. To achieve this effect, the B content is preferably set to 0.0003% or more. Furthermore, the B content is more preferably set to 0.0008% or more. The B content is even more preferably 0.0010% or more. From the viewpoint of fixing N, B is preferably added in combination with 0.002% or more Ti. On the other hand, when the B content exceeds 0.0100%, not only is its effect saturated, but it can also sometimes delay the solid solution rate of cementite during annealing, leaving undissolved cementite residues and thus deteriorating the resistance to delayed fracture at the shear end face. Based on the above, the B content is 0.0100% or less. The B content is preferably 0.0065% or less, more preferably 0.0030% or less, and even more preferably 0.0025% or less.

[0078] One or more of Nb and Ti: totaling more than 0.005% and less than 0.080%.

[0079] Nb and Ti contribute to increased strength by refining the internal structure of martensite, and improve resistance to delayed fracture by refining the original γ grain size. From this perspective, the steel billet contains at least 0.005% of either Nb or Ti. The total content of Nb and Ti is preferably 0.010% or more, more preferably 0.020% or more. On the other hand, when the total content of either Nb or Ti exceeds 0.080%, Nb and Ti are not completely dissolved during billet reheating, leading to an increase in coarse inclusion particles such as TiN, Ti(C,N), NbN, Nb(C,N), and (Nb,Ti)(C,N), which deteriorates the resistance to delayed fracture. Therefore, the upper limit for the total content of Nb and Ti is 0.080%. The total content of Nb and Ti is preferably 0.07% or less, more preferably 0.06% or less.

[0080] The steel plate of the present invention has a composition containing the above-mentioned components, with the balance being Fe (iron) and unavoidable impurities. In particular, the steel plate of one embodiment of the present invention preferably has a composition containing the above-mentioned components, with the balance being Fe and unavoidable impurities.

[0081] Based on the above-mentioned basic components, the steel plate of the present invention may also contain the following optional elements.

[0082] It should be noted that, in this invention, when these optional components are contained at values ​​lower than the preferred lower limits described below, the element is contained as an unavoidable impurity.

[0083] Cu: below 1.0%

[0084] Cu improves the corrosion resistance of automobiles in the operating environment. Furthermore, by containing Cu, corrosion products are coated onto the steel sheet surface, inhibiting hydrogen penetration into the steel sheet. Additionally, Cu is an element incorporated when effectively utilizing waste materials as raw materials; by allowing Cu to be incorporated, recycled materials can be effectively utilized as raw materials, reducing manufacturing costs. From the above viewpoints, Cu is preferably contained at 0.01% or more, and further, from the viewpoint of improving resistance to delayed fracture, Cu is preferably contained at 0.05% or more. The Cu content is more preferably 0.10% or more. However, excessive Cu content can cause surface defects; therefore, the Cu content is preferably set to 1.0% or less. Based on the above, when Cu is present, the Cu content is set to 1.0% or less. The Cu content is more preferably 0.50% or less, and even more preferably 0.30% or less.

[0085] Ni: below 1.0%

[0086] Ni also improves corrosion resistance. Furthermore, Ni reduces surface defects that are easily formed when Cu is present. Therefore, from the above perspective, it is preferable that Ni contains 0.01% or more. More preferably, Ni content is 0.05% or more, and even more preferably 0.10% or more. However, excessive Ni content leads to uneven oxide scale formation in the furnace, causing surface defects and significantly increasing costs. Therefore, when Ni is present, the Ni content is set to 1.0% or less. More preferably, Ni content is 0.50% or less, and even more preferably 0.30% or less.

[0087] Cr: less than 1.0%

[0088] Cr can be added to improve the hardenability of steel. To achieve this effect, a Cr content of 0.01% or more is preferred. A Cr content of 0.05% or more is more preferred, and 0.10% or more is even more preferred. However, when the Cr content exceeds 1.0%, the solid solution rate of cementite during annealing is delayed, leaving undissolved cementite residue, thereby deteriorating the resistance to delayed fracture at the shear end face. Furthermore, it also deteriorates resistance to pitting corrosion. In addition, it deteriorates the chemical conversion treatment properties. Therefore, when Cr is present, the Cr content is set to 1.0% or less. Resistance to delayed fracture, pitting corrosion, and chemical conversion treatment properties all tend to deteriorate when the Cr content exceeds 0.2%. Therefore, from the viewpoint of preventing these issues, a Cr content of 0.2% or less is more preferred.

[0089] Mo: less than 0.3%

[0090] Mo can be added to improve the hardenability of steel, generate fine Mo-containing carbides that act as hydrogen trapping points, and improve resistance to delayed fracture by refining martensite. When large amounts of Nb and Ti are added, coarse precipitates are formed, which deteriorates the resistance to delayed fracture; however, the solid solution limit of Mo is relatively large compared to Nb and Ti. When Nb and Ti are added in combination, fine precipitates are formed by their combination with Mo, which has the effect of refining the microstructure. Therefore, by adding Mo in combination with a small amount of Nb and Ti, the microstructure can be refined without leaving coarse precipitates, and fine carbides can be dispersed in large quantities, thereby improving the resistance to delayed fracture. To achieve this effect, Mo is preferably 0.01% or more. The Mo content is more preferably 0.03% or more, and even more preferably 0.05% or more. However, when Mo is 0.3% or more, the chemical conversion treatment properties deteriorate. Therefore, when Mo is present, the Mo content is set to be less than 0.3%. The Mo content is preferably 0.2% or less.

[0091] V: Below 0.5%

[0092] V can be added for the purpose of improving the hardenability of steel, generating fine V-containing carbides that serve as hydrogen trapping points, and improving resistance to delayed fracture by refining martensite. To achieve this effect, the V content is preferably set to 0.003% or more. More preferably, the V content is 0.03% or more, and even more preferably 0.05% or more. However, when the V content exceeds 0.5%, castability deteriorates significantly. Therefore, when V is present, the V content is set to 0.5% or less. More preferably, the V content is 0.3% or less, and even more preferably 0.2% or less. The V content is even more preferably 0.1% or less.

[0093] Zr: below 0.2%

[0094] Zr contributes to increased strength and improved resistance to delayed fracture by refining the original γ grain size and the resulting internal martensite structure. Furthermore, it enhances strength and improves resistance to delayed fracture by forming fine Zr-based carbides / carbonitrides that act as hydrogen trapping points. Additionally, Zr improves castability. From this perspective, the Zr content is preferably set to 0.005% or more. More preferably, it is 0.010% or more, and even more preferably 0.015% or more. However, when a large amount of Zr is added, the amount of coarse ZrN and ZrS precipitates remaining due to incomplete solution treatment during the hot rolling process increases, deteriorating the resistance to delayed fracture at the shear face. Therefore, when Zr is present, the Zr content is set to 0.2% or less. More preferably, it is 0.1% or less, and even more preferably 0.04% or less.

[0095] W: below 0.2%

[0096] W contributes to increased strength and improved resistance to delayed fracture by forming fine W-based carbides / carbonitrides that act as hydrogen trapping points. From this perspective, W content is preferably 0.005% or more. More preferably, it is 0.010% or more, and even more preferably 0.030% or more. However, when W content is high, the amount of coarse precipitates remaining due to lack of solid solution during billet heating in the hot rolling process increases, deteriorating the resistance to delayed fracture at the shear end face. Therefore, when W is present, the W content is set to 0.2% or less. More preferably, it is 0.1% or less.

[0097] Ca: below 0.0030%

[0098] Ca improves the resistance to delayed fracture by fixing S in the form of CaS. To achieve this effect, it is preferable to contain 0.0002% or more of Ca. More preferably, the Ca content is 0.0005% or more, and even more preferably 0.0010% or more. However, adding a large amount of Ca deteriorates surface quality and flexibility; therefore, the Ca content is preferably 0.0030% or less. Based on the above, when Ca is present, the Ca content is set to 0.0030% or less. More preferably, the Ca content is 0.0025% or less, and even more preferably 0.0020% or less.

[0099] Ce: below 0.0030%

[0100] Ce also fixes S, thus improving resistance to delayed fracture. To achieve this effect, it is preferable to contain 0.0002% or more Ce. More preferably, the Ce content is 0.0003% or more, and even more preferably 0.0005% or more. However, adding a large amount of Ce deteriorates surface quality and flexibility; therefore, the Ce content is preferably 0.0030% or less. Based on the above, when Ce is present, the Ce content is set to 0.0030% or less. More preferably, the Ce content is 0.0020% or less, and even more preferably 0.0015% or less.

[0101] La: below 0.0030%

[0102] La also fixes S, thus improving resistance to delayed fracture. To achieve this effect, it is preferable to contain 0.0002% or more of La. More preferably, the La content is 0.0005% or more, and even more preferably 0.0010% or more. However, adding a large amount of La deteriorates surface quality and flexibility; therefore, the La content is preferably 0.0030% or less. Based on the above, when La is present, the La content is set to 0.0030% or less. More preferably, the La content is 0.0020% or less, and even more preferably 0.0015% or less.

[0103] REM: below 0.0030%

[0104] REM also fixes S, thus improving resistance to delayed fracture. To achieve this effect, it is preferable to contain 0.0002% or more REM. A more preferable REM content is 0.0003% or more, and even more preferable is 0.0005% or more. However, adding a large amount of REM deteriorates surface quality and flexibility; therefore, a REM content of 0.0030% or less is preferable. Based on the above, when REM is present, the REM content is set to 0.0030% or less. A more preferable REM content is 0.0020% or less, and even more preferable is 0.0015% or less.

[0105] It should be noted that, in this invention, REM refers to scandium (Sc) atom number 21, yttrium (Y) atom number 39, and lanthanum (La) atom number 57 to lutetium (Lu) atom number 71, excluding Ce and La. The REM concentration in this invention refers to the total content of one or more elements selected from the aforementioned REM elements.

[0106] Mg: less than 0.0030%

[0107] Mg improves resistance to delayed fracture by fixing O in the form of MgO. To achieve this effect, it is preferable to contain 0.0002% or more of Mg. More preferably, the Mg content is 0.0005% or more, and even more preferably 0.0010% or more. However, adding a large amount of Mg deteriorates surface quality and flexibility; therefore, the Mg content is preferably 0.0030% or less. Based on the above, when Mg is present, the Mg content is set to 0.0030% or less. More preferably, the Mg content is 0.0020% or less, and even more preferably 0.0015% or less.

[0108] Sb: below 0.1%

[0109] Sb inhibits surface oxidation and nitriding, thus suppressing the resulting reduction of carbon (C) and boron (B). By suppressing the reduction of C and B, the formation of ferrite in the surface layer is suppressed, which contributes to increased strength and improved resistance to delayed fracture. From this perspective, the Sb content is preferably set to 0.002% or more. The Sb content is more preferably 0.004% or more, and even more preferably 0.006% or more. However, when the Sb content exceeds 0.1%, castability deteriorates, and Sb segregation at the original γ grain boundaries deteriorates the resistance to delayed fracture at the shear end face. Therefore, the Sb content is preferably 0.1% or less. Based on the above, in the case of containing Sb, the Sb content is set to 0.1% or less. The Sb content is more preferably 0.05% or less, and even more preferably 0.02% or less.

[0110] Sn: less than 0.1%

[0111] Sn inhibits surface oxidation and nitriding, thereby suppressing the resulting decrease in the content of C and B in the surface layer. By suppressing the reduction of C and B, the formation of ferrite in the surface layer is suppressed, which contributes to increased strength and improved resistance to delayed fracture. From this point of view, the Sn content is preferably set to 0.002% or more. The Sn content is preferably 0.003% or more. However, when the Sn content exceeds 0.1%, castability deteriorates, and Sn segregation at the original γ grain boundaries deteriorates the resistance to delayed fracture at the shear end face. Therefore, in the case of Sn content, the Sn content is set to 0.1% or less. The Sn content is more preferably 0.05% or less, and even more preferably 0.01% or less.

[0112] steel structure

[0113] The steel plate of the present invention has the following steel structure.

[0114] (Composition 1) The area ratio of martensite relative to the overall structure is more than 95% and less than 100%, and the balance is composed of one or more of bainite, ferrite and retained austenite.

[0115] (Composition 2) The average grain size of the original austenite grains is less than 18 μm.

[0116] (Composition 3) contains more than 90% by mass of the total content of Nb and Ti in the form of carbonitrides with an equivalent circle diameter of more than 100 nm.

[0117] (Composition 4) Nb carbonitrides and Ti carbonitrides with an equivalent circle diameter of 1.0 μm or more, at a density of 800 per mm. 2 The following exists.

[0118] The following is an explanation of each component.

[0119] (Composition 1) The area ratio of martensite relative to the overall structure is more than 95% and less than 100%, and the balance is composed of one or more of bainite, ferrite and retained austenite.

[0120] To balance high strength (TS≥1310MPa) and excellent resistance to delayed fracture, the total martensite area fraction in the steel microstructure is set to 95% or more. More preferably, it is 99% or more, and even more preferably 100%. It should be noted that, in the case of containing martensite and bainite, the balance is ferrite and retained austenite (retained γ). The remaining components are trace amounts of carbides, sulfides, nitrides, and oxides. Furthermore, the martensite also includes martensite that has not undergone tempering, including self-tempering during continuous cooling, by holding at approximately 150°C or above for a certain period of time. It should be noted that the martensite area fraction can also be 100% without any balance.

[0121] (Composition 2) The average grain size of the original austenite grains is less than 18 μm.

[0122] The delayed fracture surface of steel with martensite as the parent phase is mostly a grain boundary fracture, and it is believed that the initiation point of delayed fracture and the crack propagation path in the early stage of delayed fracture are at the original γ grain boundaries. By refining the original γ grain size, grain boundary fracture is suppressed, and the resistance to delayed fracture is significantly improved. As a mechanism, it is believed that due to the refinement of the original γ grain size, the volume fraction of the original γ grain boundaries increases, and the concentration of grain boundary embrittlement elements such as P at the grain boundaries decreases. From the viewpoint of resistance to delayed fracture, the average grain size of the original austenite grains (average original γ grain size) is 18 μm or less. This average grain size is preferably 15 μm or less, more preferably 10 μm or less, further preferably 7 μm or less, and even more preferably 5 μm or less.

[0123] (Composition 3) contains more than 90% by mass of the total content of Nb and Ti in the form of carbonitrides with an equivalent circle diameter of more than 100 nm.

[0124] It is believed that Nb and Ti precipitate during hot rolling and coiling processes, and through the pinning effect during coiling and annealing, refine the original γ-particle size into fine precipitates smaller than 50 nm. These precipitates effectively suppress delayed fracture by non-diffusively capturing hydrogen in the steel at the parent phase interface. However, for cold-pressed steel sheets with continuous annealing processes, the Nb-based and Ti-based precipitates become coarser. Therefore, to capture enough hydrogen to suppress delayed fracture, a large amount of Nb and Ti needs to be added. Furthermore, precipitates with hydrogen-capturing capabilities increase the amount of hydrogen penetrating the steel, potentially worsening the delayed fracture characteristics.

[0125] On the other hand, the inventors have discovered that by utilizing Nb-based and Ti-based precipitates larger than 100 nm, which have long been considered ineffective in suppressing delayed fracture due to their lack of hydrogen capture capability, delayed fracture characteristics are significantly improved. This effect is achieved by ensuring that at least 90% by mass of the combined Nb and Ti content in the steel is composed of Nb and Ti forming carbonitrides larger than 100 nm. Although the mechanism is not necessarily clear, the present invention suggests that Nb and Ti-based carbonitrides larger than 100 nm dispersed in the steel influence the crack progression of delayed fracture, thereby suppressing delayed fracture.

[0126] Based on the above, in the steel plate of the present invention, 90% or more of the total content of Nb and Ti exists in the form of carbonitrides with an equivalent circle diameter of 100 nm or more.

[0127] Furthermore, while there is no specific upper limit for the size of carbonitrides, the newly precipitated Nb and Ti precipitates during the hot rolling and coiling processes are mostly below 500 nm. Therefore, in this invention, the aforementioned Nb and Ti carbonitrides are preferably set to carbonitrides with an equivalent circle diameter of 100 nm or more and 500 nm or less.

[0128] (Component 4) Nb carbonitrides and Ti carbonitrides with an equivalent circle diameter of 1.0 μm or more, totaling 800 per mm. 2 The following exists.

[0129] In steels with sufficiently fine original γ grain size and suppressed grain boundary fracture, inclusions larger than 1.0 μm become the initiation point for delayed fracture. Therefore, reducing inclusions larger than 1.0 μm is crucial. Nb and Ti precipitates have high melting temperatures and constitute a particularly high proportion of inclusions larger than 1.0 μm. Therefore, in this invention, to improve resistance to delayed fracture, the total number of Nb carbonitrides and Ti carbonitrides with an equivalent circle diameter of 1.0 μm or larger is set to 800 per mm. 2 The following is preferred: 100 pieces / mm. 2 The following is a further preferred value: 50 pieces / mm 2 It should be noted that, in most cases, the total number of Nb carbonitrides and Ti carbonitrides with an equivalent circle diameter of 1.0 μm or more is 5 per mm. 2 above.

[0130] The methods for measuring each component in the above-mentioned steel structure are explained.

[0131] The area ratios of martensite, bainite, and ferrite were determined as follows: The L-section of the steel plate (parallel to the rolling direction and perpendicular to the plate surface) was ground and etched with a nitric acid-ethanol solution. At a position 1 / 4 thickness from the plate surface, four fields of view were observed using a SEM at 2000x magnification. Image analysis was performed on the photographs to determine the area ratios of martensite, bainite, and ferrite. Here, martensite and bainite refer to the gray or white structure observed in the SEM. Ferrite, on the other hand, is the region appearing as black contrast in the SEM. It should be noted that while martensite and bainite contain trace amounts of carbides, nitrides, sulfides, and oxides, these are difficult to eliminate; therefore, the area ratio of the region containing these substances is used as the area ratio.

[0132] Here, bainite has the following characteristics: it has an aspect ratio of 2.5 or greater and a plate-like morphology; it is a slightly darker structure compared to martensite. The width (minor axis) of these plates is 0.3–1.7 μm. The distribution density of carbides with diameters of 10–200 nm within the bainite body is 0–3 per μm. 2 .

[0133] The determination of retained austenite (retained γ) is as follows: The surface 200 μm of the steel plate was chemically ground with oxalic acid, and the retained austenite (retained γ) was determined by X-ray diffraction intensity method using the plate surface as the object. It was calculated based on the integrated intensity of the diffraction peaks of (200)α, (211)α, (220)α, (200)γ, (220)γ, and (311)γ measured using Mo-Kα rays.

[0134] The average grain size of the original austenite grains (original γ grain size) is determined as follows: After grinding the L-section of the steel plate (a section parallel to the rolling direction and perpendicular to the steel plate surface), it is etched using a reagent that corrodes the original γ grain boundaries (e.g., a saturated picric acid aqueous solution or a solution to which ferric chloride has been added). Four fields of view are observed at 500x magnification at a position 1 / 4 thickness from the steel plate surface. In the obtained photographs, 15 lines are drawn at intervals of at least 10 μm in both the plate thickness direction and the rolling direction, and the number of intersections between the grain boundaries and the lines is counted. Furthermore, by multiplying the value obtained by dividing the line length by the number of intersections by 1.13, the original γ grain size (average grain size of the original austenite grains) can be determined.

[0135] The ratio of Nb and Ti in the Nb and Ti contained in the carbonitrides that form equivalent circular diameters of 100 nm or more can be determined by the following method.

[0136] After electrolyzing a specified amount of the sample in an electrolyte, the sample sheet was removed from the electrolyte and immersed in a dispersible solution. Next, the precipitates contained in the solution were filtered using a filter with a pore size of 100 nm. The precipitates captured by this 100 nm filter were carbonitrides with a diameter (equivalent circle diameter) of 100 nm or more. The Nb and Ti contents of the residue on the filter and the filtrate after filtration were analyzed to determine the Nb and Ti contents in carbonitrides with a diameter of 100 nm or more and in carbonitrides with a diameter of less than 100 nm. Inductively coupled plasma (ICP) emission spectrometry could be used in the analysis. Then, the ratio of the total Nb and Ti content in carbonitrides with a diameter of 100 nm or more to the total Nb and Ti content in the steel was calculated.

[0137] Nb carbonitrides and Ti carbonitrides per 1 mm 2 The number (distribution density) can be calculated as follows: After grinding the L-section of the steel plate (the section parallel to the rolling direction and perpendicular to the surface of the steel plate), without corrosion, continuously take 2mm SEM images in the region from 1 / 5 to 4 / 5 of the plate thickness, that is, from 1 / 5 of the plate thickness from the surface of the steel plate, sandwiched between the center of the plate thickness and the 4 / 5 position. 2The number of such carbonitrides in the above area was determined by SEM images, which allowed for the calculation of the number of Nb and Ti carbonitrides per 1 mm. 2 The number (distribution density) of inclusions. Here, the SEM image is preferably a reflectance electron image. In addition, the magnification is set to 2000x. However, if it is difficult to accurately determine the size of the precipitates at 2000x, the inclusion particles can be magnified to 10000x to draw the carbonitrides described above.

[0138] Tensile strength (TS): ≥1310MPa

[0139] The deterioration of resistance to delayed fracture becomes significant when the tensile strength of the raw material is 1310 MPa or higher. One of the features of this invention is that the resistance to delayed fracture remains good even at 1310 MPa or higher. Therefore, in this invention, the tensile strength is set to 1310 MPa or higher. The tensile strength of the steel sheet of this invention can be set to 2100 MPa or lower.

[0140] Regarding tensile strength, it can be determined by cutting JIS No. 5 tensile test pieces at 1 / 4 of the roll width with the rolling right angle as the length direction, and then performing the tensile test according to JIS Z2241.

[0141] The steel plate of the present invention preferably has the following fracture stress σ0 before immersion in a solution containing 10% by mass of ammonium thiocyanate aqueous solution and McIlvaine buffer solution at pH 3, fracture stress σ1 after immersion in the above solution, and tensile strength satisfying (A), (B) or (C).

[0142] (A) Tensile strength is 1310 MPa or more and less than 1500 MPa, and σ1 / σ0 is 0.80 or more;

[0143] (B) Tensile strength is above 1500MPa and below 1800MPa, and σ1 / σ0 is above 0.50;

[0144] (C) Tensile strength is above 1800MPa, and σ1 / σ0 is above 0.35.

[0145] Regarding the aforementioned fracture stresses σ0 and σ1, they can be obtained by cutting a conventional strain rate test piece (CSRT) from a position 1 / 4 of the width of the steel plate with the rolling right angle as the length direction, and using this test piece. The CSRT test piece can be configured as a tensile test piece with semi-circular notches of 3 mm radius at both ends of the parallel portion, which has a parallel portion width of 12.5 mm and a parallel portion length of 25 mm.

[0146] A 10% (w / w) aqueous solution of ammonium thiocyanate can be mixed with McIlvaine buffer at pH 3 in a 1:1 ratio. CSRT can then be applied to test strips at 1 cm intervals. 2 The sample was immersed in a solution at 20°C for 24 hours, with the liquid volume adjusted to reach 20 ml to allow hydrogen to penetrate and diffuse into the test piece. Immediately after 24 hours, a tensile test was performed at a crosshead speed of 1 mm / min, and the fracture stress σ1 was measured. Alternatively, the fracture stress obtained under the same conditions without the above immersion can be defined as σ0.

[0147] The steel plate of the present invention can be a steel plate with a coating on its surface. The coating can be a Zn coating or a coating of other metals. Alternatively, it can be any one of hot-dip galvanizing or electroplating.

[0148] Next, the method for manufacturing the steel plate of the present invention will be described.

[0149] The method for manufacturing the steel plate of the present invention is as follows: a steel billet having the above-described composition is heated from 1000°C to a heating holding temperature of 1250°C or higher at an average heating rate of 10°C / min or lower using a steel billet surface thermometer, and held at this heating holding temperature for 30 minutes or more. Then, hot finishing rolling is performed with the finishing rolling temperature set to Ar3 or higher. Cooling is then performed with an average cooling rate set to 40°C / second or higher in the range from the finishing rolling temperature to 650°C. Finally, the steel plate is cooled and coiled at a coiling temperature of 600°C or lower to produce a hot-rolled steel plate. Cold-rolled steel sheets are produced by cold rolling with a reduction rate of 40% or more, and then subjected to continuous annealing as follows: the annealing temperature is set to 800-950°C, the cold-rolled steel sheet is heated from 700°C to the annealing temperature at an average heating rate of 0.4°C / second or more, held at the annealing temperature for 600 seconds or less, cooled from the annealing temperature to 420°C at a first average cooling rate of 2°C / second or more, cooled from 420°C to a cooling stop temperature of 280°C or less at a second average cooling rate of 10°C / second or more, reheated as needed, and then held at a holding temperature of 120-260°C for 20-1500 seconds.

[0150] Hot rolling

[0151] In the heating of the billet before hot rolling, by setting the average heating rate to 10°C / min or less until the heating and holding temperature is from 1000°C to 1250°C or higher, solid solution of sulfides can be promoted, thereby reducing the size and number of inclusions. Since Nb and Ti have high melting temperatures, by setting the heating and holding temperature to 1250°C or higher using a billet surface thermometer and setting the holding time to 30 minutes or more, solid solution of Nb and Ti can be promoted, thereby reducing the size and number of inclusions. The heating and holding temperature is preferably set to 1300°C or higher, more preferably 1350°C or higher. It should be noted that the average heating rate from 1000°C to 1250°C or higher is preferably 2°C / min or higher. Furthermore, the heating and holding temperature using a billet surface thermometer is preferably 1380°C or lower. The holding time of the billet at the heating and holding temperature is preferably 250 minutes or lower.

[0152] Here, the average heating rate is defined as "(temperature at the end of billet heating (°C) - temperature at the beginning of billet heating (°C)) / heating time from the start to the end of heating (minutes)".

[0153] In hot finishing rolling, when the finishing temperature is below the Ar3 point, ferrite is formed, and stress concentration occurs at the ferrite interface of the final product, thus promoting delayed fracture. Therefore, the finishing temperature (FT) is set above the Ar3 point.

[0154] During cooling after hot finishing rolling, an average cooling rate of 40°C / second or higher is set within the range from the finishing rolling temperature to 650°C. When the average cooling rate is less than 40°C / second, the desired resistance to delayed fracture is not achieved due to the coarsening of Nb and Ti carbonitrides and the increase in carbonitrides with an equivalent circle diameter of 1.0 μm or higher. Preferably, the average cooling rate is 250°C / second or lower, more preferably 200°C / second or lower.

[0155] It should be noted that the average cooling rate is defined as "(temperature at the start of cooling (finishing temperature) (°C) - temperature at the end of cooling (°C) (650°C)) / cooling time from the start to the end of cooling (seconds)".

[0156] When the winding temperature exceeds 600℃, only the coarsening of Nb and Ti-based precipitates in the fine austenite region occurs. Therefore, the increase in coarse precipitates reduces the delayed fracture characteristics. Thus, the winding temperature is set below 600℃.

[0157] It should be noted that point Ar3 is obtained as follows.

[0158] Ar3 point (℃)=910-310×[C]-80×[Mn]-20×[Cu]-15×[Cr]-55×[Ni]-80×[Mo]

[0159] (In the above formula, [M] is the content (mass%) of element M in the steel billet, and the value of elements that are not present is set to zero (0).)

[0160] cold rolling

[0161] In cold rolling, setting the reduction rate (cold rolling ratio) to 40% or more can stabilize the recrystallization behavior and texture orientation during subsequent continuous annealing. When it is less than 40%, some of the austenite grains during annealing may become coarse, leading to a decrease in strength. Furthermore, the cold rolling ratio is preferably 80% or less. More preferably, it is 70% or less.

[0162] Continuous annealing

[0163] For cold-rolled steel sheets, annealing and tempering treatment as needed, and leveling rolling are carried out using a continuous annealing production line (CAL).

[0164] In order to refine the original γ particle size, increasing the heating rate is effective. To reduce the original γ particle size to below 10 μm, the average heating rate above 700℃ should be above 0.4℃ / second.

[0165] It should be noted that the average heating rate here refers to "annealing temperature (°C) - 700 (°C) / heating time (seconds) from 700°C to the annealing temperature".

[0166] To minimize the amount of cementite particles and other carbides remaining after annealing due to incomplete solution treatment, annealing is performed at high temperatures for an extended period. Specifically, the annealing temperature needs to be set to 800°C or higher. Annealing at temperatures exceeding 950°C results in excessively large original gamma particle sizes; therefore, the annealing temperature is set to 950°C or lower. More preferably, it is set to 900°C or lower. Furthermore, prolonged soaking time (holding time) also leads to excessively large original gamma particle sizes; therefore, soaking time is set to 600 seconds or lower.

[0167] To reduce ferrite and residual gamma, and to achieve a martensite area ratio of over 95%, a first average cooling rate of 2°C / second or higher is required to cool from the annealing temperature to 420°C. When the first average cooling rate is less than 2°C / second, a large amount of ferrite is formed, carbon is enriched in gamma, martensite hardens, and resistance to delayed fracture deteriorates. Therefore, the first average cooling rate is set to be 2°C / second or higher.

[0168] There is no particular limit to the upper limit of the first average cooling rate, but the first average cooling rate is preferably set to 100°C / second or less.

[0169] To suppress the formation of bainitic ferrite and lower bainite, and to achieve a martensite area ratio of 95% or more, it is necessary to cool from 420°C to a cooling stop temperature below 280°C at a second average cooling rate of 10°C / second or higher. In microstructures with abundant bainite formation, strength decreases and residual γ increases, thus deteriorating resistance to delayed fracture. Therefore, the second average cooling rate from 420°C to the cooling stop temperature below 280°C is set to 10°C / second or higher. Preferably, the second average cooling rate is 20°C / second or higher, more preferably 70°C / second or higher.

[0170] It should be noted that the average cooling rate here refers to "(cooling start temperature (°C) - cooling stop temperature (°C)) / cooling time (seconds) from the start of cooling to the end of cooling".

[0171] More specifically, the first average cooling rate is defined as "(annealing temperature (°C) - 420°C)) / cooling time (seconds) from the annealing temperature to 420°C".

[0172] In addition, the second average cooling rate is defined as "(420 (°C) - cooling stop temperature (°C)) / cooling time (seconds) from 420°C to cooling stop temperature".

[0173] Carbides distributed within the martensite form during low-temperature holding after quenching, becoming hydrogen trapping points in the steel plate and thus contributing to excellent resistance to delayed fracture. Therefore, to ensure a tensile strength of ≥1310 MPa (TS≥1310 MPa) and the formation of the aforementioned carbides, it is necessary to appropriately control the holding temperature and holding time. For this purpose, a holding temperature of 120–260°C for 20–1500 seconds is required. If the holding temperature is below the lower limit of 120°C or the holding time is short, the carbide distribution density within the phase transformation phase becomes insufficient, and the resistance to delayed fracture deteriorates. Furthermore, at a high temperature above the upper limit of 260°C, the coarsening of carbides within grains and at bulk grain boundaries may become significant, further deteriorating the resistance to delayed fracture. Here, a holding time of 60 seconds or more is preferred.

[0174] It should be noted that holding at 120℃~260℃ can be achieved by rapidly cooling to near room temperature and then reheating to 120~260℃ and holding for 20~1500 seconds, or by setting the cooling stop temperature to 120~260℃ and controlling the holding time to 20~1500 seconds. Alternatively, the above thermal process can be achieved by reducing the cooling rate within a temperature range below 260℃, or by intermittent annealing after cooling to room temperature.

[0175] From the viewpoint of stabilizing the compressive formability by adjusting the surface roughness and flattening the sheet shape, the steel sheet obtained in this way can be subjected to surface finishing rolling. In this case, the surface finishing elongation is preferably set to 0.1 to 0.6%. In this case, the surface finishing roll is a rough-surface roll, and from the viewpoint of shape flattening, the roughness Ra of the steel sheet is preferably adjusted to 0.8 to 1.8 μm.

[0176] Alternatively, the obtained steel sheet can be plated. By performing a plating process, a steel sheet with a coating on its surface can be obtained. There is no particular limitation on the type of plating process; it can be either hot-dip galvanizing or electroplating. Alternatively, an alloying plating process can be performed after hot-dip galvanizing. It should be noted that when performing a plating process, the aforementioned surface finishing rolling is performed after the plating process.

[0177] According to the present invention, the delayed fracture resistance of high-strength cold-rolled steel sheets is significantly improved, which contributes to the increased strength and weight reduction of components resulting from the application of high-strength steel sheets. Preferably, the steel sheet of the present invention has a thickness of 0.5 mm or more. Furthermore, the steel sheet of the present invention preferably has a thickness of 2.0 mm or less.

[0178] Next, the components of the present invention and their manufacturing method will be described.

[0179] The component of the present invention is formed by performing at least one of forming and welding on the steel plate of the present invention. Furthermore, the manufacturing method of the component of the present invention includes the steps of performing at least one of forming and welding on the steel plate manufactured by the steel plate manufacturing method of the present invention.

[0180] The steel sheet of the present invention has a tensile strength of 1310 MPa or higher and exhibits excellent resistance to delayed fracture. Therefore, components obtained using the steel sheet of the present invention are also high-strength, exhibiting superior resistance to delayed fracture compared to conventional high-strength components. Furthermore, using components of the present invention enables weight reduction. Therefore, components of the present invention can be suitable for applications such as vehicle body frame components.

[0181] Forming can be performed without restriction using general processing methods such as pressure processing. Furthermore, welding can be performed without restriction using general welding methods such as spot welding and arc welding.

[0182] Example

[0183] [Example 1]

[0184] The embodiments of the present invention will be described below.

[0185] The steel with the composition shown in Table 1 was melted and then cast into steel billets.

[0186] The steel billet was subjected to the heat treatment and rolling shown in Table 2 to obtain a steel plate with a thickness of 1.4 mm.

[0187] Specifically, a steel billet with each component composition is heated at an average heating rate of 6°C / min to the heating holding temperature shown in Table 2 (using a steel billet surface thermometer), and held for the heating holding time shown in Table 2. Then, hot finishing rolling is performed at the finishing rolling temperatures shown in Table 2, followed by cooling at an average cooling rate of 50°C / second from the finishing rolling temperature to 650°C.

[0188] Then, the hot-rolled steel sheet is produced by cooling and coiling at the coiling temperature shown in Table 2, and cold-rolled steel sheet is produced by cold rolling the hot-rolled steel sheet at the reduction rate (cold rolling reduction rate) shown in Table 2.

[0189] Then, the cold-rolled steel sheet is heated from 700°C to the annealing temperature shown in Table 2 at the average heating rate shown in Table 2, and then homogenized at the annealing temperature for the homogenization time shown in Table 2.

[0190] Then, the annealing temperature is cooled to 420°C at the first average cooling rate shown in Table 2, and then cooled to the cooling stop temperature shown in Table 2 at the second average cooling rate shown in Table 2. The temperature is then reheated as needed, and then the temperature is held at the holding time shown in Table 2 for the holding time shown in Table 2, thereby performing continuous annealing.

[0191] In Table 2, cold-rolled steel sheet No. 6 (CR) is subjected to electro-galvanizing treatment to produce electro-galvanized steel sheet (EG). In Table 2, the surface temperature of the steel billet is measured by a radiation thermometer, and the center temperature of the steel billet is calculated by heat transfer calculation.

[0192] [Table 1]

[0193]

[0194] The balance other than those mentioned above is Fe and unavoidable impurities.

[0195] (*1) Ar3 point = 910 - 310 × [c] - 80 × [Mn] - 20 × [Cu] - 15 × [Cr] - 55 × [Ni] - 80 × [Mo] (where [M] is the content of element M (mass%))

[0196] [Table 2]

[0197]

[0198] The obtained steel plate is then quantified using the above method, followed by tensile testing and evaluation of its resistance to delayed fracture.

[0199] Specifically, the tissue assay is performed as follows.

[0200] The area ratios of martensite, bainite, and ferrite were determined as follows: The L-section of the steel plate (parallel to the rolling direction and perpendicular to the plate surface) was ground and etched with nitric acid-ethanol solution. Four fields of view were observed using a SEM at 2000x magnification at a distance of 1 / 4 thickness from the plate surface. Image analysis was performed on the photographs to determine the area ratios of martensite, bainite, and ferrite. Here, martensite and bainite refer to structures that appear gray or white in the SEM. Bainite is characterized by the following: an aspect ratio of 2.5 or higher, a plate-like morphology, and a slightly darker appearance compared to martensite. The width (minor axis) of the plate is 0.3–1.7 μm. The distribution density of carbides with diameters of 10–200 nm within the bainite is 0–3 per μm. 2 On the other hand, ferrite is a region with black contrast in SEM. It should be noted that there are trace amounts of carbides, nitrides, sulfides, and oxides inside martensite and bainite, but it is difficult to exclude them. Therefore, the area ratio of the region containing them is used as the area ratio. The determination of retained austenite (retained γ) is as follows: the surface 200 μm of the steel plate is chemically ground with oxalic acid. The retained austenite (retained γ) is determined by X-ray diffraction intensity method using the plate surface as the object. It is calculated based on the integrated intensity of the diffraction peaks of (200)α, (211)α, (220)α, (200)γ, (220)γ, and (311)γ measured by Mo-Kα rays.

[0201] The average grain size of the original austenite grains (original γ grain size) was determined as follows: After grinding the L-section of the steel plate (a section parallel to the rolling direction and perpendicular to the steel plate surface), it was etched using a reagent that corrodes the original γ grain boundaries (e.g., a saturated picric acid aqueous solution or a solution to which ferric chloride has been added). Four fields of view were observed at 500x magnification from a position 1 / 4 thickness away from the steel plate surface. In the obtained photographs, 15 lines were drawn at intervals of at least 10 μm in both the thickness direction and the rolling direction, and the number of intersections between the grain boundaries and the lines was counted. The original γ grain size was calculated by multiplying the value obtained by dividing the line length by the number of intersections by 1.13.

[0202] The ratio of Nb and Ti in the Nb and Ti contained in the carbonitrides that form equivalent circular diameters of 100 nm or more can be determined by the following method.

[0203] After electrolyzing a specified amount of the sample in an electrolyte, the sample sheet was removed from the electrolyte and immersed in a dispersible solution. Next, the precipitates contained in the solution were filtered using a filter with a pore size of 100 nm. The precipitates captured by this 100 nm filter were carbonitrides with a diameter (equivalent circle diameter) of 100 nm or more. The Nb and Ti contents of the residue on the filter and the filtrate after filtration were analyzed to determine the Nb and Ti contents in carbonitrides with a diameter of 100 nm or more and in carbonitrides with a diameter of less than 100 nm. Inductively coupled plasma (ICP) emission spectrometry could be used in the analysis. Then, the ratio of the total Nb and Ti content in carbonitrides with a diameter of 100 nm or more to the total Nb and Ti content in the steel was calculated.

[0204] Nb carbonitrides and Ti carbonitrides per 1 mm 2 The number (distribution density) can be calculated as follows: After grinding the L-section (the perpendicular section parallel to the rolling direction) of the steel plate, without corrosion, continuously take 2mm SEM images in the region from 1 / 5 to 4 / 5 of the plate thickness, that is, from 1 / 5 of the plate thickness from the surface, sandwiched between the center of the plate thickness and the 4 / 5 position. 2 The number of such carbonitrides in the above area was determined by SEM images, which allowed for the calculation of the number of Nb and Ti carbonitrides per 1 mm. 2 The number (distribution density) of inclusions. Here, the SEM image is preferably a reflectance electron image. In addition, the magnification is set to 2000x. However, if it is difficult to accurately determine the size of the precipitates at 2000x, the inclusion particles can be magnified to 10000x to draw the carbonitrides described above.

[0205] In addition, in the tensile test, JIS No. 5 tensile test pieces are cut at 1 / 4 of the width of the roll material with the rolling right angle as the length direction, and tensile tests are carried out (according to JIS Z2241) to evaluate YP, TS, and El.

[0206] The evaluation of resistance to delayed fracture was conducted as follows. Conventional Strain Rate Test (CSRT) specimens were cut from the steel sheet (coil) at 1 / 4 of its width along the rolling right angle. The CSRT specimens were tensile test pieces with semi-circular notches of 3 mm radius at both ends of the parallel section, which was 12.5 mm wide and 25 mm long. A 10% (w / w) ammonium thiocyanate aqueous solution was mixed with McIlvaine buffer solution at pH 3 at a 1:1 volume ratio. The CSRT specimens were then subjected to... 2The sample was immersed in a solution (pH 3) at 20°C for 24 hours, with the liquid volume adjusted to reach 20 ml to allow hydrogen to penetrate and diffuse into the test piece. Immediately after 24 hours, a tensile test was performed at a crosshead speed of 1 mm / min to determine the fracture stress. The fracture stress before immersion was defined as σ0, and the fracture stress after hydrogen penetration and diffusion through immersion was defined as σ1. The resistance to delayed fracture was evaluated using σ1 / σ0. Samples with σ1 / σ0 of 0.80 or higher at a TS of 1310 MPa or higher but less than 1500 MPa, σ1 / σ0 of 0.50 or higher at a TS of 1500 MPa or higher but less than 1800 MPa, and σ1 / σ0 of 0.35 or higher at a TS of 1800 MPa or higher were judged to have excellent resistance to delayed fracture.

[0207] The microstructure and properties of the obtained steel plates are shown in Table 3.

[0208] [Table 3]

[0209]

[0210] (*1) Residual area ratio: The total area ratio of bainite, ferrite and retained austenite

[0211] (*2) Number density of Nb and Ti carbonitrides: The combined number density of Nb and Ti carbonitrides with an equivalent diameter of 1.0 μm or larger.

[0212] The steel plates within the scope of this invention are high-strength and have excellent resistance to delayed fracture.

[0213] On the other hand, regarding No.16 (steel grade P), the C content is excessive, and insufficient resistance to delayed fracture is not achieved.

[0214] Regarding No.17 (steel grade Q), the C content is insufficient, and the TS (steel ore) has not been fully obtained.

[0215] Regarding No.18 (steel grade R), the Si content is excessive, and insufficient resistance to delayed fracture is not achieved.

[0216] Regarding No.19 (steel grade S), the Mn content is insufficient, resulting in the formation of ferrite and insufficient formation of martensite, thus failing to achieve adequate resistance to delayed fracture.

[0217] Regarding No.20 (steel grade T), the P content is excessive, and insufficient resistance to delayed fracture is not achieved.

[0218] Regarding No.21 (steel grade U), the S content is excessive, and insufficient resistance to delayed fracture is not achieved.

[0219] Regarding No.22 (steel grade V), the sol.Al content is excessive, and insufficient resistance to delayed fracture is not achieved.

[0220] Regarding No.23 (steel grade W), the nitrogen content is excessive, and insufficient resistance to delayed fracture is not achieved.

[0221] Regarding No.24 (steel grade X), the combined content of Nb and Ti is excessive, and there is an excessive amount of precipitates larger than 1.0 μm, resulting in insufficient resistance to delayed fracture.

[0222] Regarding No.25 (steel grade Y), the combined content of Nb and Ti is insufficient, and the delayed fracture resistance characteristics are not fully obtained.

[0223] Regarding No.26 (steel grade Z), the B content is excessive, and insufficient resistance to delayed fracture is not achieved.

[0224] Regarding No.27 (steel grade C), the heating temperature (slab surface temperature (SRT)) is low, and Nb and Ti are not fully dissolved. Therefore, a large number of Nb and Ti precipitates larger than 1.0 μm exist, and the delayed fracture resistance characteristics are not fully obtained.

[0225] Regarding No.28 (steel grade C), the billet heating time was too short, and Nb and Ti were not fully dissolved. As a result, a large number of Nb and Ti precipitates larger than 1.0 μm were present, and the delayed fracture resistance characteristics were not fully obtained.

[0226] Regarding No.29 (steel grade C), the coiling temperature (CT) is high, and a large amount of Nb and Ti precipitates above 1.0 μm are present, resulting in insufficient resistance to delayed fracture.

[0227] Regarding No.30 (steel grade C), the average heating rate during annealing is low, and sufficient resistance to delayed fracture is not achieved.

[0228] Regarding No.31 (steel grade C), the first average cooling rate is low, resulting in insufficient formation of ferrite and martensite, thus failing to achieve adequate delayed fracture characteristics.

[0229] Regarding No.32 (steel grade C), the second average cooling rate is low, resulting in insufficient formation of bainite and martensite, thus failing to achieve adequate delayed fracture characteristics.

[0230] Regarding No. 33 (steel grade C), the temperature was kept high, and sufficient TS was not obtained.

[0231] Regarding No.35 (steel grade C), the annealing temperature was high, and the original γ grain size was not sufficiently refined, thus failing to obtain sufficient resistance to delayed fracture.

[0232] [Example 2]

[0233] A galvanized steel sheet obtained by galvanizing under manufacturing conditions No. 6 (suitable example) of Table 2 in Example 1 is press-formed to manufacture a component of the present invention. Furthermore, a galvanized steel sheet obtained by galvanizing under manufacturing conditions No. 6 (suitable example) of Table 2 in Example 1 is joined with a galvanized steel sheet obtained by galvanizing under manufacturing conditions No. 7 (suitable example) of Table 2 in Example 1 by spot welding to manufacture a component of the present invention.

[0234] Regarding the components of these examples of the present invention, the tensile strength TS is 1800 MPa or more, the σ1 / σ0 is 0.40 or more and 0.35 or more, and the delayed fracture resistance is excellent. Therefore, it can be seen that these components are suitable for use in automotive parts, etc.

[0235] Similarly, the steel sheet according to manufacturing condition No. 6 (suitable example) in Table 2 of Example 1 was press-formed to manufacture the component of the present invention. Furthermore, the steel sheet according to manufacturing condition No. 6 (suitable example) in Table 2 of Example 1 was joined with the steel sheet according to manufacturing condition No. 7 (suitable example) in Table 2 of Example 1 by spot welding to manufacture the component of the present invention. Regarding these components of the present invention, the tensile strength TS is 1800 MPa or more, the σ1 / σ0 value is 0.40 or more and 0.35 or more, and the delayed fracture resistance is excellent; therefore, these components are suitable for use in automotive parts, etc.

Claims

1. A type of steel plate, It contains, by mass%, C: 0.12% to 0.40%, Si: 1.5% to 1.5%, Mn: 1.8% to 4.0%, P: 0.03% to 0.03%, S: less than 0.0023%, sol.Al: 0.20% to 0.005%, N: 0.005%, B: 0.0100%, and one or more of Nb and Ti: totaling 0.005% to 0.080%, with the balance being Fe and unavoidable impurities. Furthermore, it has a microstructure in which the area ratio of martensite relative to the overall microstructure is more than 95% and less than 100%, and the balance is composed of one or more of bainite, ferrite, and retained austenite. The average grain size of the original austenite grains is less than 18 μm. More than 90% by mass of the total Nb and Ti content exists in the form of carbonitrides with an equivalent circle diameter of more than 100 nm. Nb carbonitrides and Ti carbonitrides with an equivalent circle diameter of 1.0 μm or more, totaling 800 per mm. 2 The following exists, The tensile strength of the steel plate is above 1310 MPa.

2. The steel plate according to claim 1, wherein, As a component, the ingredient also contains, by mass%, at least one group of elements selected from groups A to D. Group A: Selected from one or both of Cu: less than 1.0% and Ni: less than 1.0%; Group B: Selected from one or more of the following: Cr: less than 1.0%, Mo: less than 0.3%, V: less than 0.5%, Zr: less than 0.2%, and W: less than 0.2%; Group C: Selected from one or more of Ca: less than 0.0030%, Ce: less than 0.0030%, La: less than 0.0030%, REM: less than 0.0030%, and Mg: less than 0.0030%, wherein REM does not include Ce and La; Group D: Selected from one or both of Sb: less than 0.1% and Sn: less than 0.1%.

3. The steel plate according to claim 1, wherein, The average grain size of the original austenite grains is less than 10 μm.

4. The steel plate according to claim 2, wherein, The average grain size of the original austenite grains is less than 10 μm.

5. The steel plate according to any one of claims 1 to 4, wherein the fracture stress σ0 before immersion in a solution containing 10% by mass of an aqueous solution of ammonium thiocyanate and a McIlvaine buffer solution at pH 3, the fracture stress σ1 after immersion in said solution, and the tensile strength satisfy the following (A), (B), or (C): (A) Tensile strength is 1310 MPa or more and less than 1500 MPa, and σ1 / σ0 is 0.80 or more; (B) Tensile strength is above 1500MPa and below 1800MPa, and σ1 / σ0 is above 0.50; (C) Tensile strength is above 1800MPa, and σ1 / σ0 is above 0.

35.

6. The steel plate according to any one of claims 1 to 4, wherein, As a component, it contains less than 0.0010% S by mass.

7. The steel plate according to claim 5, wherein, As a component, it contains less than 0.0010% S by mass.

8. The steel plate according to any one of claims 1 to 4, wherein, The steel plate has a coating.

9. The steel plate according to claim 5, wherein, The steel plate has a coating.

10. The steel plate according to claim 6, wherein, The steel plate has a coating.

11. The steel plate according to claim 7, wherein, The steel plate has a coating.

12. A component formed by forming and welding at least one of the steel plates according to any one of claims 1 to 11.

13. A method for manufacturing a steel plate, wherein, A steel billet having the composition described in any one of claims 1 to 7 is heated from 1000°C to a holding temperature of 1250°C or higher using a steel billet surface thermometer at an average heating rate of 10°C / min or less, and held at said holding temperature for 30 minutes or more. Perform hot finishing rolling with the finishing temperature set above Ar3. The cooling process is performed with an average cooling rate set to 40°C / second or higher within the range from the finishing rolling temperature to 650°C. Then, it is cooled and wound at a winding temperature below 600°C to produce hot-rolled steel sheet. The hot-rolled steel sheet is cold-rolled to produce a cold-rolled steel sheet by cold rolling with a reduction rate of more than 40%. The following continuous annealing process is performed: the annealing temperature is set to 800–950°C, the cold-rolled steel sheet is heated from 700°C to the annealing temperature at an average heating rate of 0.4°C / second or higher and less than 2.0°C / second, held at the annealing temperature for less than 600 seconds, cooled from the annealing temperature to 420°C at a first average cooling rate of 2°C / second or higher, cooled from 420°C to a cooling stop temperature of less than 280°C at a second average cooling rate of 10°C / second or higher, and then held at a holding temperature of 120–260°C for 20–1500 seconds.

14. The method for manufacturing a steel plate according to claim 13, wherein, The steel plate surface is plated after the continuous annealing.

15. A method for manufacturing a component, comprising at least one of forming and welding a steel plate manufactured by the method for manufacturing a steel plate according to claim 13 or 14.

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