Steel sheet, member, and method for manufacturing the same

Abstract

The present invention provides a steel sheet having a tensile strength of 780 MPa or more, having high ductility and excellent stretch flange formability, and having excellent energy absorption characteristics at the time of collision, a member, and a method for manufacturing the same. A steel sheet having a composition in which C, Si, Mn, P, S, sol. Al, N are specified, and having a steel structure in which the total of ferrite and bainite ferrite is 5 to 60%, tempered martensite is 20 to 80%, fresh martensite is 20% or less (including 0%), residual austenite is 5 to 25%, a structure composed of one or two or more of ferrite, tempered martensite, fresh martensite, bainite, and residual austenite is 90% or more (including 100%), and in the ferrite, high-Mn ferrite is 20% or more and 70% or less in area ratio, S C≥0.5 / S C≥0.3 x 100 is 20% or more.

Description

Technical Field

[0001] This invention relates to steel plates, components, and methods for manufacturing them used in various applications such as automobiles and home appliances. Background Technology

[0002] In recent years, due to the increasing demand for lightweight automobile bodies, high-strength steel sheets with tensile strengths exceeding 780 MPa are being used in automotive frame components and seat components. However, when using high-strength steel sheets with tensile strengths exceeding 780 MPa in automotive parts, stamping cracks are prone to occur due to reduced ductility and reduced formability of the extended flange. Therefore, it is desirable for these high-strength steel sheets to have better formability than before.

[0003] Furthermore, from the perspective of occupant safety, it is necessary to suppress deformation around the cab during a collision. Therefore, energy-absorbing components such as longitudinal beams are required to absorb collision energy by deforming individual components during a collision. However, in high-strength steel plates with tensile strength of 780 MPa or higher, due to the reduction in axial crushing characteristics, the parts subjected to single-processing based on forming become the starting point for component fracture during a collision, posing a problem of not being able to stably exert the collision energy absorption capacity.

[0004] Against this backdrop, TRIP steel, which disperses residual γ in the microstructure of steel plates, was developed as a technology to improve the ductility of steel plates. For example, Patent Document 1 discloses, by mass%, the following components are present: C: 0.15% or more and 0.30% or less; P: 0.040% or less; S: 0.0100% or less; N: 0.0100% or less; O: 0.0060% or less; one or two of Si and Al: totaling 0.70% or more and 2.50% or less; one or two of Mn and Cr: totaling 1.50% or more and 3.50% or less; Mo: 0% or more and 1.00% or less; Ni: 0% or more and 1.00% or less; Cu: 0% or more and 1.00% or less; Nb: 0% or more and 0.30% or less; Ti: 0% or more and 0.30% or less; V: 0% or more and 0.30% or less; B: 0% or more and 0.0050% or less; Ca: 0% or more and 0.0050% or less. The steel sheet has the following composition: 0.0400% or less, Mg: 0% or more and 0.0400% or less, and REM: 0% or more and 0.0400% or less, with the balance consisting of Fe and impurities. As a percentage of the overall microstructure, it has one or both of ferrite and granular bainite, totaling 10% or more and 50% or less; one or both of upper bainite and lower bainite, totaling 10% or more and 50% or less; tempered martensite: greater than 0% and 30% or less; retained austenite: 5% or more; and one or more of pearlite, cementite, and martensite, totaling 0% or more and 10% or less. The area percentage of the ferrite is less than 25% of the total area percentage of the ferrite and granular bainite. This results in a steel sheet with a tensile strength of 980 MPa or more and excellent elongation and porosity.

[0005] Patent document 2 discloses that, as a steel composition, it contains, by mass percent, C: 0.07–0.20%, Si: 0.1–2.0%, Mn: 2.0–3.5%, P: 0.05% or less, S: 0.05% or less, Sol.Al: 0.005–0.1%, with the balance consisting of Fe and unavoidable impurities. The steel microstructure, by area percentage, is ferrite: 60% or less, tempered martensite: 40% or more, and fresh martensite: 10% or less. Furthermore, the void number density of the bent portion in the VDA bending test is 1500 voids / mm. 2 The following results in a high-strength hot-dip galvanized steel sheet with a tensile strength of over 980 MPa and excellent fracture resistance upon impact.

[0006] Patent document 3 discloses a high-strength cold-rolled steel sheet with excellent strength, ductility, and porosity, containing, by mass percent, 0.10-0.40% C, 0.5-4.0% Mn, 0.005-2.5% Si, 0.005-2.5% Al, and 1.0% or less Cr, with the balance consisting of iron and unavoidable impurities, limited to 0.05% or less P, 0.02% or less S, and 0.006% or less N. As a steel structure, it contains 2-30% retained austenite by area ratio, limits martensite to 20% or less, and the average particle size of cementite is 0.01 μm or more and 1 μm or less. The cementite contains 30% or more and 100% or less cementite with an aspect ratio of 1 or more and 3 or less.

[0007] Existing technical documents

[0008] Patent documents

[0009] Patent Document 1: Japanese Patent No. 6338038

[0010] Patent Document 2: Japanese Patent No. 6795122

[0011] Patent Document 3: Japanese Patent No. 4903915 Summary of the Invention

[0012] The problem that the invention aims to solve

[0013] In the technology described in Patent Document 1, by effectively utilizing granular bainite, the hardness difference between different phases in the composite steel sheet is greatly reduced, and the decrease in elongated flange formability caused by the increase of ferrite is suppressed. As a result, it is possible to manufacture steel sheets with excellent ductility and elongated flange formability. However, since the area ratio of tempered martensite is less than 30%, there is a tendency for the axial crushing characteristics to deteriorate.

[0014] In the technology described in Patent Document 2, steel plates with excellent axial crushing characteristics can be manufactured by reducing the void density. However, since the bainite phase, which has reached the intermediate hardness between ferrite and tempered martensite, is not present, the hardness difference between the different phases is large, and the formability of the elongated flange becomes a problem.

[0015] Patent document 3 describes a steel sheet with excellent ductility and elongation flange formability obtained by controlling the shape of cementite, but it does not consider axial crushing characteristics. During bending deformation, cracks are generated starting from cementite, and the axial crushing characteristics deteriorate.

[0016] Thus, in the prior art, at least one of the following is poor: ductility, elongation flange formability, and energy absorption characteristics during impact (axial crushing characteristics).

[0017] The present invention was made to solve such problems, and its object is to provide steel plates, components, and methods thereof having tensile strength of 780 MPa or more, high ductility and excellent elongation flange formability, and excellent energy absorption characteristics upon impact.

[0018] In this invention, tensile strength is determined by a tensile test according to JIS Z 2241 (2011).

[0019] Furthermore, in this invention, high ductility refers to the following for the total elongation (T-El) measured by tensile testing based on JIS Z 2241 (2011): (A) when TS is 780 MPa or more and less than 980 MPa, T-El is 18.0% or more; (B) when TS is 980 MPa or more and less than 1180 MPa, T-El is 16.0% or more; (C) when TS is 1180 MPa or more and less than 1320 MPa, T-El is 14.0% or more; and (D) when TS is 1320 MPa or more, T-El is 13.0% or more.

[0020] Furthermore, in this invention, excellent extension flange forming performance means that after punching a 100mm × 100mm square sample using a punching tool with a punch diameter of 10mm and a die diameter of 10.3mm (13% clearance), a conical punch with a 60-degree apex angle is used to expand the hole so that the burrs generated during the forming of the punched hole are on the outside until a crack penetrating the plate thickness is generated. d0 is the initial hole diameter (mm), d is the hole diameter (mm) when the crack is generated, and the hole expansion rate λ (%) = {(d-d0) / d0} × 100 is 30% or more.

[0021] Furthermore, in this invention, excellent energy absorption characteristics during collision means that the area up to the maximum stress in the nominal stress-nominal strain curve of the tensile test according to JIS Z 2241 (2011) is regarded as the energy absorption during deformation, and the energy absorption is 13000 MPa·% or more.

[0022] Methods for solving problems

[0023] The inventors have conducted in-depth research on methods for achieving high ductility, excellent elongated flange formability, and excellent energy absorption characteristics during impact. The results showed that by forming a microstructure in which the total area ratio of ferrite, tempered martensite, fresh martensite, bainite, and retained austenite is 90% or more, high ductility and excellent elongated flange formability can be ensured. Furthermore, by appropriately controlling the phase fraction of each microstructure and broadly adjusting the amount of dissolved Mn in the ferrite within an appropriate range, excellent energy absorption characteristics during impact can also be obtained. In addition, it was clarified that these microstructures can be achieved by annealing at a specified temperature followed by a cooling process within a temperature range up to -15°C at an average cooling rate CR1 of 0.01°C / s or more and 5°C / s or less, allowing the ferrite phase transformation to occur without Mn diffusion, thereby including both ferrite generated during annealing and ferrite generated during cooling.

[0024] This invention is based on the above insights and specifically provides the following invention.

[0025] [1] A steel plate having, by mass%, a composition comprising: C: 0.06% or more and 0.25% or less, Si: 0.4% or more and 2.5% or less, Mn: 1.5% or more and 3.5% or less, P: 0.10% or less, S: 0.010% or less, sol.Al: 1.0% or less, N: 0.015% or less, with the balance being Fe and unavoidable impurities. It also possesses the following steel microstructure: by area percentage, the total amount of ferrite and bainitic ferrite is 5% to 60%; tempered martensite is 20% to 80%; fresh martensite is 20% to 0% (inclusive); retained austenite by volume percentage is 5% to 25%; and the total area percentage of ferrite, tempered martensite, fresh martensite, bainite, and retained austenite is 90% to 100% (inclusive). When the ferrite content is greater than 0%, the proportion of ferrite with a solid solution Mn content of 2.0% or more in the total ferrite is 20% or more and less than 70% by area ratio. The area S of the region where the C concentration is 0.5% by mass or higher C≥0.5 The area S relative to the region where the C concentration is 0.3% by mass or higher C≥0.3 Ratio: (S) C≥0.5 / S C≥0.3 )×100 is more than 20%.

[0026] [2] The steel plate according to [1] above, wherein, as the composition, it further contains, by mass%, one or more of the following: Ti: less than 0.1%, B: less than 0.01%, Cu: less than 1%, Ni: less than 1%, Cr: less than 1.0%, Mo: less than 0.5%, V: less than 0.5%, Nb: less than 0.1%, Zr: less than 0.2%, W: less than 0.2%, Ca: less than 0.0040%, Ce: less than 0.0040%, La: less than 0.0040%, Mg: less than 0.0030%, Sb: less than 0.1%, Sn: less than 0.1%.

[0027] [3] According to the steel plate described in [1] or [2] above, wherein the steel structure further contains, in terms of area ratio, 3% or more and 40% or less of internal carbides per 10 μm 2 It contains less than 20 bainitic ferrites.

[0028] [4] The steel plate according to any one of [1] to [3] above, wherein a zinc coating is provided on the surface.

[0029] [5] A component made of steel plate as described in any one of [1] to [4] above.

[0030] [6] A method for manufacturing a steel plate, wherein, After hot rolling and cold rolling of a steel billet having the composition described in [1] or [2] above, the resulting cold-rolled steel sheet is annealed. The annealing process includes, in sequence: The process of holding the product at an annealing temperature above 775°C and below 830°C; A cooling process performed within a temperature range from the annealing temperature to -15°C with an average cooling rate CR1 of 0.01°C / s or more and 5°C / s or less. A cooling process performed at an average cooling rate of CR2: 3°C / s or higher within a temperature range from annealing temperature to cooling stop temperature above 200°C and below 300°C. A heating process performed at an average heating rate of 2°C / s or more within a temperature range from the cooling stop temperature to 380°C; Processes involving a temperature range of 340°C to 590°C with an average cooling rate CR4 of 0.01–5°C / s and a dwell time of 20 seconds to 3000 seconds; and A process that cools to a temperature below 50°C at an average cooling rate CR5 of 0.1°C / s or higher.

[0031] [7] A method for manufacturing a steel plate, wherein, After hot rolling and cold rolling of a steel billet having the composition described in [1] or [2] above, the resulting cold-rolled steel sheet is annealed. The annealing process includes, in sequence: The process of holding the product at an annealing temperature above 775°C and below 830°C; A cooling process performed within a temperature range from the annealing temperature to -15°C with an average cooling rate CR1 of 0.01°C / s or more and 5°C / s or less. A cooling process performed at an average cooling rate of CR2A: 3°C / s or higher within a temperature range from -15°C to 500°C. A process with an average cooling rate of CR3 of less than 10℃ and a dwell time of more than 10 seconds and less than 60 seconds within a temperature range from 500℃ to the dwell temperature above the martensitic transformation start temperature Ms and above the dwell temperature above 320℃; A cooling process performed at an average cooling rate of CR2B: 3°C / s or more within a temperature range from the stated stopping temperature to a cooling stopping temperature of 200°C or more and 300°C or less; A heating process performed at an average heating rate of 2°C / s or more within a temperature range from the cooling stop temperature to 380°C; Processes involving a temperature range of 340°C to 590°C with an average cooling rate CR4 of 0.01–5°C / s and a dwell time of 20 seconds to 3000 seconds; and A process that cools to a temperature below 50°C at an average cooling rate CR5 of 0.1°C / s or higher.

[0032] [8] In the steel plate manufacturing method described in [6] or [7] above, the steel plate is subjected to hot-dip galvanizing or alloyed hot-dip galvanizing in the process of holding at the average cooling rate CR4: 0.01 to 5°C / s.

[0033] [9] The method for manufacturing steel plates according to [6] or [7] above includes a step of electro-galvanizing after a step of cooling at an average cooling rate CR5 of 0.1°C / s or higher.

[0034]

[10] A method for manufacturing a component, comprising a step of forming a component by performing at least one of forming or joining processes on a steel plate as described in any one of [1] to [4].

[0035] Invention Effects

[0036] According to the present invention, a steel sheet with high ductility and excellent extended flange formability, as well as excellent energy absorption characteristics during impact, can be obtained. Furthermore, according to the present invention, high strength can also be achieved. If the steel sheet of the present invention is applied to automotive parts, lightweighting of the automotive parts can be achieved, and improved fuel efficiency can be expected. Attached Figure Description

[0037] Figure 1 This is an example of a SEM image of the steel structure in a steel plate.

[0038] Figure 2 This is a diagram illustrating the method for measuring the steel structure of the steel plate of the present invention.

[0039] Figure 3 These are diagrams illustrating a method for manufacturing the steel plate of the present invention: (a) is a diagram illustrating a manufacturing method without dwell time, and (b) is a diagram illustrating a manufacturing method with dwell time. Detailed Implementation

[0040] The present invention will now be described in detail. It should be noted that the present invention is not limited to the following embodiments.

[0041] <steel plate>

[0042] The steel plate of the present invention comprises, by mass percent, C: 0.06% or more and 0.25% or less, Si: 0.4% or more and 2.5% or less, Mn: 1.5% or more and 3.5% or less, P: 0.10% or less, S: 0.010% or less, sol.Al: 1.0% or less, N: 0.015% or less, with the balance being Fe and unavoidable impurities, and has a steel structure in which the total amount of ferrite and bainitic ferrite, based on the area percentage in the entire structure, is 5% or more and 60% or less. The ferrite content is 20% to 80% or more, the fresh martensite content is 20% to 0% or less (including 0%), the retained austenite content is 5% to 25% by volume, and the total content of ferrite, tempered martensite, fresh martensite, bainite, and retained austenite is 90% to 100% or more. In cases where ferrite content is greater than 0%, the proportion of ferrite with a dissolved Mn content of 2.0% or more in the total ferrite is 20% to 70% by area, and the area S of the region with a C concentration of 0.5% or more is... C≥0.5 The area S relative to the region where the C concentration is 0.3% by mass or higher C≥0.3 Ratio: (S) C≥0.5 / S C≥0.3 )×100 is more than 20%.

[0043] The steel plate of the present invention will be described below in the order of composition and steel structure.

[0044] The steel plate of the present invention contains the following components. In the following description, the unit "%" for the content of the components refers to "mass%".

[0045] C: Above 0.06% and below 0.25%

[0046] From the perspective of ensuring the area ratio of tempered martensite to ensure the specified strength, ensuring the area ratio (volume ratio) of retained austenite (retained γ) to improve ductility, and enriching in retained γ to stabilize it and improve ductility, C is required. When the C content is less than 0.06%, these effects cannot be sufficiently ensured; therefore, a lower limit of 0.06% is set. The C content is preferably 0.09% or more, and more preferably 0.11% or more.

[0047] On the other hand, if the C content is greater than 0.25%, it will result in excessive strength, reduced ductility, and an increase in blocky martensite, which may sometimes deteriorate the formability of the extended flange.

[0048] Therefore, the upper limit of C content is set at 0.25%. From the viewpoint of improving ductility, the C content is preferably set at 0.22% or less. From the viewpoint of further improving ductility, the C content is more preferably set at 0.20% or less.

[0049] Si: 0.4% or more and 2.5% or less

[0050] From the perspective of increasing strength by strengthening ferrite, and improving ductility by suppressing carbide formation in martensite and bainite and enhancing the stability of residual γ, Si is included. Based on these considerations, the Si content is set to be 0.4% or higher.

[0051] From the viewpoint of improving ductility, the Si content is preferably 0.6% or more. More preferably, the Si content is 0.8% or more.

[0052] On the other hand, when the Si content is greater than 2.5%, the high strength of the blocky fresh martensite due to the excessive increase in tempering softening resistance reduces the elongation of the flange. Furthermore, the rolling load during hot rolling becomes extremely high, making the manufacture of thin sheets difficult. In addition, chemical conversion treatment properties and weld toughness sometimes deteriorate. Therefore, the Si content is set to 2.5% or less.

[0053] From the viewpoint of ensuring chemical conversion processability, the toughness of raw materials and welds, the Si content is preferably set to less than 2.0%. From the viewpoint of ensuring the toughness of welds, the Si content is preferably set to less than 1.8%, and more preferably to less than 1.5%.

[0054] Mn: 1.5% or more and 3.5% or less

[0055] From the viewpoint of ensuring strength by guaranteeing tempered martensite and / or bainite with a specified area ratio, and from the viewpoint of improving ductility by stabilizing residual γ through a reduction in its Ms point, Mn is an important element. Furthermore, like Si, Mn is an important element from the viewpoint of improving ductility by suppressing carbide formation in bainite, and from the viewpoint of improving ductility by increasing the volume fraction of residual γ. To obtain these effects, the Mn content is set to 1.5% or more. From the viewpoint of improving ductility by stabilizing residual γ, the Mn content is preferably 2.5% or more. More preferably, the Mn content is 2.6% or more, and even more preferably 2.7% or more.

[0056] On the other hand, if the Mn content is greater than 3.5%, the bainitic phase transformation is significantly delayed, thus reducing ductility. In addition, if the Mn content is greater than 3.5%, it is difficult to suppress the formation of blocky coarse γ-rays and blocky coarse martensite, and the formability of the extended flange is also deteriorated.

[0057] Furthermore, when the Mn content is greater than 3.5%, the hardenability is excessively increased. Therefore, in processes where cooling occurs at an average cooling rate CR1 of 0.01℃ / s or more and 5℃ / s or less within the temperature range from the annealing temperature to -15℃, the amount of ferrite with a solid solution Mn content of 2.0% by mass or more becomes insufficient, and sometimes the energy absorption characteristics during collisions are not adequately obtained.

[0058] Therefore, the Mn content is set to 3.5% or less. From the viewpoint of promoting ferrite and bainite phase transformations to ensure high ductility and energy absorption characteristics during collisions, the Mn content is preferably set to 3.2% or less. The Mn content is more preferably 3.1% or less.

[0059] P: below 0.10%

[0060] Phosphorus (P) is an element that strengthens steel, but a high P content deteriorates spot weldability. Therefore, the P content is set to 0.10% or less. Preferably, the P content is 0.02% or less. From the viewpoint of improving spot weldability, the P content is more preferably 0.01% or less. It should be noted that P can be absent, but from the viewpoint of manufacturing cost, the P content is preferably 0.001% or more.

[0061] S: below 0.010%

[0062] S is an element that improves the peeling properties of oxide scale during hot rolling and inhibits nitriding during annealing, but it has a negative impact on spot weldability, bendability, and porosity. To reduce these adverse effects, the S content is set to 0.010% or less. In this invention, since high contents of C, Si, and Mn can easily deteriorate spot weldability, from the viewpoint of improving spot weldability, the S content is preferably set to 0.0020% or less, and more preferably to be less than 0.0010%.

[0063] It should be noted that it may also be free of sulfur, but from a manufacturing cost perspective, the sulfur content is preferably 0.0001% or more. More preferably, the sulfur content is 0.0005% or more.

[0064] sol.Al: 1.0% or less

[0065] Al is included for deoxidation or as a substitute for Si to stabilize residual γ. There is no specific lower limit for sol.Al, but for stable deoxidation, the sol.Al content is preferably 0.005% or more. More preferably, the sol.Al content is 0.01% or more, further preferably 0.02% or more, and even more preferably 0.03% or more.

[0066] On the other hand, when the sol.Al content is greater than 1.0%, the strength of the raw material is drastically reduced, and it also has an adverse effect on the chemical conversion processability. Therefore, the sol.Al content is set to 1.0% or less. To obtain high strength, the sol.Al content is preferably less than 0.50%, more preferably less than 0.20%. The sol.Al content is further preferably less than 0.15%, and even more preferably less than 0.10%.

[0067] N: below 0.015%

[0068] Nitrogen (N) is an element that forms nitrides such as boron (BN), alnitride (AlN), and nitride (TiN) in steel, reducing its thermal ductility and surface quality. Furthermore, in steels containing boron (B), the formation of boron (BN) negates its beneficial effects. When the N content exceeds 0.015%, the surface quality deteriorates significantly. Therefore, the N content is set to 0.015% or less. Preferably, the N content is 0.010% or less.

[0069] It should be noted that it may also be nitrogen-free, but from a manufacturing cost perspective, the nitrogen content is preferably 0.0001% or more. More preferably, the nitrogen content is 0.001% or more.

[0070] The balance other than those mentioned above consists of Fe and unavoidable impurities. The steel plate of the present invention preferably has a composition containing the above-mentioned basic components, with the balance consisting of Fe and unavoidable impurities.

[0071] In addition to the components described above, the steel plate of the present invention may also contain one or more of the following optional elements.

[0072] Ti: less than 0.1%, B: less than 0.01%, Cu: less than 1%, Ni: less than 1%, Cr: less than 1.0%, Mo: less than 0.5%, V: less than 0.5%, Nb: less than 0.1%, Zr: less than 0.2%, W: less than 0.2%, Ca: less than 0.0040%, Ce: less than 0.0040%, La: less than 0.0040%, Mg: less than 0.0030%, Sb: less than 0.1%, Sn: less than 0.1%.

[0073] Ti: below 0.1%

[0074] Ti has the effect of fixing nitrogen in steel as TiN, thereby improving hot ductility, and also improves the hardenability of boron. Furthermore, the precipitation of TiC has the effect of refining the microstructure. To obtain these effects, the Ti content is preferably 0.002% or more. From the viewpoint of sufficiently fixing nitrogen, the Ti content is more preferably 0.008% or more. The Ti content is even more preferably 0.010% or more.

[0075] On the other hand, if the Ti content is greater than 0.1%, it leads to an increase in rolling load and a decrease in ductility due to the increase in precipitation strengthening. Therefore, in the case of Ti content, the Ti content is set to 0.1% or less. The Ti content is preferably 0.05% or less. To ensure high ductility, the Ti content is more preferably set to 0.03% or less.

[0076] B: Below 0.01%

[0077] Boron (B) is an element that improves the hardenability of steel, and has the advantage of easily forming tempered martensite and / or bainite with a specified area ratio. Furthermore, due to the residual B from solid solution, the resistance to delayed fracture is improved. To obtain these effects of B, the B content is preferably set to 0.0002% or more. More preferably, the B content is 0.0005% or more. Even more preferably, the B content is 0.0010% or more.

[0078] On the other hand, when the boron content is greater than 0.01%, not only does its effect become saturated, but it also leads to a significant reduction in thermal ductility, resulting in surface defects. Therefore, when boron is present, the boron content is set to 0.01% or less. The boron content is preferably 0.0050% or less. More preferably, the boron content is 0.0030% or less.

[0079] Cu: less than 1%

[0080] Cu improves the corrosion resistance of automobiles in their operating environment. Furthermore, the corrosion products of Cu coat the surface of the steel sheet, inhibiting hydrogen penetration. Cu is an element incorporated when waste materials are used as raw materials; by allowing Cu inclusion, recycled materials can be used as raw materials, reducing manufacturing costs. From this perspective, a content of 0.005% or more Cu is preferred, and from the viewpoint of improving resistance to delayed fracture, a content of 0.05% or more Cu is more preferred. Even more preferred is a content of 0.10% or more Cu.

[0081] However, excessive Cu content can lead to surface defects. Therefore, when Cu is present, the Cu content is set to 1% or less. Preferably, the Cu content is 0.4% or less, and more preferably 0.2% or less.

[0082] Ni: Below 1%

[0083] Like Cu, Ni is an element that improves corrosion resistance. Furthermore, Ni helps suppress the formation of surface defects that are prone to occur in the presence of Cu. Therefore, it is preferable to contain 0.01% or more Ni. More preferably, the Ni content is 0.04% or more, and even more preferably 0.06% or more.

[0084] However, if the Ni content is too high, the oxide scale formation in the furnace becomes uneven, which in turn causes surface defects. Furthermore, it increases costs. Therefore, when Ni is present, the Ni content is set to 1% or less. Preferably, the Ni content is 0.4% or less, more preferably 0.2% or less.

[0085] Cr: less than 1.0%

[0086] Cr can be included to improve the hardenability of steel and suppress the formation of carbides in martensite and upper / lower bainite. To achieve these effects, the Cr content is preferably 0.01% or more. More preferably, the Cr content is 0.03% or more, and even more preferably 0.06% or more.

[0087] On the other hand, if there is an excessive amount of Cr, the resistance to pitting corrosion deteriorates. Therefore, when Cr is present, the Cr content is set to 1.0% or less. The Cr content is preferably 0.8% or less, more preferably 0.4% or less. The Cr content is even more preferably 0.2% or less, and even more preferably 0.1% or less.

[0088] Mo: 0.5% or less

[0089] Mo can be included to improve the hardenability of steel and suppress the formation of carbides in martensite and upper / lower bainite. To achieve these effects, the Mo content is preferably 0.01% or more. More preferably, it is 0.03% or more, and even more preferably, it is 0.06% or more.

[0090] On the other hand, Mo significantly deteriorates the chemical conversion processability of cold-rolled steel sheets; therefore, in the case of Mo content, the Mo content is set to 0.5% or less. From the viewpoint of improving chemical conversion processability, the Mo content is preferably set to 0.15% or less.

[0091] V: Below 0.5%

[0092] From the perspectives of improving the hardenability of steel, suppressing the formation of carbides in martensite, upper bainite / lower bainite, refining the microstructure, and improving resistance to delayed fracture by precipitating carbides, V may be included. To achieve these effects, the V content is preferably 0.003% or more. More preferably, it is 0.005% or more, and even more preferably, it is 0.010% or more.

[0093] However, if the content of V is high, the castability will be significantly deteriorated. Therefore, when V is present, the V content is set to 0.5% or less. The V content is preferably 0.3% or less, more preferably 0.1% or less, and even more preferably 0.05% or less.

[0094] Nb: below 0.1%

[0095] Nb can be included to achieve the effects of refining the steel microstructure and increasing its strength, promoting bainitic phase transformation through grain refinement, improving bending properties, and enhancing resistance to delayed fracture. To obtain these effects, the Nb content is preferably 0.002% or more. More preferably, the Nb content is 0.004% or more, even more preferably 0.010% or more, and still more preferably 0.020% or more.

[0096] On the other hand, if a large amount of Nb is present, the precipitation strengthening becomes excessive, reducing ductility. Furthermore, it leads to increased rolling load and deterioration of castability. Therefore, when Nb is present, the Nb content is set to 0.1% or less. The Nb content is preferably 0.05% or less, and more preferably 0.03% or less.

[0097] Zr: below 0.2%

[0098] Zr can be included to improve the hardenability of steel, suppress carbide formation in bainite, refine the microstructure, and improve resistance to delayed fracture by precipitating carbides. To achieve these effects, the Zr content is preferably 0.005% or more. More preferably, the Zr content is 0.008% or more, and even more preferably 0.010% or more.

[0099] On the other hand, if the Zr content is high, the amount of coarse precipitates such as ZrN and ZrS remaining after the slab is heated before hot rolling increases, leading to a deterioration in the resistance to delayed fracture. Therefore, when Zr is present, the Zr content is set to 0.2% or less. The Zr content is preferably 0.15% or less, more preferably 0.08% or less, and even more preferably 0.05% or less.

[0100] W: Below 0.2%

[0101] W can be included to improve the hardenability of steel, suppress carbide formation in bainite, refine the microstructure, and improve resistance to delayed fracture by precipitating carbides. To achieve these effects, the W content is preferably 0.005% or more. More preferably, the W content is 0.008% or more, and even more preferably 0.010% or more.

[0102] On the other hand, if the W content is high, the amount of coarse precipitates such as WN and WS that remain undissolved during the heating of the slab before hot rolling increases, leading to a deterioration in the resistance to delayed fracture. Therefore, when W is present, the W content is set to 0.2% or less. The W content is preferably 0.15% or less, more preferably 0.08% or less, and even more preferably 0.05% or less.

[0103] Ca: below 0.0040%

[0104] Ca fixes S in the form of CaS, which helps improve flexural properties and resistance to delayed fracture. Therefore, the Ca content is preferably 0.0002% or more. The Ca content is more preferably 0.0005% or more, and even more preferably 0.0010% or more.

[0105] On the other hand, adding a large amount of Ca will degrade the surface quality and flexibility. Therefore, when Ca is present, the Ca content is set to 0.0040% or less. The Ca content is preferably 0.0035% or less, and more preferably 0.0020% or less.

[0106] Ce: below 0.0040%

[0107] Like Ca, Ce helps to improve bending properties and resistance to delayed fracture by fixing S. Therefore, the Ce content is preferably 0.0002% or more. More preferably, the Ce content is 0.0004% or more, and even more preferably 0.0006% or more.

[0108] On the other hand, if a large amount of Ce is added, the surface quality and flexibility will deteriorate. Therefore, when Ce is present, the Ce content is set to 0.0040% or less. The Ce content is preferably 0.0035% or less, and more preferably 0.0020% or less.

[0109] La: below 0.0040%

[0110] Like Ca, La helps improve flexural properties and resistance to delayed fracture by fixing S. Therefore, the La content is preferably 0.0002% or more. More preferably, the La content is 0.0004% or more, and even more preferably 0.0006% or more.

[0111] On the other hand, if a large amount of La is added, the surface quality and flexibility will deteriorate. Therefore, when La is present, the La content is set to 0.0040% or less. The La content is preferably 0.0035% or less, and more preferably 0.0020% or less.

[0112] Mg: below 0.0030%

[0113] Mg fixes O in the form of MgO, which helps improve resistance to delayed fracture. Therefore, the Mg content is preferably 0.0002% or more. More preferably, the Mg content is 0.0004% or more, and even more preferably 0.0006% or more.

[0114] On the other hand, if a large amount of Mg is added, the surface quality and flexibility will deteriorate. Therefore, when Mg is present, the Mg content is set to 0.0030% or less. The Mg content is preferably 0.0025% or less, and more preferably 0.0010% or less.

[0115] Sb: below 0.1%

[0116] Sb suppresses oxidation and nitriding in the surface layer of the steel plate, thereby suppressing the resulting decrease in the content of carbon (C) and boron (B) in the surface layer. Furthermore, by suppressing the aforementioned decrease in the content of C and B, ferrite formation in the surface layer of the steel plate is suppressed, achieving high strength and improving resistance to delayed fracture. From this perspective, the Sb content is preferably 0.002% or more. More preferably, the Sb content is 0.004% or more, and even more preferably 0.006% or more.

[0117] On the other hand, if the Sb content is greater than 0.1%, the castability deteriorates, and segregation at the original γ grain boundaries worsens the resistance to delayed fracture at the shear end face. Therefore, in the case of Sb, the Sb content is set to 0.1% or less. The Sb content is preferably 0.04% or less, more preferably 0.03% or less, and even more preferably 0.02% or less.

[0118] Sn: less than 0.1%

[0119] Sn inhibits oxidation and nitriding in the surface layer of the steel plate, thereby suppressing the resulting decrease in the content of carbon (C) and boron (B) in the surface layer. Furthermore, by suppressing the aforementioned decrease in the content of C and B, ferrite formation in the surface layer of the steel plate is suppressed, resulting in increased strength and improved resistance to delayed fracture. From this perspective, the Sn content is preferably 0.002% or more. More preferably, the Sn content is 0.004% or more, and even more preferably 0.006% or more.

[0120] On the other hand, if the Sn content is greater than 0.1%, the castability deteriorates. Furthermore, Sn segregates at the original γ grain boundaries, worsening the resistance to delayed fracture at the shear end face. Therefore, when Sn is present, the Sn content is set to 0.1% or less. More preferably, it is 0.04% or less, and even more preferably, it is 0.03% or less.

[0121] When the optional components are contained at a value below the preferred lower limit, the presence of optional elements at a value below the lower limit will not impair the effects of the present invention. Therefore, when the optional elements are contained at a value below the lower limit, the optional elements are contained as unavoidable impurities.

[0122] Next, the steel structure of the steel plate of the present invention will be described.

[0123] The combined area ratio of ferrite and bainitic ferrite is between 5% and 60%.

[0124] Ferrite formed during annealing or cooling, and bainitic ferrite in upper bainite formed during cooling, subsequent heating, or residence time, contributes to improved ductility. Furthermore, enrichment of carbon in the surrounding untransformed austenite contributes to the stabilization of retained austenite. On the other hand, excessive ferrite and bainitic ferrite can lead to reduced strength and create a hardness difference with surrounding hard phases such as martensite. During bending, cracks propagate from the interface with the hard phases, sometimes resulting in reduced extended flange formability and axial crushing properties. Therefore, the combined area fraction of ferrite and bainitic ferrite is set to 5% or more and 60% or less.

[0125] The combined area fraction of ferrite and bainitic ferrite is preferably 10% or more, more preferably 15% or more. Furthermore, the combined area fraction of ferrite and bainitic ferrite is preferably 55% or less, more preferably 50% or less.

[0126] It should be noted that the aforementioned bainitic ferrite refers to the BCC phase remaining after removing carbides and other precipitates from the bainite described later.

[0127] In addition, ferrite content is preferably greater than 0%.

[0128] The area ratio of tempered martensite is between 20% and 80%.

[0129] To achieve the specified strength and elongated flange formability, the tempered martensite content is set to 20% or more in terms of area ratio. The area ratio of tempered martensite is preferably 30% or more, and more preferably 40% or more.

[0130] On the other hand, if the area ratio of tempered martensite is greater than 80%, the ductility will decrease due to excessive high strength. Therefore, the area ratio of tempered martensite is set to 80% or less, preferably 70% or less, and more preferably 60% or less.

[0131] Area fraction of fresh martensite: 20% or less (including 0%)

[0132] Because it leads to a reduction in at least one of the ductility and the formability of the extended flange, the area fraction of fresh martensite is set to 20% or less, preferably 15% or less, and more preferably 10% or less. Alternatively, the fresh martensite may be 0%. Fresh martensite may be 5% or more, or 10% or more.

[0133] Volume fraction of retained austenite: 5% or more but less than 25%

[0134] To ensure high ductility, the retained austenite (retained γ) is set to 5% or more in volume fraction relative to the overall steel structure. The volume fraction of retained γ is preferably 7% or more, more preferably 9% or more. This amount of retained γ includes the volume fraction of retained γ formed adjacent to bainite. If the volume fraction of retained γ is greater than 25%, it leads to reduced strength, reduced elongation flange formability, and deterioration of resistance to delayed fracture. Therefore, the volume fraction of retained γ is set to 25% or less. The volume fraction of retained γ is preferably 20% or less, more preferably 18% or less. Furthermore, regarding retained γ, the proportion of "volume fraction" measured by the measurement method described later can also be regarded as the proportion of "area fraction".

[0135] The total area ratio of ferrite, tempered martensite, fresh martensite, bainite, and retained austenite is 90% or more (including 100%).

[0136] To ensure the specified strength, ductility, and elongation flange formability, the total area ratio of ferrite, tempered martensite, fresh martensite, bainite, and retained austenite is required to be above 90%. Specifically, the bainite contains the aforementioned bainitic ferrite, and the bainite contains internal carbides at a concentration of [missing information - likely a unit of area] per 10 μm. 2 It contains fewer than 20 bainitic ferrite particles. Additionally, the bainite may contain internal carbides at a density of 10 μm. 2 More than 20 bainitic ferrites.

[0137] The area fraction of ferrite with a dissolved Mn content of 2.0% or more (high-Mn ferrite) in the total ferrite is 20% or more and less than 70%.

[0138] When the ferrite content is greater than 0%, it exhibits high work hardening over a wide strain range and has excellent energy absorption characteristics. Therefore, the ferrite content with a solid solution Mn content of 2.0% or more (high Mn ferrite) is 20% or more and less than 70%.

[0139] The area fraction of high-Mn ferrite is preferably 25% or more, more preferably 30% or more. Furthermore, the area fraction of high-Mn ferrite is preferably 65% ​​or less, more preferably 60% or less.

[0140] It should be noted that the area fraction of ferrite with a dissolved Mn content of less than 2.0% by mass (low-Mn ferrite) in the overall ferrite is 30% or more and 80% or less, preferably 75% or less, and more preferably 70% or less. Furthermore, the area fraction of low-Mn ferrite in the overall ferrite is preferably 35% or more, and more preferably 40% or more.

[0141] The area S of the region where the C concentration is 0.5% by mass or higher C≥0.5 The area S relative to the region where the C concentration is 0.3% by mass or higher C≥0.3 Ratio: (S) C≥0.5 / S C≥0.3 )×100: more than 20%

[0142] To ensure high ductility, the area S of the region with a C concentration of 0.5% by mass or higher is limited. C≥0.5 The area S relative to the region where the C concentration is 0.3% by mass or higher C≥0.3 proportion (S) C≥0.5 / S C≥0.3 The percentage is 20% or more (100 × 100). The preferred percentage is 25% or more, and more preferably 30% or more.

[0143] There is no particular upper limit to the above ratio, but the ratio is preferably 70% or less, and more preferably 60% or less.

[0144] The internal carbides are per 10 μm 2 Area fraction of bainitic ferrite with fewer than 20 particles: 3% to 40%

[0145] Bainite content is preferably greater than 0%. When bainite content is greater than 0%, it is achieved by ensuring that the bainite contains internal carbides per 10 μm. 2 With fewer than 20 bainitic ferrite particles, carbon efficiently accumulates in the residual γ surrounding the bainitic ferrite, further improving ductility. To achieve this effect, the internal carbides are at a density of 10 μm per 10 μm. 2 The area fraction of bainitic ferrite with 20 or fewer particles is preferably 3% or more. This area fraction is more preferably 5% or more, and even more preferably 7% or more.

[0146] On the other hand, in order to suppress the decrease in strength, the internal carbides are at a density of 10 μm. 2 The area fraction of bainitic ferrite with 20 or fewer particles is preferably 40% or less. More preferably, this area fraction is 30% or less, and even more preferably 25% or less.

[0147] It should be noted that, in this invention, the internal carbides can exist at a density of 10 μm. 2 More than 20 bainitic ferrites.

[0148] It should be noted that the aforementioned bainitic ferrite refers to the BCC phase remaining after removing carbides and other precipitates from the aforementioned bainite.

[0149] Next, the method for measuring the steel structure of the steel plate of the present invention will be described.

[0150] In determining the area ratios of ferrite, bainitic ferrite, tempered martensite, and fresh martensite, a section of plate thickness parallel to the rolling direction was cut out. Then, the cut section was mirror-polished and etched with a 3% (v / v) nitric acid ethanol solution. At the 1 / 4 thickness position, 10 fields of view were observed using SEM at 5000x magnification, with a field size of 30 μm × 40 μm. Figure 1 An example is shown: a magnified SEM image of the steel microstructure of a steel plate. For example... Figure 1 As shown, ferrite ( Figure 1 In the middle, the reference symbol F) is almost devoid of carbides internally, compared to equiaxed polygonal ferrite. It appears as the darkest region in SEM. Bainitic ferrite ( Figure 1 In the diagram, reference symbol BF) refers to the ferrite microstructure that appears white in SEM as a result of the formation of carbides or residual γ. Fresh martensite and retained austenite (reference symbol BF) Figure 1 The symbols FM and RA in SEM refer to whitish, blocky areas that appear as if the underlying tissue is not visible.

[0151] The area ratio of fresh martensite can be obtained by subtracting the area ratio of the white blocky region from the volume ratio of the residual γ determined by the method described later, which is considered as the area ratio.

[0152] In this invention, when it is difficult to distinguish between ferrite and bainitic ferrite, regions of polygonal ferrite with an aspect ratio ≤ 2.5 are classified as ferrite, and regions with an aspect ratio > 2.5 are classified as bainitic ferrite, and the area ratio is calculated.

[0153] Figure 2 This is a diagram illustrating the method for determining the steel microstructure of the steel plate according to the present invention. (See diagram for example.) Figure 2 As shown, regarding the aspect ratio, the longest major axis length 'a' of the particle is calculated, and the length of the particle that traverses the particle for the longest distance in the direction perpendicular to it is set as the minor axis length 'b'. Then, a / b is taken as the aspect ratio.

[0154] Regarding the internal carbides, per 10 μm 2 The area fraction of bainitic ferrite with fewer than 20 particles was determined by SEM images at 5000x magnification (30 μm × 40 μm field of view) and the number of carbides within each bainitic ferrite. Additionally, the area fraction of each bainitic ferrite particle with fewer than 20 particles per 10 μm was measured, along with the number of carbides within it. 2 The number of carbides in the conversion of bainitic ferrite is calculated by summing the area percentages of each bainitic ferrite with fewer than 20 carbides in the overall microstructure.

[0155] The total area ratio of ferrite, tempered martensite, fresh martensite, bainite, and retained austenite can be obtained by subtracting the area ratio of the remaining microstructures other than these from the overall steel microstructure. Here, the area ratio of carbides is very small, and therefore is included in the area ratios of the aforementioned microstructures (ferrite, tempered martensite, fresh martensite, bainite, and retained austenite). As remaining microstructures, precipitates other than pearlite and carbides can be cited; their area ratios can be determined using SEM.

[0156] The volume fraction of retained austenite (retained γ) was determined by X-ray diffraction after chemical grinding at a depth of 1 / 4 thickness from the steel plate surface. A Co-Kα X-ray source was used for the incident X-rays, and the volume fraction of retained austenite was calculated from the intensity ratio of the (200), (211), and (220) planes of ferrite to the (200), (220), and (311) planes of austenite. Since the retained γ is randomly distributed, the volume fraction of retained γ obtained by X-ray diffraction can be treated as the area fraction of retained γ in the steel microstructure.

[0157] The area fraction of ferrite with a solid Mn content of 2.0% or more (high Mn ferrite) and the area S of the region with a C concentration of 0.5% or more. C≥0.5The area S of regions with a C concentration of 0.3% by mass or higher C≥0.3 In the measurement, at the 1 / 4 position of the plate thickness section parallel to the rolling direction, a field emission electron microscopy (FE-EPMA) JXA-8500F manufactured by NEC was used, with an accelerating voltage of 6kV and an irradiation current of 7×10⁻⁶ kV. -8 A. The Mn and C concentration distributions are mapped and analyzed using the method with the smallest beam diameter, and then measured. However, when measuring the C concentration distribution, to eliminate the influence of contamination, the background component is subtracted by ensuring that the average C concentration obtained through analysis is equal to the carbon content of the base material (C content of the steel plate). That is, if the average measured carbon content is greater than the carbon content of the base material, the increase is considered as contamination, and this increase is subtracted from the analytical values ​​at each location to obtain the true C concentration at each location.

[0158] Furthermore, ferrite and bainitic ferrite can be distinguished by the SEM observations described above. Therefore, by performing SEM observations on the regions where C concentration distribution mapping analysis was conducted, the proportion of high-Mn ferrite and low-Mn ferrite within the ferrite can be determined.

[0159] The tensile strength (TS) of the steel plate of the present invention is 780 MPa or more. More preferably, it is 980 MPa or more. Regarding the upper limit of the tensile strength, from the viewpoint of taking into account other properties, it is preferably 1469 MPa or less, and more preferably 1320 MPa or less.

[0160] In the steel sheet of the present invention, the total elongation T-E1 is ensured to be 18.0% or more when T-E is 780 MPa or more and less than 980 MPa, 16.0% or more when T-E is 980 MPa or more and less than 1180 MPa, 14.0% or more when T-E is 1180 MPa or more and less than 1320 MPa, and 13.0% or more when T-E is 1320 MPa or more, thereby significantly improving the stability of the forming process. The hole expansion ratio λ is ensured to be 30% or more. There is no particular upper limit to λ; from the viewpoint of considering other properties, λ is preferably 90% or less, and more preferably 80% or less, at any strength level.

[0161] Furthermore, in the steel plate of the present invention, from the viewpoint of ensuring excellent energy absorption characteristics during impact, it is preferable that the area (energy absorption during deformation) in the nominal stress-nominal strain curve of the tensile test up to the maximum stress is 13000 MPa·% or more. More preferably, it is 14000 MPa·% or more.

[0162] The steel plate of the present invention described above can also be a steel plate with a zinc coating on its surface (one side or both sides). The coating can be any of the hot-dip galvanized coating or electroplated coating.

[0163] Next, the method for manufacturing the steel plate of the present invention will be described.

[0164] The following description will focus on the first embodiment, which involves a dwell time under specified conditions during the cooling process following the annealing temperature. The second embodiment will also describe the case where a dwell time under specified conditions is performed during the cooling process following the annealing temperature.

[0165] It should be noted that the temperatures specified in each process of this invention refer to the surface temperature of the slab (steel billet) or steel plate, which can be measured using a radiation thermometer or the like. Furthermore, the average cooling rate (°C / s) is calculated as “(cooling start temperature - cooling stop temperature)(°C) / cooling time(s)”, and the average heating rate (°C / s) is calculated as “(heating stop temperature - heating start temperature)(°C) / heating time(s)”.

[0166] in addition, Figure 3 This diagram illustrates the method for manufacturing the steel sheet according to the present invention, specifically showing the time-varying surface temperature of the slab (steel billet) or steel sheet. Details of each step, including this temperature variation over time, are described below. Figure 3 (a) Shows the time-varying surface temperature of the slab (steel billet) or steel sheet in the steel sheet manufacturing method of the first embodiment (without a dwell process). Additionally, Figure 3 (b) Shows the time change of the surface temperature of the slab (steel billet) or steel plate in the steel plate manufacturing method of the second embodiment (in the case of dwell treatment).

[0167] <First Implementation Method (No-Distance Processing)>

[0168] In the manufacturing method of the steel plate of the first embodiment of the present invention, as follows: Figure 3As shown in (a), after hot rolling and cold rolling of the steel billet having the above-mentioned composition, the resulting cold-rolled steel sheet is annealed. The annealing process includes: a holding process at an annealing temperature of 775°C or higher and 830°C or lower; a cooling process at an average cooling rate CR1 of 0.01°C / s or higher and 5°C / s or lower within a temperature range from the annealing temperature to -15°C; a cooling process at an average cooling rate CR2 of 3°C / s or higher within a temperature range from -15°C to a cooling stop temperature of 200°C or higher and 300°C; a heating process at an average heating rate of 2°C / s or higher within a temperature range from the cooling stop temperature to 380°C; a holding process at an average cooling rate CR4 of 0.01 to 5°C / s for 20 s or more and 3000 s within a temperature range of 340°C or higher and 590°C; and a cooling process at an average cooling rate CR5 of 0.1°C / s or higher to a temperature of 50°C or lower.

[0169] Hot rolling

[0170] For hot rolling of steel billets, there are methods such as heating the slab before rolling, rolling the slab directly after continuous casting without heating, and short-time heating treatment followed by rolling of the slab after continuous casting. Hot rolling can be carried out using conventional methods; for example, the slab heating temperature can be 1100°C or higher. Alternatively, the slab heating temperature can be 1300°C or lower. Furthermore, the soaking time can be 20 minutes or higher. Alternatively, the soaking time can be 300 minutes or lower. Furthermore, the finishing rolling temperature can be above the Ar3 phase transformation point. Alternatively, the finishing rolling temperature can be below the Ar3 phase transformation point +200°C. Furthermore, the coiling temperature can be set to 400°C or higher. Alternatively, the coiling temperature can be set to 720°C or lower. From the viewpoint of suppressing thickness variation and consistently ensuring high strength, controlling the coiling temperature is preferable. Specifically, the coiling temperature is preferably set to 430°C or higher. Alternatively, the coiling temperature is preferably set to 630°C or lower.

[0171] cold rolling

[0172] In cold rolling, a rolling percentage (cumulative rolling percentage) of 30% or higher is acceptable. Furthermore, a rolling percentage (cumulative rolling percentage) of 85% or lower is preferable. From the viewpoint of consistently ensuring high strength and minimizing anisotropy, this control is preferred. Specifically, a rolling percentage of 35% or higher is preferred. Additionally, a rolling percentage of 85% or lower is preferred. It should be noted that when the rolling load is high, a softening annealing treatment can be performed at 450–730°C using a CAL (continuous annealing line) or BAF (box annealing furnace).

[0173] annealing

[0174] After hot rolling and cold rolling, the steel billet with the above-mentioned composition is annealed under the following specified conditions. There are no particular limitations on the annealing equipment, but from the viewpoint of ensuring productivity and desired heating and cooling rates, it is preferable to carry out the annealing in a continuous annealing line (CAL) or a continuous hot-dip galvanizing line (CGL).

[0175] Maintain at an annealing temperature above 775°C and below 830°C

[0176] To ensure the specified area ratio of ferrite, tempered martensite, bainite, and residual γ, the annealing temperature is set to 775°C or higher and 830°C or lower. To ensure that the total area ratio of ferrite and bainitic ferrite is 5% or higher and 60% or lower, the annealing temperature is preferably adjusted to annealing in the ferrite + austenite dual-phase region. The preferred annealing temperature is 780°C or higher.

[0177] On the other hand, when the annealing temperature is above 830°C, the γ particle size becomes too large, the diffusion distance of C atoms required to obtain the desired residual γ area becomes longer, and the required amount of ferrite cannot be obtained, thereby reducing ductility. Therefore, the annealing temperature is set below 830°C.

[0178] The holding time at the annealing temperature is not particularly limited, but from the viewpoint of allowing sufficient recrystallization and reverse phase transformation from the cold-rolled structure, it is preferably set to 10 seconds or more, and more preferably 30 seconds or more. In addition, from the viewpoint of avoiding excessive coarsening of the structure after recrystallization and reverse phase transformation, the holding time is preferably set to 600 seconds or less, and more preferably 500 seconds or less.

[0179] Cooling is performed within a temperature range from the annealing temperature to -15°C at an average cooling rate CR1 of 0.01°C / s or higher and 5°C / s or lower.

[0180] Ferrite formed during annealing may be accompanied by Mn diffusion, resulting in a lower Mn concentration in the ferrite. However, by cooling at an average cooling rate CR1 of 0.01℃ / s or higher and 5℃ / s or lower within a temperature range from the annealing temperature to -15℃, ferrite with a high Mn concentration can be obtained without Mn diffusion. As a result, ferrite with a wide Mn concentration distribution can be obtained, exhibiting high work hardening ability over a wider strain range during deformation and improved energy absorption characteristics.

[0181] When the average cooling rate CR1 in the temperature range from the annealing temperature to -15°C is greater than 5°C / s, a sufficient amount of ferrite with high Mn concentration cannot be obtained. Therefore, the average cooling rate CR1 is set to 5°C / s or less, preferably 4°C / s or less.

[0182] On the other hand, when the average cooling rate CR1 in the temperature range from the annealing temperature to -15°C is less than 0.01°C / s, Mn diffusion occurs, and a sufficient amount of ferrite with high Mn concentration cannot be obtained. Therefore, the average cooling rate CR1 is set to 0.01°C / s or more. Preferably, it is 0.1°C / s or more.

[0183] Here, the average cooling rate CR1 (°C / s) is obtained by "15 (°C) / (cooling time (s) from annealing temperature to annealing temperature - 15°C)".

[0184] Cooling is performed at an average cooling rate CR2 of 3°C or higher within a temperature range from the annealing temperature to the cooling stop temperature above 200°C and below 300°C.

[0185] After cooling at an average cooling rate CR1 of 0.01°C / s or more and 5°C / s or less within a temperature range from the annealing temperature to -15°C, cooling is then performed at an average cooling rate CR2 of 3°C / s or more within a temperature range up to a cooling stop temperature of 200°C or more and 300°C. This results in a specified amount of tempered martensite and residual γ in the final microstructure, improving strength, elongation flange formability, and energy absorption characteristics. When the average cooling rate CR2 is less than 3°C / s, excessive formation of ferrite, bainite, and pearlite occurs during cooling, thereby reducing strength, elongation flange formability, and energy absorption characteristics. Therefore, the average cooling rate CR2 for the temperature range from -15°C to a cooling stop temperature of 200°C or more and 300°C is set to 3°C / s or more. The average cooling rate CR2 is preferably 5°C / s or more, and more preferably 8°C / s or more. Furthermore, if the average cooling rate CR2 within this temperature range is too high, the plate shape will deteriorate. Therefore, the average cooling rate CR2 within this temperature range is preferably set to 100°C / s or less. More preferably, the average cooling rate CR2 is 50°C / s or less.

[0186] To ensure the specified amount of tempered martensite and retained austenite, the cooling stop temperature is set to 200°C or higher. Preferably, the cooling stop temperature is 210°C or higher, more preferably 220°C or higher. If the cooling stop temperature exceeds 300°C, a large amount of blocky, untransformed austenite remains, increasing the amount of fresh martensite during final cooling and reducing the formability of the extended flange. Therefore, the cooling stop temperature is set to 300°C or lower. Preferably, the cooling stop temperature is 280°C or lower.

[0187] Here, the average cooling rate CR2 (°C / s) is calculated by “(annealing temperature (°C) - 15 (°C) - cooling stop temperature (°C)) / (cooling time (s) from annealing temperature - 15°C to cooling stop temperature)”.

[0188] Heating is performed at an average heating rate of 2°C / s or higher within the temperature range from the aforementioned cooling stop temperature to 380°C.

[0189] By heating for a short time within the temperature range from the aforementioned cooling stop temperature to 380°C, carbide precipitation can be suppressed, ensuring high ductility. Furthermore, when the martensite or bainite formed during cooling is reheated to 380°C or higher, upper bainite is formed. If the average heating rate up to 380°C is slow, these effects cannot be achieved. As a result, the residual γ content decreases, and ductility is reduced. Therefore, the average heating rate within the temperature range from the cooling stop temperature to 380°C is set to 2°C / s or higher. From the viewpoint of suppressing carbide precipitation and promoting the formation of upper bainite during reheating, the average heating rate is preferably set to 5°C / s or higher, and more preferably 10°C / s or higher. The upper limit of the above average heating rate is not particularly limited, but is preferably 50°C / s or lower, and more preferably 30°C / s or lower.

[0190] Here, the average heating rate (°C / s) is calculated by “(380 (°C) (heating end temperature) - cooling stop temperature (°C)) / (heating time from cooling stop temperature to 380°C (s))”.

[0191] It should be noted that the heating end temperature (380°C) here refers to the end temperature when the average heating temperature is used as the calculation object. After the above heating, heating can continue before staying in the temperature range of 340°C or higher and 590°C as described later.

[0192] Within a temperature range of 340℃ to 590℃, maintain an average cooling rate CR4 of 0.01–5℃ / s for a duration of 20 seconds to 3000 seconds.

[0193] From the perspective of stabilizing residual γ by distributing C to them and improving ductility, and from the perspective of increasing λ by subdividing the regions where untransformed γ is distributed in a blocky form through bainitic phase transformation, a temperature range of 340°C to 590°C is maintained (slowly cooled) for 20 to 3000 s. Furthermore, to suppress the formation of blocky structures caused by excessive C distribution to residual γ, and to increase λ through self-tempering of fresh martensite, slow cooling is performed within this temperature range at an average cooling rate CR4 of 0.01–5°C / s. When the average cooling rate CR4 is below 0.01°C / s, C excessively distributes to residual γ, resulting in a decrease in λ due to the formation of blocky structures. Therefore, the average cooling rate CR4 is set to 0.01°C / s or higher.

[0194] On the other hand, if the average cooling rate CR4 is greater than 5 °C / s, the distribution of C to the residual γ is suppressed, and a sufficient amount of C-enriched region cannot be obtained. Furthermore, the formation of fresh martensite leads to the deterioration of γ. Therefore, the average cooling rate CR4 is set to be below 5 °C / s.

[0195] Here, the average cooling rate CR4 is calculated by "(cooling start temperature (°C) - cooling stop temperature (°C)) / (cooling time from cooling start temperature to cooling stop temperature (s))".

[0196] Here, the cooling start temperature and cooling stop temperature are not particularly limited as long as they are in the range of 340°C or higher and 590°C or lower, but the cooling start temperature is preferably set to 360°C or higher. Furthermore, the cooling start temperature is preferably 580°C or lower. Furthermore, the cooling stop temperature is preferably 350°C or higher. Furthermore, the cooling stop temperature is preferably 450°C or lower.

[0197] If the cooling start temperature exceeds 380°C, after the above-mentioned heating process at an average heating rate of 2°C / s or higher, heating is performed again until the cooling start temperature is reached. If the cooling start temperature is below 380°C, after the above-mentioned heating process at an average heating rate of 2°C / s or higher, cooling is performed again until the cooling start temperature is reached.

[0198] It should be noted that holding (dwelling, slow cooling) within a temperature range of 340°C to 590°C can be used for both hot-dip galvanizing and alloying hot-dip galvanizing. That is, in the aforementioned process of holding at an average cooling rate CR4 of 0.01–5°C / s, the steel sheet can be subjected to either hot-dip galvanizing or alloying hot-dip galvanizing. When performing hot-dip galvanizing, it is preferable to immerse the steel sheet in a zinc bath at a temperature of 440°C to 500°C, and then adjust the coating adhesion by methods such as gas wiping. A zinc bath with an Al content of 0.10% to 0.22% is preferably used for hot-dip galvanizing. Alternatively, as an alloying hot-dip galvanizing process, an alloying treatment of the zinc coating can be performed after the hot-dip galvanizing. When performing the alloying treatment of the zinc coating, it is preferable to perform it within a temperature range of 470°C to 590°C.

[0199] It should be noted that this process involves cooling (dwelling, slow cooling), but as long as the temperature range, dwell time range, and average cooling rate CR4 range mentioned above are met, hot-dip galvanizing and zinc coating alloying can be performed in this process. Temperature can also be increased during hot-dip galvanizing and zinc coating alloying.

[0200] With an average cooling rate CR5: cooling to a temperature below 50°C at a rate of 0.1°C / s or higher.

[0201] Subsequently, from the viewpoint of preventing softening caused by excessive tempering and reduced ductility caused by carbide precipitation, the steel sheet is cooled to a temperature below 50°C at an average cooling rate CR5 of 0.1°C / s or higher. From the viewpoint of stabilizing stamping formability by adjusting surface roughness and flattening the sheet shape, and from the viewpoint of increasing YS, the steel sheet can be surface-rolled. The surface-rolling elongation is preferably set to 0.1 to 0.5%. In addition, the sheet shape can also be flattened using a leveling machine. The average cooling rate CR5 up to the temperature below 50°C is preferably 5°C / s or higher. Furthermore, the average cooling rate CR5 is preferably 100°C / s or lower.

[0202] Here, the average cooling rate CR5 is calculated by "(340 (°C) (cooling start temperature) - cooling stop temperature below 50°C (°C)) / (cooling time from cooling start temperature to cooling stop temperature (s))".

[0203] From the viewpoint of improving the formability of the extended flange, a low-temperature heat treatment can also be performed at 100–300°C for 30 seconds to 10 days after the aforementioned annealing (heat treatment) or after surface rolling. This treatment allows hydrogen that has penetrated the steel sheet during tempering or annealing of the martensite generated during final cooling or surface rolling to escape from the steel sheet. Through low-temperature heat treatment, hydrogen levels can be reduced to less than 0.1 ppm.

[0204] Alternatively, electroplating can be performed. That is, the steel sheet can be electroplated with zinc after the above-mentioned cooling process at an average cooling rate CR5 of 0.1℃ / s or higher. After electroplating, from the viewpoint of reducing hydrogen in the steel, the above-mentioned low-temperature heat treatment is preferred.

[0205] <Second Implementation Method (with Dwell Time Processing)>

[0206] In the manufacturing method of the steel plate of the second embodiment of the present invention, as follows: Figure 3(b) shows that after hot rolling and cold rolling of the steel billet having the above composition, the resulting cold-rolled steel sheet is annealed. The annealing process includes: a holding process at an annealing temperature of 775°C or higher and 830°C or lower; a cooling process at an average cooling rate CR1 of 0.01°C / s or higher and 5°C / s or lower within a temperature range from the annealing temperature to -15°C; a cooling process at an average cooling rate CR2A of 3°C / s or higher within a temperature range from -15°C to 500°C; and a cooling process at an average cooling rate CR2A of 3°C / s or higher within a temperature range from 500°C to a holding temperature of 320°C or higher within a temperature range from the martensitic transformation start temperature Ms to the holding temperature. Processes with a temperature CR3 of 10°C / s or less and a dwell time of 10 to 60 seconds or less; processes with an average cooling rate CR2B of 3°C / s or more in the temperature range from the dwell temperature to the cooling stop temperature of 200°C to 300°C; processes with an average heating rate of 2°C / s or more in the temperature range from the cooling stop temperature to 380°C; processes with an average cooling rate CR4 of 0.01 to 5°C / s and a dwell time of 20 to 3000 seconds in the temperature range of 340°C to 590°C; and processes with an average cooling rate CR5 of 0.1°C / s or more to a temperature of 50°C or less.

[0207] In the second embodiment, hot rolling and cold rolling can be carried out under the same conditions as in the first embodiment.

[0208] Furthermore, in the second embodiment, the process of holding at an annealing temperature of 775°C or higher and 830°C or lower, and the process of cooling at an average cooling rate CR1 of 0.01°C / s or higher and 5°C / s or lower, can be performed under the same conditions as in the first embodiment.

[0209] Furthermore, in the second embodiment, the process of cooling at an average cooling rate of CR2: 3°C / s or higher in the first embodiment is configured as a process of cooling at an average cooling rate of CR2A: 3°C / s or higher, a process of holding at an average cooling rate of CR3: 10°C / s or lower for 10s or more and 60s or less, and a process of cooling at an average cooling rate of CR2B: 3°C / s or higher.

[0210] Furthermore, in the second embodiment, the process of heating at an average heating rate of 2°C / s or higher, the process of holding at an average cooling rate CR4 of 0.01 to 5°C / s for 20 seconds or more and 3000 seconds or less, and the process of cooling at an average cooling rate CR5 of 0.1°C / s or higher can be performed under the same conditions as in the first embodiment.

[0211] In addition, in the second embodiment, as other treatments, hot-dip galvanizing, surface rolling, low-temperature heat treatment after annealing, electroplating, etc., can also be carried out under the same conditions as in the first embodiment.

[0212] The following mainly describes the process of cooling at an average cooling rate of CR2A: 3℃ / s or higher, the process of cooling at an average cooling rate of CR3: 10℃ / s or lower for 10s or more and 60s or less, and the process of cooling at an average cooling rate of CR2B: 3℃ / s or higher.

[0213] Cooling at an average cooling rate of CR2A above 3°C / s within a temperature range from -15°C to 500°C.

[0214] Within a temperature range from 500℃ to the dwell temperature above the martensitic transformation start temperature Ms and above 320℃, with an average cooling rate CR3 of less than 10℃ / s, the dwell time is more than 10s and less than 60s.

[0215] Cooling is performed at an average cooling rate of CR2B of 3°C or higher within a temperature range from the dwell temperature to the cooling stop temperature of 200°C to 300°C.

[0216] In the process of cooling from the annealing temperature of -15°C to a cooling stop temperature of 200°C or higher and 300°C or lower, a slow cooling process is included, which involves a dwell (slow cooling) process with an average cooling rate of CR3 of 10°C or lower and a dwell time of 10 seconds or higher and 60 seconds or lower within a temperature range from 500°C to a dwell stop temperature of 320°C or higher than the martensitic transformation start temperature Ms. This process can generate bainitic ferrite with low carbide density, generate residual γ with high C concentration adjacent to the bainitic ferrite, and obtain steel sheets with better ductility.

[0217] When the stopping temperature is below Ms or below 320°C, martensite is formed first, and the bainitic phase transformation proceeds excessively due to the swing back phenomenon, sometimes leading to a decrease in strength and elongation flange formability.

[0218] On the other hand, if the temperature exceeds 500°C, the driving force for the bainitic phase transformation decreases, and the amount of bainitic ferrite with low carbide density decreases. Therefore, when the residence treatment is performed during cooling, the residence temperature range is Ms or higher and 320°C or higher, and is set to 500°C or lower. This temperature range is preferably 380°C or higher, more preferably 420°C or higher. Furthermore, this temperature range is preferably 480°C or lower, more preferably 460°C or lower.

[0219] The reasons for setting the average cooling rate CR2A to 3°C / s or more and the average cooling rate CR2B to 3°C / s or more are the same as those explained in the first embodiment for setting the average cooling rate CR2 to 3°C / s or more. Both average cooling rates CR2A and CR2B are preferably 5°C / s or more, more preferably 8°C / s or more. Furthermore, if both average cooling rates CR2A and CR2B are too high, the plate shape will deteriorate; therefore, average cooling rates CR2A and CR2B are preferably set to 100°C / s or less. Both average cooling rates CR2A and CR2B are more preferably 50°C / s or less.

[0220] Here, the average cooling rate CR2A (°C / s) is obtained by “(annealing temperature (°C) - 15 (°C) - 500 (°C)) / (cooling time (s) from annealing temperature - 15°C to 500°C)”.

[0221] In addition, the average cooling rate CR2B (°C / s) is obtained by dividing (the dwell temperature (°C) - the cooling stop temperature (°C)) by (the cooling time (s) from the dwell temperature to the cooling stop temperature).

[0222] If the average cooling rate CR3 is greater than 10℃ / s, the amount of bainite phase decreases, and the aforementioned improvement in ductility becomes insufficient. Therefore, when a residence treatment is performed during cooling, the average cooling rate CR3 is set to be 10℃ / s or less.

[0223] Furthermore, if the dwell time is less than 10 s, the amount of bainite phase decreases, and the aforementioned improvement in ductility becomes insufficient. On the other hand, if it is greater than 60 s, carbon accumulates from bainite to bulky, untransformed γ, leading to an increase in the amount of residual bulky structure. Therefore, when the dwell time is performed during cooling, it is set to be 10 s or more and 60 s or less. From the viewpoint of ensuring bainitic ferrite and retained austenite and improving ductility, a dwell time of 20 s or more is preferred. In addition, from the viewpoint of improving the formability of the extended flange by reducing the bulky structure, a dwell time of 50 s or less is preferred.

[0224] Here, the average cooling rate CR3 (°C / s) is calculated by "(500 (°C) - dwell temperature (°C)) / cooling time from 500°C to dwell temperature (s))".

[0225] It should be noted that the martensitic transformation initiation temperature Ms can be determined as follows: using a cylindrical test piece (3mm in diameter × 10mm in height), after holding it at the specified annealing temperature using a Formaster testing machine, the height change of the test piece is measured when it is quenched with helium, and thus the temperature is determined.

[0226] The steel plate of the present invention preferably has a thickness of 0.5 mm or more. Alternatively, the thickness is preferably 2.0 mm or less.

[0227] <Components>

[0228] Next, the components of the present invention and their manufacturing method will be described.

[0229] The component of the present invention is formed by performing at least one of forming and joining processes on the steel plate of the present invention. Furthermore, the manufacturing method of the component of the present invention includes a step of forming and joining processes on the steel plate of the present invention to form the component.

[0230] The steel sheet of the present invention has a tensile strength of 780 MPa or higher, and exhibits high ductility, excellent extension flange formability, and excellent energy absorption characteristics during impact. Therefore, components obtained using the steel sheet of the present invention also possess high strength, exhibiting high ductility, excellent extension flange formability, and excellent energy absorption characteristics during impact compared to existing high-strength components. Furthermore, the use of components of the present invention allows for weight reduction. Therefore, components of the present invention can be applied, for example, to vehicle body frame components. The components of the present invention also include welded joints.

[0231] Forming processes can utilize general processing methods such as stamping without limitation. Furthermore, joining processes can utilize general welding methods such as spot welding and arc welding, as well as riveting and rivet joining without limitation.

[0232] Example

[0233] Cold-rolled steel sheets with a thickness of 1.4 mm having the composition shown in Table 1 were processed under the annealing conditions shown in Table 2 to manufacture the steel sheets of the present invention and the comparative examples.

[0234] Each cold-rolled steel sheet is obtained by hot rolling (slab heating temperature: 1200℃, soaking time: 60 minutes, finishing rolling temperature: 900℃, coiling temperature: 500℃) and cold rolling (rolling rate (cumulative rolling rate): 50%) of steel billets with the composition shown in Table 1.

[0235] In Table 2, the martensitic phase transformation initiation temperature Ms was determined by the following method: using a cylindrical test piece (3 mm in diameter × 10 mm in height), after holding it at the specified annealing temperature using a Formaster testing machine, the height change of the test piece was measured when it was quenched with helium, and thus the temperature was determined.

[0236]

[0237] It should be noted that for a certain type of steel sheet (cold-rolled steel sheet: CR), hot-dip galvanizing is performed in a process within a temperature range of 340°C to 590°C, with an average cooling rate of CR4: 0.01~5°C / s, and a holding time of 20 seconds to 3000 seconds, to produce hot-dip galvanized steel sheet (GI). Here, the steel sheet is immersed in a zinc bath at a temperature of 440°C to 500°C for hot-dip galvanizing, and then the coating adhesion is adjusted by gas wiping or other methods. The zinc bath used for hot-dip galvanizing has an Al content of 0.10% to 0.22%. Furthermore, for a certain type of hot-dip galvanized steel sheet, as an alloying hot-dip galvanizing treatment, alloying treatment is performed after the above hot-dip galvanizing treatment to produce alloyed hot-dip galvanized steel sheet (GA). Here, alloying treatment is performed within a temperature range of 460°C to 590°C. In addition, for some steel sheets (cold-rolled steel sheets: CR), electroplating is carried out after a cooling process with an average cooling rate of CR5: 0.1℃ / s or higher to produce electro-galvanized steel sheets (EG).

[0238] In addition, some steel plates are set to have a dwell time under annealing conditions (in Table 2, the steel plates whose values ​​are not "-" in the columns of CR3, dwell time, and dwell stop temperature). For CR2A and CR2B in this case, CR2A=CR2B, and CR2A and CR2B together are represented as CR2 in Table 2.

[0239] The steel microstructure was determined using the method described above. The results are shown in Table 3.

[0240]

[0241] JIS 5 tensile test pieces and reamer test pieces were cut from the obtained steel plate and tensile tests were performed (according to JIS Z2241 (2011). TS and T-El are shown in Table 3.

[0242] A tensile strength of 780 MPa or higher is considered to be of excellent strength.

[0243] The following conditions are considered to indicate excellent ductility: total elongation (T-El) of 18.0% or more when TS is less than 980 MPa, 16.0% or more when TS is 980 MPa or more and less than 1180 MPa, 14.0% or more when TS is 1180 MPa or more and less than 1320 MPa, and 13.0% or more when TS is 1320 MPa or more.

[0244] In addition, the area up to the maximum stress in the nominal stress-nominal strain curve of the tensile test is regarded as the energy absorbed during deformation, and the case of 13000 MPa·% or more is regarded as having excellent energy absorption characteristics during impact.

[0245] Furthermore, regarding the formability of the extended flange, test pieces for hole expansion were cut from the heat-treated steel sheet and evaluated using a hole expansion test as specified in the Japanese Iron and Steel Federation standard JFST1001. Specifically, a 100mm × 100mm square sample was punched using a punching tool with a punch diameter of 10mm and a die diameter of 10.3mm (13% clearance). Then, a conical punch with a 60-degree apex angle was used to expand the hole until a crack penetrating the plate thickness was formed, with the burrs from the punching process on the outer side. At this point, d0 is the initial hole diameter (mm), and d is the hole diameter at crack initiation (mm). The hole expansion rate λ (%) is calculated as {(d-d0) / d0} × 100, as shown in Table 3. Steel with λ of 30% or more was considered to have excellent extended flange formability (hole expansion).

[0246] The examples of the present invention shown in Tables 2 and 3 have excellent strength, ductility, extended flange formability, and energy absorption characteristics, while the comparative examples are inferior in one aspect.

[0247] Furthermore, it is understood that, regarding components obtained by forming, joining, and forming and joining processes using the steel plate of the present invention, since the steel plate of the present invention has high strength, high ductility, excellent extension flange forming properties, and excellent energy absorption characteristics during impact, it also has high strength, high ductility, excellent extension flange forming properties, and excellent energy absorption characteristics during impact, just like the steel plate of the present invention.

[0248] Industrial availability

[0249] This invention has extremely high ductility, excellent elongation flange formability and excellent energy absorption characteristics, and can be preferably applied to stamped parts used in automobiles, home appliances and other industries that have undergone stamping processes.

[0250] Symbol Explanation

[0251] F Ferrite

[0252] TM tempered martensite

[0253] BF bainitic ferrite

[0254] FM Fresh Martensite

[0255] RA Residual Austenite

Claims

1. A steel plate having, by mass percent, a composition comprising: C: 0.06% or more and 0.25% or less; Si: 0.4% or more and 2.5% or less; Mn: 1.5% or more and 3.5% or less; P: 0.10% or less; S: 0.010% or less; sol.Al: 1.0% or less; N: 0.015% or less; with the balance being Fe and unavoidable impurities. It also possesses the following steel microstructure: by area percentage, the total amount of ferrite and bainitic ferrite is 5% to 60%; tempered martensite is 20% to 80%; fresh martensite is 20% to 0% (inclusive); retained austenite by volume percentage is 5% to 25%; and the total area percentage of ferrite, tempered martensite, fresh martensite, bainite, and retained austenite is 90% to 100% (inclusive). When the ferrite content is greater than 0%, the proportion of ferrite with a solid solution Mn content of 2.0% or more in the total ferrite is 20% or more and less than 70% by area ratio. The area S of the region where the C concentration is 0.5% by mass or higher C≥0.5 The area S relative to the region where the C concentration is 0.3% by mass or higher C≥0.3 Ratio: (S) C≥0.5 / S C≥0.3 )×100 is more than 20%.

2. The steel plate according to claim 1, wherein, As part of the composition, the material further comprises, by mass%, one or more of the following: Ti: less than 0.1%, B: less than 0.01%, Cu: less than 1%, Ni: less than 1%, Cr: less than 1.0%, Mo: less than 0.5%, V: less than 0.5%, Nb: less than 0.1%, Zr: less than 0.2%, W: less than 0.2%, Ca: less than 0.0040%, Ce: less than 0.0040%, La: less than 0.0040%, Mg: less than 0.0030%, Sb: less than 0.1%, and Sn: less than 0.1%.

3. The steel plate according to claim 1 or 2, wherein, The steel microstructure also contains internal carbides at a concentration of 3% to 40% by area, per 10 μm. 2 It contains less than 20 bainitic ferrites.

4. The steel plate according to any one of claims 1 to 3, wherein, It has a zinc coating on the surface.

5. A component made of steel plate according to any one of claims 1 to 4.

6. A method for manufacturing a steel plate, wherein, After hot rolling and cold rolling of a steel billet having the composition described in claim 1 or 2, the resulting cold-rolled steel sheet is annealed. The annealing process includes, in sequence: The process of holding the product at an annealing temperature above 775°C and below 830°C; A cooling process performed within a temperature range from the annealing temperature to -15°C with an average cooling rate CR1 of 0.01°C / s or more and 5°C / s or less. A cooling process performed at an average cooling rate of CR2: 3°C / s or higher within a temperature range from annealing temperature to cooling stop temperature above 200°C and below 300°C. A heating process performed at an average heating rate of 2°C / s or more within a temperature range from the cooling stop temperature to 380°C; Processes involving a temperature range of 340°C to 590°C with an average cooling rate CR4 of 0.01–5°C / s and a dwell time of 20 seconds to 3000 seconds; and A process that cools to a temperature below 50°C at an average cooling rate CR5 of 0.1°C / s or higher.

7. A method for manufacturing a steel plate, wherein, After hot rolling and cold rolling of a steel billet having the composition described in claim 1 or 2, the resulting cold-rolled steel sheet is annealed. The annealing process includes, in sequence: The process of holding the product at an annealing temperature above 775°C and below 830°C; A cooling process performed within a temperature range from the annealing temperature to -15°C with an average cooling rate CR1 of 0.01°C / s or more and 5°C / s or less. A cooling process performed at an average cooling rate of CR2A: 3°C / s or higher within a temperature range from -15°C to 500°C. A process with an average cooling rate of CR3 of less than 10℃ and a dwell time of more than 10 seconds and less than 60 seconds within a temperature range from 500℃ to the dwell temperature above the martensitic transformation start temperature Ms and above the dwell temperature above 320℃; A cooling process performed at an average cooling rate of CR2B: 3°C / s or more within a temperature range from the stated stopping temperature to a cooling stopping temperature of 200°C or more and 300°C or less; A heating process performed at an average heating rate of 2°C / s or more within a temperature range from the cooling stop temperature to 380°C; Processes involving a temperature range of 340°C to 590°C with an average cooling rate CR4 of 0.01–5°C / s and a dwell time of 20 seconds to 3000 seconds; and A process that cools to a temperature below 50°C at an average cooling rate CR5 of 0.1°C / s or higher.

8. The method for manufacturing a steel plate according to claim 6 or 7, wherein, In the process of holding at the average cooling rate CR4: 0.01~5℃ / s, the steel plate is subjected to hot-dip galvanizing or alloyed hot-dip galvanizing.

9. The method for manufacturing a steel plate according to claim 6 or 7, wherein, This includes a process of electroplating zinc after a cooling process at an average cooling rate of CR5 of 0.1°C / s or higher.

10. A method for manufacturing a component, comprising a step of forming a component by performing at least one of forming or joining processes on a steel plate as described in any one of claims 1 to 4.

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