Steel sheet, member, and method for manufacturing the same

Abstract

The present invention provides a steel sheet having high strength, high ductility, excellent stretch flange formability, and excellent laser weldability, a member, and a method for manufacturing the same. A steel sheet having a composition consisting of, in mass%, C: 0.06 to 0.25%, Si: 0.4 to 2.5%, Mn: 1.5 to 3.5%, P: 0.02% or less, S: 0.01% or less, sol. Al: less than 1.0%, N: less than 0.015%, the balance consisting of Fe and unavoidable impurities, and having a steel structure in which the polygonal ferrite is 10% or less (including 0%) in area ratio, the tempered martensite is 40% or more, the fresh martensite is 20% or less (including 0%) in area ratio, the carbide in the center of 10 μm 2 or less in number is 3 to 40% of the bainite ferrite, the residual austenite is 5 to 20% in volume ratio, S C≥0.5 / S C≥0.3 × 100 is 20% or more.

Description

Technical Field

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

[0002] In recent years, due to the increasing demand for lightweight automotive bodies, the application of 980-1180MPa high-strength steel sheets in automotive frame components and sheet metal parts is being promoted. However, when using 980-1180MPa high-strength steel sheets in automotive parts, stamping cracks are prone to occur due to decreased ductility and reduced tensile flange formability. Therefore, for these high-strength steel sheets, better formability than before is desired. In addition, efforts are being made to achieve lightweight and high rigidity of components through the effective use of laser welding. For example, efforts are being made to effectively utilize welded blanks, which are formed by joining steel sheets of different thicknesses and strengths before stamping; to achieve closed section structures through laser welding of the ends and flanges of stamped parts; to shorten flanges; and to improve the strength of welded areas compared to conventional spot welding.

[0003] However, when high-strength steel sheets with strengths ranging from 590 to 1470 MPa are used in automotive parts, stamping cracks are prone to occur due to decreased ductility. Furthermore, these steel sheets often utilize hard martensite in their microstructure to improve strength and ductility. Therefore, if laser welding is performed on these sheets, significant softening occurs in the heat-affected zone (HAZ) due to martensite softening. This results in problems such as fracture occurring in the softened HAZ during stamping or preferential fracture of the softened HAZ during part deformation, leading to reduced part strength. Therefore, for these high-strength steel sheets, excellent formability and minimal HAZ softening are desirable.

[0004] Against this backdrop, TRIP steel, which disperses retained austenite (retained γ) in the microstructure of steel plates, was developed as a technology to improve the ductility of steel plates.

[0005] For example, Patent Document 1 discloses that steel containing 0.10–0.45% C, 0.5–1.8% Si, and 0.5–3.0% Mn is subjected to aging treatment at 350–500°C for 1–30 minutes after annealing to generate residual γ, thereby obtaining a tensile strength (TS) of 80 kgf / mm². 2 The above and TS×El≥2500kgf / mm 2 •% of steel plates with high ductility.

[0006] Patent document 2 discloses that steel containing C: 0.10-0.25%, Si: 1.0-2.0%, and Mn: 1.5-3.0% is cooled to 450-300°C at a rate of 10°C / second or higher after annealing and held for 180-600 seconds. The residual austenite is controlled to be 5% or more, the bainitic ferrite to be 60% or more, and the polygonal ferrite to be 20% or less, based on the duty cycle, thereby obtaining a steel sheet with excellent ductility (E1) and tensile flange formability (λ).

[0007] Patent document 3 discloses a method of annealing a steel sheet with a specific composition, cooling it to a temperature range of 150–350°C, and then heating it to 350–600°C and holding it thereafter. This yields a microstructure containing ferrite, tempered martensite, and retained austenite, imparting high ductility and high tensile flange formability to the steel sheet. This utilizes the principle of so-called Q&P (Quenching & Partitioning): during cooling, the sheet is temporarily cooled to a temperature range between the martensitic transformation initiation temperature (Ms point) and the martensitic transformation completion temperature (Mf point), and then reheated and held to stabilize the retained γ. In recent years, the development of high-strength steels with high ductility and high tensile flange formability utilizing this principle has been ongoing.

[0008] Patent document 4 discloses an improved method for the above-mentioned Q&P treatment. That is, in steel with a specific composition, in order to make the polygonal ferrite content less than 5%, annealing is performed at a temperature of more than 10°C above the Ae3 point, and then cooling is stopped at a higher temperature of Ms-10°C to Ms-100°C, thereby generating upper bainite when reheated to 350 to 450°C, thus obtaining high ductility and high tensile flange formability.

[0009] Furthermore, Patent Document 5 discloses a method that effectively utilizes bainite formed at low temperatures and bainite formed at high temperatures to obtain a steel sheet with excellent ductility and low-temperature toughness. Specifically, steel containing 0.10–0.5% C is annealed and then cooled to 150–400°C at a cooling rate of 10°C / second or higher, held within this temperature range for 10–200 seconds to generate bainite in the low-temperature range. This is then reheated to a temperature range above 400°C but below 540°C and held for at least 50 seconds to generate bainite in the high-temperature range, thereby obtaining a steel sheet with excellent ductility and low-temperature toughness.

[0010] Furthermore, Patent Document 6 shows that a steel sheet containing C: 0.01-0.3%, Si: 0.005-2.5%, Mn: 0.01-3%, Mo: 0.01-0.3%, and Nb: 0.001-0.1% is annealed in a high-temperature range to achieve a near-γ single-phase structure, then cooled to a temperature range of 200-450°C and held thereafter. This results in a steel sheet containing 50-97% bainite or bainitic ferrite as the main phase and 3-50% austenite as the second phase, thereby obtaining a steel sheet with excellent ductility, porosity, and weldability.

[0011] Existing technical documents

[0012] Patent documents

[0013] Patent Document 1: Japanese Patent Publication No. 6-35619

[0014] Patent Document 2: Japanese Patent No. 4411221

[0015] Patent Document 3: Japanese Patent No. 5463685

[0016] Patent Document 4: Japanese Patent No. 5780086

[0017] Patent Document 5: Japanese Patent No. 5728115

[0018] Patent Document 6: Japanese Patent No. 3854506 Summary of the Invention

[0019] The problem that the invention aims to solve

[0020] However, although the conventional TRIP steel described in Patent Document 1 has excellent ductility, it has the problem of very low tensile flange formability.

[0021] In the technology described in Patent Document 2, the microstructure primarily utilizes bainitic ferrite, with relatively little ferrite suppression. Therefore, while the formability of the drawn flange is excellent, its ductility may not be high. Consequently, if its application to difficult-to-form parts is considered, further improvement in ductility is required.

[0022] The technology described in Patent Document 3 achieves higher ductility and high tensile flange formability compared to conventional TRIP steel and steel that effectively utilizes bainitic ferrite. However, fracture has been observed during the forming of difficult-to-form parts such as central pillars, indicating a need for further improvement in ductility. It is clear that the uniform deformation amount, which represents the ease of fracture, may not be sufficient for the steel sheet using this technology. This uniform deformation amount, denoted by U.E1, represents the elongation up to the point where necking begins to occur in E1, which is an indicator of ductility, and further increases in U.E1 are needed.

[0023] In the technology described in Patent Document 4, the amount of polygonal ferrite generated is reduced in order to reduce bulk martensite, but sufficient ductility cannot be ensured. In addition, the cooling stop temperature is set high in order to increase El, and a large amount of untransformed γ remains when cooling stops, so bulk martensite is prone to remain.

[0024] In the technology described in Patent Document 5, both low-temperature and high-temperature bainite phase transformations are effectively utilized to improve ductility. However, bainite that undergoes phase transformation at low temperatures contributes little to improving ductility, and when using bainite generated at high temperatures, blocky structures are easily left behind. Therefore, it is difficult to simultaneously impart high ductility and high tensile flange formability.

[0025] While the technology described in Patent Document 6 achieves improvements in raw materials with higher strength and is effective under conditions of high laser welding speed and low heat input, it exhibits significant softening in the HAZ (Heat Zone) under typical or low-speed welding conditions such as laser output power of 4–6 kW and welding speed of 3–5 mpm. Tensile tests on joints containing laser weld lines in a direction perpendicular to the tensile axis reveal the possibility of fracture in the HAZ. Furthermore, its ductility may not be high, and further improvements in ductility are desired.

[0026] Thus, in the existing technology, steel sheets with sufficiently high ductility and high tensile flange formability have not yet been obtained.

[0027] The present invention was made to solve such problems and provides steel plates, components, and methods for manufacturing the same, having tensile strength of 980 MPa or more, high ductility and excellent tensile flange formability, and excellent laser weldability.

[0028] It should be noted that the steel plates mentioned here also include galvanized steel plates whose surfaces have been galvanized.

[0029] In this invention, a tensile strength of 980 MPa or above refers to a tensile strength of 980 MPa or above according to JIS Z2241.

[0030] In addition, high ductility means that the total elongation T-El according to JIS Z2241 is 16.0% or more when TS is less than 1180MPa, 14.0% or more when TS is 1180MPa or more and less than 1320MPa, and 13.0% or more when TS is 1320MPa or more.

[0031] In addition, excellent stretch 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: initial hole diameter (mm), d: hole diameter when crack is generated (mm), and the hole expansion rate λ (%) = {(d-d0) / d0}×100 is 30% or more.

[0032] In addition, excellent laser weldability means that, after laser welding, fracture morphology determination test, and notch tensile test, the fracture morphology is that the base material fractures and the HAZ strength is ≥ base material TS + 50MPa.

[0033] (1) First, cut two steel plates with a diameter of 120 mm in the right-angle direction of rolling and 200 mm in the rolling direction (the end faces are ground) from the steel plate. Align the two plates together in the rolling direction and then perform laser welding at the joint. The gap between the joint surfaces is set to 0 mm.

[0034] (2) Laser welding was performed using an Nd-YAG laser, with the following settings: spot diameter at the focal point: 0.6 mm; focal point location: 4 mm above the top of the steel plate; shielding gas: Ar; laser output power: 4.2 kW; welding speed: 3.7 m / min. Tensile test pieces were cut from the welded component with the weld line perpendicular to the tensile axis and located at the center of the test piece's length (refer to...). Figure 1 (a) The fracture morphology is evaluated by performing a tensile test (fracture morphology determination test).

[0035] (3) When the fracture location is more than 2.0 mm away from the weld line (in some cases where it is more than 2.0 mm away), the fracture of the base material, the fracture of less than 2.0 mm away and the fracture along the weld line (the fracture is in the HAZ part and the molten part) are judged as the fracture of the weld.

[0036] (4) Furthermore, a notched test piece is cut from the welded component, with the weld line perpendicular to the tensile axis and located at the center of the test piece's length, and the welded portion has been notched (refer to...). Figure 1 (b) Perform a tensile test (notched tensile test).

[0037] (5) Thus, the strength of the HAZ itself is evaluated by forcibly breaking the HAZ by deforming only the HAZ and its surrounding small area of ​​the welded part.

[0038] Methods for solving problems

[0039] The inventors have conducted in-depth research on methods that possess high ductility, excellent tensile flange forming properties, and excellent laser welding properties, and have reached the following conclusions.

[0040] (i) After hot rolling and cold rolling of a steel billet with a specified composition, it is held at an annealing temperature of 810 to 900°C, thereby ensuring the specified residual austenite (residual γ) in the final microstructure, for which polygonal ferrite is controlled to be below 10%.

[0041] (ii) During the cooling process following the holding in the annealing process, the temperature is maintained within a range from 500°C to a holding temperature (T1) above the martensitic transformation initiation temperature Ms and above 320°C, with an average cooling rate (CR2) of less than 10°C / second for a holding time of 10 to 60 seconds. This generates bainitic ferrite with low carbide content and residual γ. The bainitic ferrite generated in this way is soft and less susceptible to the thermal effects of laser welding. Furthermore, during the holding time, carbon is efficiently distributed from the bainitic ferrite to the residual γ, thereby resulting in a high-carbon-concentration residual γ in the final microstructure that contributes to improved ductility.

[0042] (iii) In the remaining untransformed γ region, before the enrichment of carbon up to the T0 composition, which is the cause of the bulk structure, rapid cooling is performed at an average cooling rate (CR3): 3 to 100 °C / second, within a temperature range from the aforementioned stopping temperature (T1) to a cooling stopping temperature (T2) of 200 °C or higher and 300 °C or lower. Through this secondary cooling, the remaining untransformed γ region is segmented by martensitic or lower bainitic phase transformation, the residual γ is dispersed, and the bulk structure is reduced.

[0043] It should be noted that, here, "massive structure" refers to fresh martensite or retained austenite that appears as blocks on SEM. Additionally, T0 composition refers to the composition where the free energies of austenite and bainite are equal, and the bainitic phase transformation has ceased.

[0044] (iv) Then, the temperature is reheated at an average heating rate of 2°C / second or higher within a temperature range from the cooling stop temperature (T2) to 380°C, and held at a temperature range of 340°C to 590°C for 20 seconds to 3000 seconds. This transforms the martensite into tempered martensite, while simultaneously allowing carbon to further distribute into the residual γ formed adjacent to the upper bainite due to the holding time during cooling, thus stabilizing the residual γ. Furthermore, slow cooling is performed at an average cooling rate (CR4) of 0.01 to 5°C / second. This suppresses the formation of blocky structures caused by the distribution of excessive C to the residual γ, and improves the tensile flange properties by utilizing the self-tempering of fresh martensite.

[0045] By effectively utilizing the bainitic phase transformation before the martensitic phase transformation and then effectively utilizing the slow cooling after reheating, it is possible to achieve both the utilization of stable residual γ, which has been difficult to achieve until now, and the reduction of bulk structure.

[0046] The result is a steel sheet with extremely high ductility, excellent tensile flange formability, and excellent laser weldability. Furthermore, according to the present invention, high strength can also be achieved.

[0047] This invention is based on the above insights and specifically provides the following solutions.

[0048] [1] A steel plate having, by mass %, the following composition: C: 0.06-0.25%, Si: 0.4-2.5%, Mn: 1.5-3.5%, P: less than 0.02%, S: less than 0.01%, sol.Al: less than 1.0%, N: less than 0.015%, with the balance being Fe and unavoidable impurities.

[0049] And it has the following surface area ratios: polygonal ferrite: 10% or less (including 0%), tempered martensite: 40% or more, fresh martensite: 20% or less (including 0%), per 10 μm 2 The internal structure of the steel consists of 3-40% bainitic ferrite with fewer than 20 carbides and 5-20% retained austenite by volume fraction.

[0050] The area S of the region where the C concentration is above 0.50% C≥0.5 The area S of the region where the C concentration is above 0.30% C≥0.3 The proportion S in C≥0.5 / S C≥0.3 ×100 is over 20%.

[0051] [2] According to the steel plate described in [1], wherein, in the above-mentioned steel structure, with each 10 μm 2 The number density of retained austenite in the interior is less than 20 bainitic ferrite particles, with a density of retained austenite particles present in adjacent areas per 10,000 μm. 2 More than 50 in the middle.

[0052] [3] The steel plate according to [1] or [2], wherein the above composition, in mass %, further contains one or two selected from Ti: less than 0.1% and B: less than 0.01%.

[0053] [4] The steel plate according to any one of [1] to [3], wherein the above composition further contains, by mass %, one or more of the following: 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% and W: less than 0.2%.

[0054] [5] The steel plate according to any one of [1] to [4], wherein the above composition further contains, by mass %, one or more of the following: 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%.

[0055] [6] The steel plate according to any one of [1] to [5], wherein a zinc coating is provided on the surface.

[0056] [7] A component made of steel plate as described in any one of [1] to [6].

[0057] [8] A method for manufacturing a steel plate, wherein a steel billet having the composition described in [1], [3], [4] or [5] is hot-rolled and cold-rolled, and then the resulting cold-rolled steel plate is annealed.

[0058] The above annealing includes:

[0059] The process of holding the annealed food at an annealing temperature of 810–900°C;

[0060] A cooling process performed at an average cooling rate CR1: 5~100℃ / s within a temperature range of 810℃ to 500℃;

[0061] A process that takes place for 10 seconds to 60 seconds at an average cooling rate of CR2 of less than 10℃ / s within a temperature range from 500℃ to a dwell stop temperature T1 above the martensitic transformation start temperature Ms (℃) and above 320℃, and a dwell stop temperature of more than 320℃.

[0062] A process of cooling at an average cooling rate CR3 of 3 to 100°C / s within a temperature range from the aforementioned stopping temperature T1 to a cooling stopping temperature T2 of 200°C or more and 300°C or less.

[0063] A process of heating at an average heating rate of 2°C / s or more within a temperature range from the aforementioned cooling stop temperature T2 to 380°C;

[0064] 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 s to 3000 s; and

[0065] A process that cools to a temperature below 50°C at an average cooling rate CR5 of 0.1°C / s or higher.

[0066] [9] According to the steel plate manufacturing method described in [8], hot-dip galvanizing or alloying hot-dip galvanizing is performed in the process of holding at the above-mentioned average cooling rate CR4: 0.01~5℃ / s.

[0067]

[10] The steel plate manufacturing method according to [8] includes a step of electro-galvanizing after the cooling process at the above-mentioned average cooling rate CR5: 0.1℃ / s or higher.

[0068]

[11] 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 [6].

[0069] Invention Effects

[0070] According to the present invention, steel sheets and components with high ductility, excellent tensile flange formability, and excellent laser weldability 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, the automotive parts can be made lighter, and improved fuel efficiency can be expected. Attached Figure Description

[0071] Figure 1 This is a diagram illustrating the method for evaluating the laser weldability of the steel plate of the present invention.

[0072] Figure 2 This is an example of a SEM image of the steel structure in a steel plate.

[0073] Figure 3 This is a diagram illustrating the method for measuring the steel structure of the steel plate of the present invention.

[0074] Figure 4 This is a diagram illustrating the method for manufacturing the steel plate of the present invention. Detailed Implementation

[0075] The present invention will now be described in detail. It should be noted that the present invention is not limited to the following embodiments.

[0076] The steel plate of the present invention comprises, by mass percent, C: 0.06 to 0.25%, Si: 0.4 to 2.5%, Mn: 1.5 to 3.5%, P: less than 0.02%, S: less than 0.01%, sol.Al: less than 1.0%, N: less than 0.015%, with the balance being Fe and unavoidable impurities, and has, by area percent, polygonal ferrite: less than 10% (inclusive), tempered martensite: more than 40%, fresh martensite: less than 20% (inclusive), per 10 μm 2 The steel microstructure consists of 3-40% bainitic ferrite with fewer than 20 carbides and 5-20% retained austenite by volume fraction, with a carbon concentration of 0.50% or higher, covering an area S. C≥0.5 The area S of the region where the C concentration is above 0.30% C≥0.3 The proportion S in C≥0.5 / S C≥0.3 ×100 is over 20%.

[0077] The steel plate of the present invention will be described below in the order of composition and steel structure.

[0078] 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 %".

[0079] C: 0.06~0.25%

[0080] C is included from the viewpoints of ensuring the area ratio of tempered martensite to ensure the specified strength, ensuring the volume ratio of residual γ to improve ductility, and enriching in residual γ to stabilize it and thus improve ductility. Furthermore, by including C, the strength of the molten portion in the weld joint and the quenched portion from the γ region is increased, and deformation in the HAZ is suppressed, thereby improving HAZ softening resistance. 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, more preferably 0.11% or more. When the C content exceeds 0.25%, the upper bainitic phase transformation during intermediate holding in the cooling process is delayed, making it difficult to form residual γ adjacent to a specified amount of upper bainite. This results in reduced ductility. In addition, the increase in blocky martensite or blocky residual γ deteriorates the formability of the stretch flange. Furthermore, various properties such as HAZ softening resistance, spot weldability, bending properties, and hole expansion properties during laser welding of the steel sheet are significantly deteriorated. Therefore, the upper limit of C content is set at 0.25%. From the viewpoint of ductility and resistance to HAZ softening, the C content is preferably set to 0.22% or less. From the viewpoint of further improving ductility and resistance to HAZ softening, the C content is more preferably set to 0.20% or less.

[0081] Si: 0.4–2.5%

[0082] Si is contained from the viewpoints of increasing strength by strengthening ferrite, improving ductility by suppressing carbide formation in martensite and bainite to enhance the stability of residual γ, and improving the resistance to HAZ softening in welds by increasing the amount of solid solution strengthening that is less susceptible to heat effects. From these viewpoints, the Si content is set to 0.4% or more. From the viewpoint of improving ductility, the Si content is preferably set to 0.6% or more. More preferably, the Si content is 0.8% or more. When the Si content exceeds 2.5%, the rolling load during hot rolling becomes extremely high, making it difficult to manufacture thin sheets. Furthermore, the chemical conversion treatment properties and the toughness of the weld deteriorate. Therefore, the Si content is set to 2.5% or less. From the viewpoint of ensuring chemical conversion treatment properties, the toughness of the raw material, and the weld, the Si content is preferably set to less than 2.0%. From the viewpoint of ensuring the toughness of the weld, the Si content is preferably set to 1.8% or less, more preferably 1.5% or less.

[0083] Mn: 1.5–3.5%

[0084] From the perspective of ensuring strength by maintaining a specified area ratio of tempered martensite and / or bainite, improving ductility by stabilizing residual γ by lowering its Ms point, enhancing ductility by suppressing carbide formation in bainite similarly to Si, and improving ductility by increasing the volume fraction of residual γ, Mn is an important element. To achieve these effects, the Mn content is set to 1.5% or more. From the perspective of stabilizing residual γ to improve ductility, the Mn content is preferably set to 2.5% or more. The Mn content is preferably 2.6% or more, more preferably 2.7% or more. When the Mn content exceeds 3.5%, the bainitic phase transformation is significantly delayed, thus reducing ductility and resistance to HAZ softening. In addition, when the Mn content exceeds 3.5%, it is difficult to suppress the formation of blocky coarse γ and blocky coarse martensite, and the stretch flange formability also deteriorates. Therefore, the Mn content is set to 3.5% or less. From the viewpoint of promoting bainitic phase transformation and thus ensuring high ductility, the Mn content is preferably set to 3.2% or less. More preferably, the Mn content is set to 3.1% or less.

[0085] P: below 0.02%

[0086] Phosphorus (P) is an element that strengthens steel, but high levels of it deteriorate spot weldability. Therefore, the P content is set to 0.02% or less. From the viewpoint of improving spot weldability, the P content is preferably set to 0.01% or less. It should be noted that P can be omitted, but from the viewpoint of manufacturing cost, the P content is preferably set to 0.001% or more.

[0087] S: below 0.01%

[0088] Sulfur (S) is an element that improves the peeling properties of oxide scale during hot rolling and inhibits nitriding during annealing, but it reduces weldability, flexibility, and porosity. From these perspectives, the S content is set to 0.01% or less. In this invention, since the contents of C, Si, and Mn are high, weldability is easily reduced. From the viewpoint of improving weldability, the S content is preferably set to 0.0020% or less, more preferably less than 0.0010%. It should be noted that S may be absent, but from the viewpoint of manufacturing cost, the S content is preferably set to 0.0001% or more. More preferably, the S content is 0.0005% or more.

[0089] sol.Al: Less than 1.0%

[0090] Al is contained for deoxidation purposes and to stabilize residual γ-rays by replacing Si. There is no particular lower limit for sol.Al, but for stable deoxidation, the sol.Al content is preferably set to 0.005% or more. Furthermore, the sol.Al content is more preferably set to 0.01% or more. On the other hand, when the sol.Al content is 1.0% or more, the strength of the raw material is drastically reduced, and it also has an adverse effect on chemical conversion treatment; therefore, the sol.Al content is set to less than 1.0%. To obtain high strength, the sol.Al content is preferably set to less than 0.50%, and more preferably to 0.20% or less.

[0091] N: Less than 0.015%

[0092] 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 the effects of boron. When the N content is 0.015% or higher, the surface quality deteriorates significantly. Therefore, the N content is set to be less than 0.015%. The N content is preferably 0.010% or less. It should be noted that the steel may be free of N, but from a manufacturing cost perspective, the N content is preferably set to 0.0001% or more. More preferably, the N content is 0.001% or more.

[0093] The balance other than those mentioned above consists of Fe and unavoidable impurities. The steel sheet of the present invention preferably has a composition containing the above-mentioned basic components, with the balance consisting of iron (Fe) and unavoidable impurities.

[0094] The composition of the steel plate of the present invention may appropriately contain one or more of the following (A), (B), (C) as optional elements, based on the above-mentioned composition.

[0095] (A) Selected by mass% from one or both of Ti: less than 0.1% and B: less than 0.01%;

[0096] (B) Selected by mass% from one or more of the following: 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%, and W: less than 0.2%;

[0097] (C) Selected by mass% from one or more of 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%.

[0098] Ti: below 0.1%

[0099] Ti has the effect of fixing nitrogen in steel as TiN, thereby improving hot ductility, and also improves hardenability of boron. Additionally, it has the effect of refining the microstructure through the precipitation of TiC. To obtain these effects, it is preferable to set the Ti content to 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. On the other hand, when the Ti content exceeds 0.1%, it sometimes leads to an increase in rolling load and a decrease in ductility due to an increase in precipitation strengthening. Therefore, in the case of Ti content, the Ti content is set to 0.1% or less. Preferably, the Ti content is 0.05% or less. To ensure high ductility, the Ti content is more preferably set to 0.03% or less.

[0100] B: Below 0.01%

[0101] Boron (B) is an element that improves the hardenability of steel, readily forming tempered martensite and / or bainite with a specified area ratio. Furthermore, B improves hardenability near the weld, forming a hard structure near the weld and thus enhancing resistance to HAZ softening. Moreover, the residual dissolved B improves resistance to delayed fracture. To achieve this effect, the B content is preferably set to 0.0002% or more. More preferably, the B content is set to 0.0005% or more. Even more preferably, the B content is 0.0010% or more. On the other hand, when the B content exceeds 0.01%, not only does its effect saturate, but it also leads to a significant decrease in hot ductility and surface defects. Therefore, when B is present, the B content is set to 0.01% or less. Preferably, the B content is 0.0050% or less. More preferably, the B content is 0.0030% or less.

[0102] Cu: less than 1%

[0103] Cu improves the corrosion resistance of automobiles in the operating environment. Furthermore, the corrosion products of Cu coat the surface of the steel sheet, inhibiting hydrogen penetration. Cu is an element incorporated when effectively utilizing waste as a raw material; by allowing the inclusion of Cu, recycled materials can be effectively utilized as raw materials, reducing manufacturing costs. From this perspective, Cu is preferably contained at 0.005% or more, and more preferably at 0.05% or more from the viewpoint of improving resistance to delayed fracture. A Cu content of 0.10% or more is even more preferred. On the other hand, excessive Cu content leads to surface defects. Therefore, when Cu is present, the Cu content is set to 1% or less. A Cu content of 0.4% or less is preferred, and more preferably 0.2% or less is more preferred.

[0104] Ni: less than 1%

[0105] Like Cu, Ni also improves corrosion resistance. Furthermore, Ni can suppress the formation of surface defects that are prone to occur when Cu is present. Therefore, Ni is preferably present at 0.01% or more. More preferably, Ni content is 0.04% or more, and even more preferably 0.06% or more. On the other hand, excessive Ni content leads to uneven oxide scale formation in the furnace, which in turn causes surface defects. It also increases costs. Therefore, when Ni is present, the Ni content is set to 1% or less. Preferably, Ni content is 0.4% or less, and more preferably 0.2% or less.

[0106] Cr: less than 1.0%

[0107] Cr can be included to improve the hardenability of steel and suppress the formation of carbides in martensite and upper / lower bainite. Furthermore, by improving hardenability near the weld, a hard phase is formed near the weld, thus improving resistance to HAZ softening. To achieve this effect, the Cr content is preferably set to 0.01% or more. More preferably, the Cr content is 0.03% or more, and even more preferably 0.06% or more. On the other hand, excessive Cr content deteriorates pitting corrosion resistance. Therefore, when Cr is present, the Cr content is set to 1.0% or less. Preferably, the Cr content is 0.8% or less, and more preferably 0.4% or less.

[0108] Mo: 0.5% or less

[0109] Mo can be included to improve the hardenability of steel and to suppress the formation of carbides in martensite and upper / lower bainite. Furthermore, by improving hardenability near the weld, a hard phase is formed near the weld, thus improving resistance to HAZ softening. To achieve this effect, the Mo content is preferably set to 0.01% or more. More preferably, it is 0.03% or more, and even more preferably 0.06% or more. On the other hand, Mo significantly deteriorates the chemical conversion treatability of cold-rolled steel sheets. Therefore, when Mo is included, the Mo content is set to 0.5% or less. From the viewpoint of improving chemical conversion treatability, the Mo content is preferably set to 0.15% or less.

[0110] V: Below 0.5%

[0111] V can be included to improve the hardenability of steel, suppress carbide formation in martensite, upper bainite / lower bainite, refine the microstructure, and improve resistance to delayed fracture by precipitating carbides. Furthermore, by improving hardenability near the weld, a hard phase is formed near the weld, improving resistance to HAZ softening. To achieve these effects, the V content is preferably set to 0.003% or more. More preferably, it is 0.005% or more, and even more preferably 0.010% or more. On the other hand, when V is present in large quantities, castability deteriorates significantly. Therefore, when V is present, the V content is set to 0.5% or less. Preferably, the V content is 0.3% or less, more preferably 0.1% or less. Even more preferably, the V content is 0.05% or less, and even more preferably 0.03% or less.

[0112] Nb: below 0.1%

[0113] Nitrogen (Nb) can be included to improve the steel's microstructure and increase its strength, promote bainitic transformation through grain refinement, improve bending properties, and enhance resistance to delayed fracture. Furthermore, Nb improves hardenability near the weld, forming a hard phase near the weld and thus enhancing resistance to HAZ softening. To achieve these effects, the Nb content is preferably set to 0.002% or more. More preferably, it is 0.004% or more, and even more preferably 0.010% or more. On the other hand, when Nb is present in large quantities, precipitation strengthening becomes excessive, reducing ductility. It also leads to increased rolling load and deterioration of castability. Therefore, when Nb is present, the Nb content is set to 0.1% or less. Preferably, it is 0.05% or less, and more preferably 0.03% or less.

[0114] Zr: below 0.2%

[0115] 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 set to 0.005% or more. More preferably, it is 0.008% or more, and even more preferably 0.010% or more. On the other hand, when Zr is present in large quantities, the amount of coarse precipitates such as ZrN and ZrS remaining due to incomplete solution formation during the heating of the billet before hot rolling increases, leading to a deterioration in resistance to delayed fracture. Therefore, when Zr is present, the Zr content is set to 0.2% or less. Preferably, the Zr content is 0.15% or less, more preferably 0.08% or less. Even more preferably, the Zr content is 0.03% or less, and even more preferably 0.02% or less.

[0116] W: below 0.2%

[0117] W can be contained to improve the hardenability of steel, inhibit 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 set to 0.005% or more. More preferably, the W content is 0.008% or more, and even more preferably 0.010% or more.

[0118] On the other hand, when the W content is high, the amount of coarse precipitates such as WN and WS remaining due to lack of solid solution during the heating of the billet before hot rolling increases, resulting in a deterioration of the delayed fracture resistance. 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. The W content is even more preferably 0.03% or less, and even more preferably 0.02% or less.

[0119] Ca: below 0.0040%

[0120] Ca fixes S in the form of CaS, which helps improve flexural properties and resistance to delayed fracture. Therefore, the Ca content is preferably set to 0.0002% or more. The Ca content is more preferably 0.0005% or more, and even more preferably 0.0010% or more. On the other hand, when Ca is added in large quantities, surface quality and flexural properties deteriorate. 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.

[0121] Ce: below 0.0040%

[0122] Ce, like Ca, fixes sulfur, which helps improve flexural properties and resistance to delayed fracture. Therefore, the Ce content is preferably set to 0.0002% or more. More preferably, the Ce content is 0.0004% or more, and even more preferably 0.0006% or more. On the other hand, adding a large amount of Ce deteriorates surface quality and flexural properties. Therefore, when Ce is present, the Ce content is set to 0.0040% or less. Preferably, the Ce content is 0.0035% or less, and more preferably 0.0020% or less.

[0123] La: below 0.0040%

[0124] Like Ca, La also fixes S, which helps improve flexural properties and resistance to delayed fracture. Therefore, the La content is preferably set to 0.0002% or more. More preferably, the La content is 0.0004% or more, and even more preferably 0.0006% or more. On the other hand, adding a large amount of La deteriorates surface quality and flexural properties. 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.

[0125] Mg: less than 0.0030%

[0126] Mg fixes O in the form of MgO, which helps improve resistance to delayed fracture. Therefore, the Mg content is preferably set to 0.0002% or more. More preferably, the Mg content is 0.0004% or more, and even more preferably 0.0006% or more. On the other hand, adding a large amount of Mg deteriorates surface quality and flexibility. Therefore, when Mg is present, the Mg content is set to 0.0030% or less. Preferably, the Mg content is 0.0025% or less, and more preferably 0.0010% or less.

[0127] Sb: below 0.1%

[0128] Sb suppresses oxidation and nitriding of the steel plate surface layer, thereby suppressing the decrease in C and B content in the surface layer. Furthermore, by suppressing the aforementioned decrease in C and B content, ferrite formation in the steel plate surface layer is suppressed, improving resistance to delayed fracture while increasing strength. From this perspective, the Sb content is preferably set to 0.002% or more. More preferably, it is 0.004% or more, and even more preferably 0.006% or more. On the other hand, when the Sb content exceeds 0.1%, 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. Preferably, it is 0.04% or less, and more preferably 0.03% or less.

[0129] Sn: less than 0.1%

[0130] Sn inhibits oxidation and nitriding in the surface layer of the steel plate, thus suppressing the resulting decrease in the content of carbon (C) and boron (B) in the surface layer. Furthermore, by suppressing the aforementioned decrease in C and B content, ferrite formation in the surface layer of the steel plate is suppressed, improving resistance to delayed fracture while increasing strength. From this perspective, the Sn content is preferably set to 0.002% or more. The Sn content is preferably 0.004% or more, and more preferably 0.006% or more. On the other hand, when the Sn content exceeds 0.1%, castability deteriorates. Additionally, Sn segregates at the original γ grain boundaries, deteriorating 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. The Sn content is preferably 0.04% or less, and more preferably 0.03% or less.

[0131] When the optional components described above are contained in amounts below a suitable lower limit, the presence of optional elements below the lower limit does not impair the effects of the present invention. Therefore, when the optional components described above are contained in amounts below a suitable lower limit, the optional elements are contained as unavoidable impurities.

[0132] Next, the steel structure of the steel plate of the present invention will be described.

[0133] Polygonal ferrite: 10% or less (including 0%)

[0134] The polygonal ferrite formed during annealing or cooling helps improve ductility, but it also reduces the formability of the drawn flange due to the hardness difference between the polygonal ferrite and the surrounding hard phases such as martensite. Since the polygonal ferrite content is 10% or less in terms of area ratio, it does not impair the effects of the present invention and can therefore be included. Therefore, in the present invention, the polygonal ferrite content is set to 10% or less in terms of area ratio. Preferably, the polygonal ferrite content is 5% or less, more preferably 2% or less. Alternatively, the polygonal ferrite content can be 0%.

[0135] Tempered martensite: 40% or more

[0136] To obtain the specified strength and tensile flange formability, the tempered martensite content is set to 40% or more in terms of area ratio. Tempered martensite content is preferably 50% or more. On the other hand, when the tempered martensite content exceeds 80%, the ductility decreases due to excessive strength; therefore, the tempered martensite content is preferably set to 80% or less. Tempered martensite content is more preferably 75% or less.

[0137] Fresh martensite: less than 20% (including 0%)

[0138] Conventionally, when a large amount of bainitic transformation is desired during the final tempering process (the process described later where the steel is held at an average cooling rate of CR4), a large amount of blocky martensite or blocky residual γ remains. Therefore, conventionally, methods to promote bainitic transformation by reducing Mn content have been employed to prevent this problem. However, reducing the Mn content impairs ductility by losing the stabilizing effect of residual γ and increasing the volume fraction. In contrast, the present invention, which involves appropriate cooling treatment of steel sheets containing a large amount of Mn, achieves both the utilization of bainitic transformation and the reduction of blocky structures.

[0139] By reducing the bulky fresh martensite structure to below 20% by area ratio, excellent tensile flange formability and resistance to HAZ softening can be ensured. Therefore, in this invention, the fresh martensite content is set to below 20% by area ratio. To ensure excellent tensile flange formability and resistance to HAZ softening, the fresh martensite content is preferably set to below 10%. More preferably, it is below 5%. Alternatively, the fresh martensite content can be 0%.

[0140] per 10μm 2 The internal carbides consist of bainitic ferrite with fewer than 20 particles: 3-40%.

[0141] By containing more than 3% in area per 10μm 2 The internal carbides are bainitic ferrite with fewer than 20 carbon atoms, and carbon efficiently accumulates in the surrounding residual γ. Furthermore, it is less susceptible to the heat effects of laser welding, which helps improve resistance to HAZ softening. Therefore, in this invention, based on area ratio, per 10 μm 2 The bainitic ferrite with fewer than 20 carbides inside the core is set to 3% or more. Preferably, the bainitic ferrite content is 5% or more, more preferably 7% or more. To suppress the reduction in strength, per 10 μm... 2 The carbides inside are bainitic ferrite with 20 or fewer carbides, and the area ratio is set to 40% or less. Preferably, it is 30% or less, and more preferably 25% or less.

[0142] Additionally, in this invention, there may also be 10 μm 2 The interior contains more than 20 bainitic ferrite carbides.

[0143] The microstructure contains one or more of the following: tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite: 90% or more (including 100%).

[0144] To ensure the specified strength, ductility, and tensile flange formability, the total area ratio of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite, which constitutes the balance microstructure of the aforementioned polygonal ferrite, is preferably set to 90% or more. This balance microstructure can be one or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite, or it can be a microstructure composed of one or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite.

[0145] Upper and lower bainite contain per 10 μm 2 The internal carbides are bainitic ferrite with fewer than 20 segments. Additionally, the upper and lower bainite may also contain bainitic ferrite per 10 μm. 2 The interior contains more than 20 bainitic ferrite carbides.

[0146] Retained austenite: 5-20%

[0147] To ensure high ductility, the retained austenite (retained γ) is set to 5% or more by volume fraction relative to the overall steel structure. Retained austenite is preferably 7% or more, more preferably 9% or more. This amount of retained γ includes retained γ formed adjacent to bainite. Excessive increase in the amount of retained γ leads to decreased strength, reduced tensile flange formability, and deterioration of resistance to delayed fracture. Therefore, the volume fraction of retained γ is set to 20% or less. Retained austenite is preferably 15% or less. Furthermore, "volume fraction" can be considered as "area fraction".

[0148] The area S of the region where the C concentration is 0.50% (mass%) or higher C≥0.5 The area S of the region where the C concentration is above 0.30% (mass%) C≥0.3 The proportion S in C≥0.5 / S C≥0.3 ×100 is over 20%

[0149] To ensure high ductility, the area S of the region with a C concentration of 0.50% or higher is... C≥0.5 The area S of the region where the C concentration is above 0.30% C≥0.3 The proportion S in C≥0.5 / S C≥0.3 The percentage of ×100 is set to 20% or more. The above percentage is preferably 25% or more, and more preferably 30% or more.

[0150] With each 10μm 2 The number density of retained austenite in the interior is less than 20 bainitic ferrite particles, with a density of retained austenite particles present in adjacent areas per 10,000 μm. 2 More than 50 (preferred criteria)

[0151] Generate per 10μm 2 When the internal carbides are 20 or fewer bainitic ferrite, C efficiently distributes to the adjacent untransformed austenite, resulting in a high carbon concentration of residual γ in the final microstructure that contributes to improved ductility. In this invention, from the viewpoint of ensuring higher ductility, it is preferable to have a carbon concentration of γ per 10 μm. 2 The number density of retained austenite in the internal carbide layer, consisting of fewer than 20 bainitic ferrite particles, is set at 10,000 μm. 2 More than 50 per 10,000 μm. More preferably, the density of the above-mentioned number is 10000 μm. 2 The number of γ particles should be 70 or more, and more preferably 100 or more. Furthermore, from the viewpoint of avoiding intensity reduction due to excessive residual γ generation, the number of γ particles per 10 μm is [not specified]. 2 The number density of retained austenite in the interior, consisting of fewer than 20 bainitic ferrite carbides, is preferably 10,000 μm. 2 The number should be less than 400, and more preferably less than 300.

[0152] Next, the method for measuring the steel structure of the steel plate of the present invention will be described.

[0153] The area ratios of polygonal ferrite, bainitic ferrite, tempered martensite, and fresh martensite were determined by the following method: a plate thickness section perpendicular to the steel plate surface and parallel to the rolling direction was cut out, mirror polished, etched with 3% nitric acid ethanol solution, and 10 fields of view were observed at 5000x magnification at 1 / 4 thickness position using SEM. Figure 2 An example of a SEM image showing the steel microstructure of a steel plate. Figure 2 The polygonal ferrite shown has almost no carbides inside, serving as a comparison with equiaxed ferrite. It appears as the darkest area in SEM. Bainitic ferrite is a ferrite structure with internal formation of carbides or residual γ-carbides that appear white in SEM.

[0154] In cases where it is difficult to distinguish between bainitic ferrite and polygonal ferrite, regions of ferrite with an aspect ratio ≤ 2.0 are classified as polygonal ferrite, and regions with an aspect ratio > 2.0 are classified as bainitic ferrite, and the area ratio is calculated.

[0155] Figure 3 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 3 As shown in (A), the major axis ratio is calculated as follows: Find the longest major axis length *a* of the particle, and define the longest transverse section of the particle in the direction perpendicular to it as the minor axis length *b*. Use *a / b* as the major axis ratio. It should be noted that when particles are in contact with each other, as... Figure 3As shown in (B), the particles are divided at positions where they are roughly equally divided, and the size of each particle is measured.

[0156] Bainitic ferrite per 10 μm 2 The number of carbides within the bainite ferrite can be calculated as follows: Count the area of ​​the bainite ferrite and the number of carbides within it in a 5000x SEM image. Divide the number of carbides by the area of ​​each bainite ferrite to convert the result to per 10 μm. 2 The value in the middle is used to calculate the per 10 μm of bainitic ferrite. 2 The number of carbides inside the middle.

[0157] Tempered martensite is the region in SEM that contains lath-like lower structures and carbide precipitation. Fresh martensite is the whitish, blocky region in SEM that does not show lower structures.

[0158] The microstructure having one or more of the following components—tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite—corresponds to the remaining microstructure other than the aforementioned polygonal ferrite, and the total area ratio of this microstructure is the area ratio of the region other than the aforementioned polygonal ferrite. Here, the area ratio of carbides is very small, and therefore is included in the area ratio of the aforementioned remaining microstructure.

[0159] The volume fraction of retained austenite (retained γ) was determined by X-ray diffraction after chemically grinding the steel plate surface to a position of 1 / 4 thickness. An incident X-ray source of Co-Kα was used, and the volume fraction of retained austenite was calculated based on 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 determined by X-ray diffraction is equal to the area fraction of retained γ in the steel microstructure.

[0160] With each 10μm 2 The number density of retained austenite adjacent to bainitic ferrite with fewer than 20 carbides inside is calculated as follows: The sample used to observe the bainitic ferrite is mirror-polished. The electron backscattering diffraction pattern (EBSD) obtained from the same field of view in SEM is mapped and measured using the EBSD analysis program OIM Data Collection ver.7. The obtained data is analyzed using TSL OIMAnalysis ver.7 (manufactured by EDAX / TSL), thereby obtaining phase mapping data. The phase mapping data is measured and measured per 10 μm. 2The number density of FCC structures, consisting of fewer than 20 bainitic ferrite carbides within the interior, is used to calculate the number density per 10 μm. 2 The internal carbides consist of a residual austenite density of fewer than 20 bainitic ferrite particles in an adjacent manner. It should be noted that "adjacent" here refers to the density of retained austenite present in a phase mapping at approximately 10 μm intervals. 2 The internal carbides are 20 or fewer bainitic ferrites. The bcc structure is connected to the fcc structure, and also includes the case where the fcc structure is enclosed within the bcc structure.

[0161] The area S of the region where the C concentration is 0.50% (mass%) or higher C≥0.5 The area S of regions where the C concentration is 0.30% (mass%) or higher. C≥0.3 The measurements were performed as follows: At a section perpendicular to the steel plate surface and parallel to the rolling direction, at a position representing 1 / 4 of the plate thickness, a field emission electron probe microanalyzer (FE-EPMA) JXA-8500F manufactured by NEC was used, with an accelerating voltage of 6 kV and an irradiation current of 7 × 10⁻⁶ kV. -8 Under condition A, the beam diameter is set to the minimum, and the concentration distribution of C is mapped and analyzed, thereby making the determination.

[0162] However, to eliminate the influence of contamination, the background amount is subtracted by ensuring that the average carbon content obtained through analysis is equal to the carbon content of the parent material. That is, if the average carbon content measured is greater than the carbon content of the parent material, the increase is considered contamination, and the value obtained by subtracting this increase from the analytical values ​​at each location is taken as the true carbon content at each location.

[0163] The steel plate of the present invention preferably has a tensile strength of 980 MPa or more. More preferably, it has a tensile strength of 1180 MPa or more. Regarding the upper limit of the tensile strength, from the viewpoint of taking into account other properties, it is preferably 1450 MPa or less, and more preferably 1400 MPa or less.

[0164] Regarding the steel sheet of the present invention, by ensuring that the total elongation T-El is 16.0% or more when TS is 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, the forming stability is significantly improved. The hole expansion ratio λ is preferably ensured to be 30% or more. From the viewpoint of taking into account other properties, the upper limit of λ is preferably 90% or less at any strength level, and more preferably 80% or less.

[0165] For the steel plate of the present invention, laser welding, fracture morphology determination test, and notch tensile test are performed, and the preferred fracture morphology is base material fracture and HAZ strength ≥ base material TS + 50MPa.

[0166] The steel plate of the present invention described above can also be a steel plate with a zinc coating on its surface. The coating can be any one of hot-dip galvanizing or electroplating.

[0167] Next, the method for manufacturing the steel plate of the present invention will be described.

[0168] In the method for manufacturing the steel sheet of the present invention, a steel billet having the above-mentioned composition is hot-rolled and cold-rolled, and then the resulting cold-rolled steel sheet is annealed. The annealing process includes: a step of holding at an annealing temperature of 810 to 900°C; a step of cooling at an average cooling rate (CR1) of 5 to 100°C / second within a temperature range from 810°C to 500°C; a step of holding at an average cooling rate (CR2) of 10°C / second or less within a temperature range from 500°C to a stop temperature (T1) of 320°C or above the martensitic transformation start temperature Ms (°C); and a step of holding at an average cooling rate (CR2) of 10°C / second or less within a temperature range from 500°C to 320°C or above the stop temperature (T1); and a step of holding at an average cooling rate (CR2) of 10°C / second or less within a temperature range from the above-mentioned stop temperature (T1) to 60 seconds. The process of cooling at an average cooling rate (CR3) of 3 to 100°C / second within a temperature range from the cooling stop temperature (T1) to a cooling stop temperature (T2) of 200°C or higher and 300°C or lower; the process of heating at an average heating rate of 2°C / second or higher within a temperature range from the aforementioned cooling stop temperature (T2) to 380°C; the process of holding at an average cooling rate (CR4) of 0.01 to 5°C / second for a period of 20 seconds or more and 3000 seconds within a temperature range of 340°C or higher and 590°C or lower; and the process of cooling to a temperature of 50°C or lower with an average cooling rate (CR5) of 0.1°C / second or higher.

[0169] It should be noted that the specific temperature in each process of this invention refers to the surface temperature of the slab (steel billet) or steel plate.

[0170] in addition, Figure 4 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. Including this time-varying temperature, each step is described in detail below.

[0171] Hot rolling

[0172] For hot rolling of steel billets, there are methods such as rolling the billets after heating, rolling the continuously cast billets directly without heating, and rolling the continuously cast billets after a short-term heat treatment. Hot rolling can be carried out using conventional methods. For example, the billet heating temperature can be set to 1100–1300℃, the soaking time can be set to 20–300 minutes, the finishing rolling temperature can be set to the Ar3 phase transformation point to the Ar3 phase transformation point + 200℃, and the coiling temperature can be set to 400–720℃. From the viewpoint of suppressing plate thickness variation and ensuring stable high strength, the coiling temperature is preferably set to 430–530℃.

[0173] Cold rolling

[0174] In cold rolling, the rolling rate (cumulative rolling rate) can be set to 30-85%. From the viewpoint of consistently ensuring high strength and reducing anisotropy, the rolling rate is preferably set to 35-85%. It should be noted that under high rolling loads, softening annealing treatment can be carried out at 450-730°C using a CAL (continuous annealing line) or BAF (box annealing furnace).

[0175] annealing

[0176] 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 the desired heating and cooling rates, it is preferable to carry out the annealing using a continuous annealing line (CAL) or a continuous hot-dip galvanizing line (CGL).

[0177] Hold at annealing temperature: 810~900℃

[0178] To ensure a specified area ratio of tempered martensite and / or bainite and a specified volume ratio of residual γ, the annealing temperature is set to 810–900°C. To ensure that the polygonal ferrite content is 5% or less, the annealing temperature is preferably adjusted to anneal in the γ single-phase region. Preferably, it is 815°C or higher. On the other hand, when the annealing temperature exceeds 900°C, the γ grain size becomes too large, and the diffusion distance of C atoms required to obtain the desired carbon concentration of residual γ becomes longer, resulting in reduced ductility. Therefore, the annealing temperature is set to 900°C or lower. Preferably, the annealing temperature is 880°C or lower.

[0179] Cooling was performed at an average cooling rate (CR1) of 5–100 °C / second within a temperature range from 810 °C to 500 °C.

[0180] After holding at 810–900°C, cooling is performed within a temperature range from 810°C to 500°C at an average cooling rate (CR1) of 5–100°C / second. If the average cooling rate (CR1) is slower than 5°C / second, a large amount of ferrite is generated, leading to reduced strength and decreased tensile flange formability. The average cooling rate (CR1) is preferably 8°C / second or higher. On the other hand, if the average cooling rate (CR1) is too fast, the plate shape deteriorates; therefore, it is set to 100°C / second or lower. The average cooling rate (CR1) is preferably 50°C / second or lower, more preferably less than 30°C / second.

[0181] Here, the average cooling rate (CR1) is "(810℃ (cooling start temperature) - 500℃ (cooling stop temperature)) / (cooling time (seconds) from the cooling start temperature of 810℃ to the cooling stop temperature of 500℃))".

[0182] Within a temperature range from 500℃ to the dwell temperature (T1) above the martensitic transformation start temperature Ms (℃) and above 320℃, with an average cooling rate (CR2) of less than 10℃ / second, a dwell time of more than 10s and less than 60s.

[0183] By maintaining a residence time (slow cooling) of 10°C or less for 10 to 60 seconds within a temperature range from 500°C to a residence temperature (T1) above the martensitic transformation start temperature Ms (°C) and above 320°C, at an average cooling rate (CR2) of 10°C / second or less, bainite with low carbide density can be generated, and residual γ with high C concentration can be generated adjacently. When the temperature range is below Ms or below 320°C, martensite forms first, followed by lower bainite, thus leading to a decrease in strength. On the other hand, when the temperature exceeds 500°C, the driving force for the bainitic transformation decreases, and the bainite phase change decreases. Therefore, the temperature range is set to be above Ms, above 320°C, and below 500°C. This temperature range is preferably above 380°C, more preferably above 420°C. Furthermore, this temperature range is preferably below 480°C, more preferably below 460°C. When the average cooling rate (CR2) exceeds 10°C / second, the bainite phase change decreases. Therefore, the average cooling rate (CR2) is set to 10°C / second or less. When the residence time is less than 10 s, the desired amount of bainite is not obtained; when it exceeds 60 s, carbon enriches from bainite into the bulky, untransformed γ phase, leading to an increase in the amount of residual bulky structure. Therefore, the residence time is set to 10 s or more and 60 s or less. From the viewpoint of ensuring bainitic ferrite and retained austenite, and improving ductility and resistance to HAZ softening, a residence time of 20 s or more is preferred. Furthermore, from the viewpoint of improving the formability of the stretch flange by reducing bulky structure, a residence time of 50 s or less is preferred.

[0184] It should be noted that the martensitic phase transformation initiation temperature Ms can be determined by measuring the volume change when a cylindrical test piece (3 mm in diameter × 10 mm in height) is rapidly cooled with helium after being held at a specified annealing temperature using a phase transformation point determination tester.

[0185] Additionally, here, the average cooling rate (CR2) is defined as "(500°C (dwell start temperature) - dwell stop temperature (T1)) / (dwell time (seconds) from 500°C to dwell stop temperature (T1))".

[0186] Cooling is performed at an average cooling rate (CR3) of 3 to 100°C / second within a temperature range from the aforementioned stopping temperature (T1) to a cooling stopping temperature (T2) above 200°C and below 300°C.

[0187] After the aforementioned pause, rapid cooling is required to prevent excessive enrichment of carbon into the γ-ray structure. When the average cooling rate (CR3) in the temperature range from the aforementioned pause temperature T1 (above 320°C) to the cooling stop temperature T2 (above 200°C and below 300°C) is less than 3°C / second, carbon enriches into the bulky, untransformed γ-ray structure, increasing the amount of fresh martensite upon final cooling and reducing the formability of the stretched flange. Therefore, from the viewpoint of improving the formability of the stretched flange, the average cooling rate (CR3) in the temperature range from the aforementioned pause temperature T1 to the cooling stop temperature T2 (above 200°C and below 300°C) is set to 3°C / second or more. More preferably, the average cooling rate (CR3) is 5°C / second or more, and even more preferably 8°C / second or more. If the average cooling rate in this temperature range is too high, the plate shape deteriorates; therefore, the average cooling rate (CR3) in this temperature range is set to 100°C / second or less. The average cooling rate (CR3) is preferably 50°C / second or less. To ensure the specified amount of residual γ, the cooling stop temperature T2 is set to 200°C or higher. Preferably, the cooling stop temperature T2 is 220°C or higher, more preferably 240°C or higher. When the cooling stop temperature T2 exceeds 300°C, a large amount of blocky, untransformed γ remains, increasing the amount of fresh martensite during final cooling and reducing the formability of the stretched flange. Therefore, the cooling stop temperature T2 is set to 300°C or lower. Preferably, the cooling stop temperature T2 is 280°C or lower.

[0188] Here, the average cooling rate (CR3) is "(stalling temperature (T1)) - (cooling temperature (T2)) / (cooling time (seconds) from stalling temperature (T1) to cooling temperature (T2))".

[0189] Heating is carried out at an average heating rate of 2°C / second or higher within a temperature range from the cooling stop temperature (T2) to 380°C.

[0190] Furthermore, by heating for a short time within the temperature range from the aforementioned cooling stop temperature (T2) to 380°C, carbide precipitation can be suppressed, ensuring high ductility. Additionally, when the martensite or bainite generated during cooling is reheated to 380°C or higher using it as a core, upper bainite is generated. If the average heating rate up to 380°C is slow, these effects are not achieved. As a result, the residual γ content decreases, leading to reduced ductility. Therefore, the average heating rate within the temperature range from the cooling stop temperature (T2) to 380°C is set to 2°C / second or higher. From the viewpoint of suppressing carbide precipitation and generating upper bainite upon reheating, the average heating rate is preferably set to 5°C / second or higher, more preferably 10°C / second or higher. The upper limit of the aforementioned average heating rate is not particularly limited, but is preferably 50°C / second or lower, more preferably 30°C / second or lower.

[0191] Here, the average heating rate is defined as “380°C (heating stop temperature) - (cooling stop temperature (T2)) / (heating time (seconds) from cooling stop temperature T2 to 380°C (heating stop temperature))”.

[0192] Within a temperature range of 340℃ to 590℃, with an average cooling rate (CR4) of 0.01 to 5℃ / second, the residence time is more than 20 seconds and less than 3000 seconds.

[0193] From the perspective of stabilizing C by distributing it in the residual γ, and from the perspective of refining the region that is distributed in a blocky manner as untransformed γ using bainitic phase transformation to improve the formability of the stretch flange, a slow cooling period of 20s to 3000s is maintained within a temperature range of 340°C to 590°C. Furthermore, in order to suppress the formation of blocky structures caused by excessive C distribution into the residual γ, and to improve the formability of the stretch flange by utilizing the self-tempering of fresh martensite, slow cooling is performed within this temperature range at an average cooling rate (CR4) of 0.01 to 5°C / s. When the average cooling rate (CR4) is less than 0.01°C / s, excessive C is distributed into the residual γ, resulting in a decrease in the formability of the stretch flange due to the formation of blocky structures. Therefore, the average cooling rate (CR4) is set to 0.01°C / s or more. On the other hand, when the average cooling rate (CR4) exceeds 5°C / s, the distribution of C into the residual γ cannot be suppressed, and a sufficient amount of C-enriched regions cannot be obtained. Furthermore, the formation of fresh martensite leads to a deterioration of λ. Therefore, the average cooling rate (CR4) is set to be below 5°C / second.

[0194] Here, the average cooling rate (CR4) is "(cooling start temperature (T3)) - (cooling stop temperature (T4)) / (cooling time (seconds) from cooling start temperature (T3) to cooling stop temperature (T4))".

[0195] Here, the cooling start temperature (T3) and cooling stop temperature (T4) are not particularly limited as long as they are in the range of 340°C or above and 590°C. The cooling start temperature (T3) is preferably in the range of 360 to 580°C, and the cooling stop temperature (T4) is preferably in the range of 350 to 450°C.

[0196] It should be noted that holding (dwelling) within a temperature range of 340–590°C can be combined with hot-dip galvanizing. That is, in a process where the steel sheet is held at the aforementioned average cooling rate (CR4): 0.01–5°C / second, hot-dip galvanizing or alloying hot-dip galvanizing can be performed. When performing hot-dip galvanizing, it is preferable to immerse the steel sheet in a galvanizing bath at a temperature of 440°C or higher and 500°C, and then adjust the coating adhesion by methods such as gas wiping. A galvanizing bath with an Al content of 0.10% or higher and 0.22% or lower is preferably used for hot-dip galvanizing. Alternatively, as an alloying hot-dip galvanizing process, alloying of the zinc coating can be performed after hot-dip galvanizing. When performing alloying of the zinc coating, it is preferable to perform it within a temperature range of 470°C or higher and 590°C.

[0197] It should be noted that this process involves cooling (dwelling, slow cooling), but as long as the above-mentioned temperature range, dwell time range, and average cooling rate CR4 range are met, hot-dip galvanizing and zinc alloying can be performed in this process. Temperature increases may occur during hot-dip galvanizing and zinc alloying.

[0198] Cooling to below 50°C at an average cooling rate (CR5) of 0.1°C / second or higher.

[0199] Then, 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 / second 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 by a leveling machine. The average cooling rate (CR5) up to the temperature below 50°C is preferably 5°C / second or higher, more preferably 100°C / second or lower.

[0200] Here, the average cooling rate (CR5) is defined as "(340°C (cooling start temperature) - cooling stop temperature below 50°C) / (cooling time (seconds) from cooling start temperature to cooling stop temperature)".

[0201] From the viewpoint of improving the formability of the stretch flange, a low-temperature heat treatment can also be performed at 100–300°C for 30 seconds to 10 days after the aforementioned heat treatment or after surface rolling. This treatment removes hydrogen that has penetrated the steel sheet during tempering and annealing of the martensite formed during final cooling or surface rolling. Through this low-temperature heat treatment, the hydrogen content can be reduced to less than 0.1 ppm. Furthermore, electroplating can be performed. That is, the steel sheet can be electroplated with zinc after a cooling process at an average cooling rate (CR5) of 0.1°C / second or higher. After electroplating, from the viewpoint of reducing hydrogen in the steel, the aforementioned low-temperature heat treatment is preferred.

[0202] The thickness of the steel plate of the present invention is preferably set to 0.5 mm or more. Alternatively, the thickness is preferably set to 2.0 mm or less.

[0203] Next, the components of the present invention and their manufacturing method will be described.

[0204] 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 produce the component.

[0205] The steel sheet of the present invention has a tensile strength of 980 MPa or higher and exhibits excellent ductility, excellent tensile flange forming ability, and excellent laser weldability. Therefore, components obtained using the steel sheet of the present invention are also high-strength, possessing superior ductility, excellent tensile flange forming ability, and excellent laser weldability compared to conventional high-strength components. Furthermore, lightweighting can be achieved by using components of the present invention. Therefore, components of the present invention can be suitable for applications such as vehicle body frame parts. The components of the present invention also include welded joints.

[0206] 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.

[0207] Example

[0208] The embodiments of the present invention will be described below.

[0209] Cold-rolled steel sheets with a thickness of 1.4 mm and having the composition shown in Table 1 were processed under the annealing conditions shown in Table 2 to produce the steel sheets of the present invention and the comparative examples.

[0210] Each cold-rolled steel sheet is obtained by hot rolling (billet 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.

[0211] In Table 2, the martensitic phase transformation initiation temperature Ms is determined by measuring the volume change when a cylindrical test piece (3 mm in diameter × 10 mm in height) is rapidly cooled with helium after being held at a specified annealing temperature using a phase transformation point determination tester.

[0212] It should be noted that a portion of the steel sheet (cold-rolled steel sheet: CR) undergoes hot-dip galvanizing in a process where it is held at an average cooling rate of 0.01 to 5°C / second for 20 to 3000 seconds within a temperature range of 340°C to 590°C to produce hot-dip galvanized steel sheet (GI). Here, the steel sheet is immersed in a galvanizing bath at 440°C to 500°C for hot-dip galvanizing, and then the coating adhesion is adjusted by methods such as gas wiping. The galvanizing bath uses an Al content of 0.10% to 0.22%. Furthermore, a portion of the hot-dip galvanized steel sheet undergoes alloying treatment 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. Additionally, a portion of the steel sheet (cold-rolled steel sheet: CR) undergoes electroplating to produce electro-galvanized steel sheet (EG).

[0213] The steel microstructure was determined using the following method. The results are shown in Table 3.

[0214] The area ratios of polygonal ferrite, bainitic ferrite, tempered martensite, and fresh martensite were determined by the following method: a plate section parallel to the rolling direction was cut out, mirror-polished, etched with a 3% nitric acid ethanol solution, and 10 fields of view were observed at 1 / 4 thickness using SEM at a magnification of 5000. Figure 2 The polygonal ferrite shown has almost no carbides inside, serving as a comparison with equiaxed ferrite. It appears as the darkest area in SEM. Bainitic ferrite is a ferrite structure with internal formation of carbides or residual γ-carbides that appear white in SEM.

[0215] In cases where it is difficult to distinguish between bainitic ferrite and polygonal ferrite, regions of ferrite with an aspect ratio ≤ 2.0 are classified as polygonal ferrite, and regions with an aspect ratio > 2.0 are classified as bainitic ferrite, and the area ratio is calculated.

[0216] like Figure 3As shown in (A), the major axis ratio is calculated as follows: Find the longest major axis length *a* of the particle, and define the longest transverse section of the particle in the direction perpendicular to it as the minor axis length *b*. Use *a / b* as the major axis ratio. It should be noted that when particles are in contact with each other, as... Figure 3 As shown in (B), the particles are divided at positions where they are roughly equally divided, and the size of each particle is measured.

[0217] Bainitic ferrite per 10 μm 2 The number of carbides within the bainite ferrite was determined as follows: The area of ​​the bainite ferrite and the number of carbides within it were counted in a 5000x SEM image. The number of carbides was then divided by the area of ​​each bainite ferrite section to convert the result to a per 10 μm area. 2 The value in the middle is used to calculate the per 10 μm of bainitic ferrite. 2 The number of carbides inside the middle.

[0218] Tempered martensite is the region in SEM that contains lath-like lower structures and carbide precipitation. Fresh martensite is the whitish, blocky region in SEM that does not show lower structures.

[0219] The microstructure having one or more of the following components—tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite—corresponds to the remaining microstructure other than the aforementioned polygonal ferrite, and the total area ratio of this microstructure is the area ratio of the region other than the aforementioned polygonal ferrite. Here, the area ratio of carbides is very small, and therefore is included in the area ratio of the aforementioned remaining microstructure.

[0220] The volume fraction of retained austenite (retained γ) was determined by X-ray diffraction after chemically grinding the steel plate surface to a position of 1 / 4 thickness. An incident X-ray source of Co-Kα was used, and the volume fraction of retained austenite was calculated based on 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 determined by X-ray diffraction is equal to the area fraction of retained γ in the steel microstructure.

[0221] With each 10μm 2The number density of retained austenite adjacent to bainitic ferrite with fewer than 20 carbides inside is calculated as follows: The sample used to observe the bainitic ferrite is mirror-polished. The electron backscattering diffraction pattern (EBSD) obtained from the same field of view in SEM is mapped and measured using the EBSD analysis program OIM Data Collection ver.7. The obtained data is analyzed using TSL OIMAnalysis ver.7 (manufactured by EDAX / TSL), thereby obtaining phase mapping data. The phase mapping data is measured and measured per 10 μm. 2 The number density of FCC structures, consisting of fewer than 20 bainitic ferrite carbides within the interior, is used to calculate the number density per 10 μm. 2 The internal carbides consist of a residual austenite density of fewer than 20 bainitic ferrite particles in an adjacent manner. It should be noted that "adjacent" here refers to the density of retained austenite present in a phase mapping at approximately 10 μm intervals. 2 The internal carbides are 20 or fewer bainitic ferrites. The bcc structure is connected to the fcc structure, and also includes the case where the fcc structure is enclosed within the bcc structure.

[0222] The area S of the region where the C concentration is 0.50% (mass%) or higher C≥0.5 The area S of regions where the C concentration is 0.30% (mass%) or higher. C≥0.3 The measurements were performed as follows: At a section perpendicular to the steel plate surface and parallel to the rolling direction, at a position representing 1 / 4 of the plate thickness, a field emission electron probe microanalyzer (FE-EPMA) JXA-8500F manufactured by NEC was used, with an accelerating voltage of 6 kV and an irradiation current of 7 × 10⁻⁶ kV. -8 Under condition A, the beam diameter is set to the minimum, and the concentration distribution of C is mapped and analyzed, thereby making the determination.

[0223] However, to eliminate the influence of contamination, the background amount is subtracted by ensuring that the average carbon content obtained through analysis is equal to the carbon content of the parent material. That is, if the average carbon content measured is greater than the carbon content of the parent material, the increase is considered contamination, and the value obtained by subtracting this increase from the analytical value at each location is taken as the true carbon content at each location.

[0224] JIS 5 tensile test pieces and expanded hole test pieces were cut from the obtained steel plate and tensile tests were performed (according to JIS Z2241). The tensile strength TS and total elongation T-El are shown in Table 3. Specimens with a tensile strength of 980 MPa or more were judged as having excellent strength. In addition, specimens with a total elongation T-El of 16.0% or more when TS is 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 were judged as having excellent ductility.

[0225] Furthermore, the tensile flange formability was evaluated using a hole-expanding test piece cut from a heat-treated steel sheet, according to the Japanese Iron and Steel Federation standard JFST1001. Specifically, a 100mm x 100mm square sample was punched using a punching tool with a 10mm punch diameter and a 10.3mm die diameter (13% clearance). Then, a conical punch with a 60-degree apex angle was used to expand the hole, with the burrs generated during punching forming on the outer side, until a crack penetrating the plate thickness was formed. At this point, d0 (initial hole diameter, mm) and d (hole diameter at crack initiation, mm) were defined, and the expansion rate λ (%) was calculated in the form of {(d-d0) / d0} × 100, as shown in Table 3. Steel with λ of 30% or more was considered to have excellent hole-expanding properties.

[0226] Next, two steel plates, each 120mm in the right-angle direction of rolling and 200mm in the rolling direction (with ground ends), were cut from the obtained steel plate. These two plates were then butt-jointed and arranged with the rolling directions aligned, and laser welding was performed at the butt joint. The gap between the butt joint surfaces was set to 0mm. An Nd-YAG laser was used for laser welding, with the following settings: spot diameter at the focal point: 0.6mm; focal point location: 4mm above the top of the steel plate; shielding gas: Ar; laser output power: 4.2kW; welding speed: 3.7m / min. The weld line was cut from the welded component with the weld line perpendicular to the tensile axis and located at the center of the test piece's length. Figure 1 Tensile test specimens of the shape shown in (a) are used to evaluate fracture morphology (fracture morphology determination test) by performing tensile tests. Fractures occurring at a distance of 2.0 mm or more from the weld line (or at least 2.0 mm for some specimens) are classified as weld fractures. Fractures occurring at a distance less than 2.0 mm with cracks progressing along the weld line (cracks in the HAZ portion, progressing in the molten portion) are classified as weld fractures. Furthermore, a test specimen is cut from the welded component with the weld line perpendicular to the tensile axis and located at the center of the specimen's length, such as... Figure 1(b) A notched test specimen with a notched weld section is shown, and a tensile test (notched tensile test) is performed. This forces fracture within the HAZ (Hard Zone) of the weld section, causing only minor deformation of the HAZ and its surrounding area. This allows for evaluation of the HAZ's strength and provides a more quantitative assessment of the laser weldability. In this evaluation, specimens exhibiting fracture mode of base material fracture in the fracture mode determination test and HAZ strength ≥ base material TS + 50 MPa in the notched tensile test are considered to have excellent laser weldability (HAZ softening resistance).

[0227] The examples of the present invention shown in Tables 2 and 3 have excellent strength, ductility, hole expansion and laser weldability (HAZ softening resistance), while the comparative examples are inferior in any of these aspects.

[0228]

[0229] [Table 2]

[0230]

[0231] *The underlined part refers to something outside of this invention.

[0232] *1: Average cooling rate CR1 over the temperature range of 810℃~500℃

[0233] *2: Average cooling rate CR2 within the temperature range of 500℃ to the dwell temperature T1

[0234] *3: Residence time within the temperature range of 500℃ to the dwell time T1

[0235] *4: Average cooling rate CR3 over the temperature range from the dwell temperature T1 to the cooling stop temperature T2

[0236] *5: Average heating rate over the temperature range from cooling stop temperature T2 to 380°C

[0237] *6: Average cooling rate CR4 in the temperature range above 340℃ and below 590℃

[0238] *7: Duration of stay in a temperature range of 340℃ or higher and 590℃ or lower

[0239] *8: Average cooling rate CR5 up to temperatures below 50°C

[0240] *9: CR: Uncoated; GA: Alloyed hot-dip galvanized steel sheet; GI: Hot-dip galvanized steel sheet (alloyed treatment without zinc coating); EG: Electro-galvanized steel sheet

[0241] [Table 3]

[0242]

[0243] *The underlined part refers to something outside of this invention.

[0244] *10: per 10μm 2 The internal carbides are bainitic ferrite with fewer than 20 atoms.

[0245] *11: Per 10μm 2 The internal carbides consist of 20 or fewer bainitic ferrite particles and residual austenite present in adjacent layers.

[0246] *12: The microstructure contains one or more of the following: tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite.

[0247] This invention has extremely high ductility and excellent stretch flange formability, as well as excellent laser weldability, making it suitable for stamping applications in automobiles, home appliances, and other industries that undergo stamping processes.

[0248] Furthermore, it is understood that since the steel plate of the present invention is high-strength and has excellent ductility, excellent tensile flange forming ability and excellent laser weldability, the components obtained by forming, joining, and further forming and joining processes using the steel plate of the present invention are also high-strength and have excellent ductility, excellent tensile flange forming ability and excellent laser weldability, just like the steel plate of the present invention.

Claims

1. A steel plate having, by mass percent, the following composition: C: 0.06–0.25%, Si: 0.4–2.5%, Mn: 1.5–3.5%, P: less than 0.02%, S: less than 0.01%, sol.Al: less than 1.0%, N: less than 0.015%, with the balance being Fe and unavoidable impurities. And it has the following composition by area ratio: polygonal ferrite: 10% or less and including 0%, tempered martensite: 40% or more, fresh martensite: 20% or less and including 0%, per 10 μm 2 The internal carbides consist of 3-40% bainitic ferrite with fewer than 20 carbides, and 5-20% retained austenite by volume fraction. The area S of the region where the C concentration is above 0.50% C≥0.5 The area S in the region where the C concentration is above 0.30% C≥0.3 The proportion S in C≥0.5 / S C≥0.3 It is over 20%.

2. The steel plate according to claim 1, wherein, In the steel microstructure, with each 10 μm 2 The number density of retained austenite in the interior is less than 20 bainitic ferrite particles, with a density of retained austenite particles present in adjacent areas per 10,000 μm. 2 More than 50 in the middle.

3. The steel plate according to claim 1, wherein, The composition, by mass%, also contains one or two of the following: Ti: less than 0.1% and B: less than 0.01%.

4. The steel plate according to claim 2, wherein, The composition, by mass%, also contains one or two of the following: Ti: less than 0.1% and B: less than 0.01%.

5. The steel plate according to any one of claims 1 to 4, wherein, The composition, by mass%, also contains one or more of the following: 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%, and W: less than 0.2%.

6. The steel plate according to any one of claims 1 to 4, wherein, The composition, by mass%, also contains one or more of the following: 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%.

7. The steel plate according to claim 5, wherein, The composition, by mass%, also contains one or more of the following: 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%.

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

9. The steel plate according to claim 5, wherein, It has a zinc coating on the surface.

10. The steel plate according to claim 6, wherein, It has a zinc coating on the surface.

11. The steel plate according to claim 7, wherein, It has a zinc coating on the surface.

12. A component made of steel plate 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, 3 to 7 is hot-rolled and cold-rolled, and then the resulting cold-rolled steel sheet is annealed. The annealing includes: The process of holding the annealed food at an annealing temperature of 810–900°C; A cooling process performed at an average cooling rate CR1: 5~100℃ / s within a temperature range of 810℃ to 500℃; A process with an average cooling rate of CR2 of less than 10℃ / s and a dwell time of more than 10s and less than 60s within a temperature range from 500℃ to a dwell temperature T1 above the martensitic transformation start temperature Ms (℃) and above 320℃, and a dwell time of more than 10s and less than 60s. A cooling process performed at an average cooling rate CR3 of 3 to 100°C / s within a temperature range from the dwell temperature T1 to the cooling stop temperature T2, which is 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 T2 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 s to 3000 s; and A process that cools to a temperature below 50°C at an average cooling rate CR5 of 0.1°C / s or higher.

14. The method for manufacturing a steel plate according to claim 13, wherein, In the process of holding at the average cooling rate CR4: 0.01~5℃ / s, hot-dip galvanizing or alloyed hot-dip galvanizing is performed.

15. The method for manufacturing a steel plate according to claim 13, wherein, After the cooling process at an average cooling rate of CR5 of 0.1°C / s or higher, an electroplating zinc treatment process is included.

16. 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 11.

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