Steel plate and method for manufacturing the same

Abstract

This steel plate contains a predetermined composition with 10-60% in total of ferrite and bainite, 40-90% in total of martensite and tempered martensite, and 0-10% in total of pearlite and retained austenite. With respect to the total number of crystal grains of ferrite and bainite, the proportion of crystal grains of ferrite and bainite having an area of 3 µm2 or less is 40% or more, and the proportion of crystal grains of ferrite and bainite having an area of 30 µm2 or greater is 5% or less. A difference ΔMn between the Mn concentration in a position at 1.0 µm from an interface between ferrite and martensite in a vertical direction with respect to the interface and toward the inside of a ferrite grain and the maximum value of the Mn concentration in a region up to 0.5 µm is 1.00 mass% or less.

Description

[Technical Field]

[0001] The present invention relates to a steel sheet and Steel plate manufacturing method By law Regarding. This application claims priority based on Japanese Patent Application Nos. 2021-060949 and 2021-060950, filed on March 31, 2021, the contents of which are incorporated herein by reference. [Background technology]

[0002] In order to reduce carbon dioxide emissions from automobiles, attempts are being made to reduce the weight of automobile bodies while ensuring safety by using high-strength steel plates.

[0003] For example, Patent Document 1 describes a steel sheet excellent in elongation, hole expandability, bending workability, and delayed fracture resistance, which contains, by mass%, C: 0.15 to 0.25%, Si: 1.00 to 2.20%, Mn: 2.00 to 3.50%, P: 0.05% or less, S: 0.005% or less, Al: 0.01 to 0.50%, N: 0.010% or less, and B: 0.0003 to 0.0050%, and also contains one or more selected from Ti: 0.005 to 0.05%, Cu: 0.003 to 0.50%, Ni: 0.003 to 0.50%, Sn: 0.003 to 0.50%, Co: 0.003 to 0.05%, and Mo: 0.003 to 0.50%. The high-strength TRIP steel sheet has a composition consisting of 15% or less (including 0%) of ferrite with an average grain size of 2 μm or less, and the balance consisting of Fe and unavoidable impurities, and the microstructure consists of 15% or less (including 0%) of volume fraction ferrite with an average grain size of 2 μm or less, 2 to 15% of volume fraction retained austenite with an average grain size of 2 μm or less, 10% or less (including 0%) of volume fraction martensite with an average grain size of 3 μm or less, and the balance being bainite and tempered martensite with an average grain size of 6 μm or less, and the bainite and tempered martensite grains contain an average of 10 or more cementite particles with a grain size of 0.04 μm or more.

[0004] Patent Document 2 discloses a high-strength cold-rolled steel sheet having a tensile strength (TS) of 980 MPa or more and excellent bendability, the steel sheet having a specific chemical composition and a specific steel structure in which the area ratio of the ferrite phase is 30% to 70%, the area ratio of the martensite phase is 30% to 70%, the average grain size of the ferrite grains is 3.5 μm or less, the standard deviation of the grain size of the ferrite grains is 1.5 μm or less, the average aspect ratio of the ferrite grains is 1.8 or less, the average grain size of the martensite grains is 3.0 μm or less, and the average aspect ratio of the martensite grains is 2.5 or less, and the high-strength cold-rolled steel sheet has a tensile strength of 980 MPa.

[0005] Patent Document 3 discloses a high-strength steel sheet having a yield strength (YS) of 780 MPa or more, a tensile strength (TS) of 1180 MPa or more, and excellent spot weldability, ductility, and bending workability, the steel sheet having a C content of 0.15% or less, an area fraction of ferrite of 8 to 45%, an area fraction of martensite of 55 to 85%, the proportion of martensite adjacent to ferrite alone in the total structure of 15% or less, an average crystal grain size of ferrite and martensite of 10 μm or less, and an area fraction of ferrite with a crystal grain size of 10 μm or more among ferrite present in a depth range of 20 μm to 100 μm from the steel sheet surface of less than 5%.

[0006] Patent Document 4 describes a steel sheet with little variation in mechanical properties (particularly strength and ductility), which has a composition containing, by mass%, C: 0.10 to 0.25%, Si: 0.5 to 2.0%, Mn: 1.0 to 3.0%, P: 0.1% or less, S: 0.01% or less, Al: 0.01 to 0.05%, N: 0.01% or less, with the balance being Fe and unavoidable impurities, and has a structure containing ferrite as a soft first phase at an area ratio of 20 to 50%, with the balance being tempered martensite and / or tempered bainite as a hard second phase, in which the total area of ​​particles of the ferrite having an average particle diameter of 10 to 25 μm accounts for 80% or more of the total area of ​​all the ferrite particles, and the dispersion state of cementite particles present in all the ferrite particles with a circle equivalent diameter of 0.3 μm or more is such that the area ratio of the ferrite particles is 1 μm or more. 2Disclosed is a high-strength cold-rolled steel sheet having a density of more than 0.15 and not more than 1.0 per square inch and a tensile strength of 980 MPa or more. [Prior art documents] [Patent documents]

[0007] [Patent Document 1] International Publication No. 2017 / 179372 [Patent Document 2] International Publication No. 2016 / 194272 [Patent Document 3] Japanese Patent Publication No. 2015-117404 [Patent Document 4] Japanese Patent Application Publication No. 2013-245397 Summary of the Invention [Problem to be solved by the invention]

[0008] Steel sheets used for automobile parts and the like often contain punched holes, and the vicinity of the punched holes is subject to large local deformation during punching, resulting in the presence of many voids and other damage. Such void damage is likely to become the initiation point for fracture. In general, increasing the strength of a steel sheet makes the steel sheet more susceptible to fracture, but fracture is particularly likely to occur in areas that have been subjected to large deformation, such as the edges of punched holes and punched end faces. Therefore, it is difficult to achieve both strength and fracture resistance in steel sheets (especially high-strength steel sheets).

[0009] Various means have been proposed for achieving a balance between strength, formability, and bendability in high-strength steel sheets. However, although Patent Documents 1 to 4 mention increasing the strength and providing good ductility and bendability, they do not disclose any technology for ensuring the strength of steel members having punched end faces. In particular, conventional steel sheets (especially high-strength steel sheets) including those described in Patent Documents 1 to 4 have not been able to sufficiently improve fracture resistance.

[0010] An object of the present invention is to provide a steel sheet that can achieve high levels of strength, formability, and fracture resistance, a method for manufacturing the same, and a method for manufacturing an intermediate steel sheet. [Means for solving the problem]

[0011] The present inventors have discovered that the formability and fracture resistance of a steel sheet can be improved by forming the metallographic structure of the steel sheet to have ferrite, bainite, martensite, and tempered martensite, and optionally further having pearlite and retained austenite, and by controlling the ferrite and bainite to be fine. Furthermore, the inventors have discovered that controlling the Mn concentration near the interface between the ferrite and martensite delays the generation of voids near the edges of punched holes formed by punching, and makes it difficult for voids to join when a steel material containing punched holes is further deformed, thereby suppressing fracture.

[0012] The present invention has been made based on the above findings, and the gist of the present invention is as follows.

[0013] (1) A steel sheet according to one aspect of the present invention has a chemical composition, in mass%, C: 0.07~0.15%, Si: 0.01 to 2.0% Mn: 1.5 to 3.0%, P: 0~0.020%, S: 0 to 0.0200%, Al: 0.001 to 1.000%, N: 0~0.020%, Co: 0 to 0.500%, Ni: 0 to 1.000%, Mo: 0 to 1.000%, Cr: 0~2.000%, O: 0 to 0.0200%, Ti: 0 to 0.50% B: 0~0.0100%, Nb: 0 to 0.50% V: 0 to 0.500%, Cu: 0-0.5% W: 0 to 0.100%, Ta: 0 to 0.100%, Sn: 0 to 0.050% Sb: 0 to 0.050% As: 0~0.050%, Mg: 0 to 0.050% Ca: 0 to 0.050%, Zr: 0 to 0.050%, and REM: 0 to 0.100% and the balance being Fe and impurities, As the tissue fraction, The area ratio of ferrite and bainite is 10% or more and 60% or less in total, the area ratio of martensite and tempered martensite is 40% or more and 90% or less in total, and the area ratio of pearlite and retained austenite is 0% or more and 10% or less in total, The area of ​​the total number of crystal grains in the ferrite and bainite is 3 μm 2 The proportion of ferrite and bainite grains is 40% or more, Area is 30 μm 2 The proportion of ferrite and bainite grains that are equal to or greater than 5% is 5% or less, The difference ΔMn between the Mn concentration at a position 1.0 μm from the interface between the ferrite and the martensite in a direction perpendicular to the interface and toward the inside of the ferrite grain and the maximum Mn concentration in a region up to 0.5 μm is within 1.00 mass %. (2) The steel sheet according to (1) above has an area of ​​3 μm 2 The average aspect ratio of the crystal grains of ferrite and bainite may be 1.0 or more and 2.0 or less. (3) In the steel sheet according to (1) or (2) above, the average carbon concentration at a depth of 10 μm from the surface of the steel sheet in the sheet thickness direction may be 0.800 times or less of the average carbon concentration at a depth of 1 / 4 from the surface of the steel sheet in the sheet thickness direction. (4) The steel sheet according to any one of (1) to (3) above, wherein the composition is, in mass%, Co: 0.010~0.500%, Ni: 0.010 to 1.000%, Mo: 0.010 to 1.000%, Cr: 0.001 to 2.000%, O: 0.0001 to 0.0200%, Ti: 0.001 to 0.50% B: 0.0001~0.0100%, Nb: 0.001 to 0.50%, V: 0.001 to 0.500%, Cu: 0.001 to 0.5%, W: 0.001 to 0.100%, Ta: 0.001 to 0.100%, Sn: 0.001 to 0.050%, Sb: 0.001 to 0.050%, As: 0.001 to 0.050%, Mg: 0.0001 to 0.050%, Ca: 0.001 to 0.050%, Zr: 0.001 to 0.050%, and REM: 0.001 to 0.100% One or more of these may be contained. (5) A method for producing an intermediate steel plate according to one aspect of the present invention comprises, in mass%, C: 0.07~0.15%, Si: 0.01 to 2.0% Mn: 1.5 to 3.0%, P: 0~0.020%, S: 0 to 0.0200%, Al: 0.001 to 1.000%, N: 0~0.020%, Co: 0 to 0.500%, Ni: 0 to 1.000%, Mo: 0 to 1.000%, Cr: 0~2.000%, O: 0 to 0.0200%, Ti: 0 to 0.50% B: 0~0.0100%, Nb: 0 to 0.50% V: 0 to 0.500%, Cu: 0-0.5% W: 0 to 0.100%, Ta: 0 to 0.100%, Sn: 0 to 0.050% Sb: 0 to 0.050% As: 0~0.050%, Mg: 0 to 0.050% Ca: 0 to 0.050%, Zr: 0 to 0.050%, and REM: 0 to 0.100% a hot rolling process in which a slab having a composition containing Fe and the balance being Fe and impurities is hot rolled in a final finishing stand at a temperature range of 900°C or less and with a thickness reduction rate of 30% or more to obtain a hot rolled steel sheet; a coiling step of coiling the hot-rolled steel sheet at a coiling temperature of 650°C or less and 450°C or more after the hot rolling step; a holding step of holding the hot-rolled steel sheet after the coiling step in a temperature range from the coiling temperature to (the coiling temperature - 50) ° C. for a holding time of 8 hours or less; a cooling step of cooling the hot-rolled steel sheet after the holding step to 300°C at an average cooling rate of 0.10°C / second or more to obtain an intermediate steel sheet; It has. (6) A method for producing a steel sheet according to one aspect of the present invention includes a cold rolling step of cold-rolling an intermediate steel sheet produced by the method for producing an intermediate steel sheet according to (5) above at a thickness reduction rate of 20% or more and 80% or less to obtain a cold-rolled steel sheet; An annealing process in which the cold-rolled steel sheet is annealed by holding it in a temperature range of 740 ° C to 900 ° C for 60 seconds or more in an atmosphere having a dew point of -80 ° C or more and 20 ° C or less; It has. (7) In the method for producing a steel sheet according to (6) above, the dew point may be higher than -15°C and 20°C or lower. (8) The method for producing a steel sheet according to (6) or (7) above may include a coating layer forming step of forming a coating layer containing zinc, aluminum, magnesium or an alloy thereof on the front and back surfaces of the steel sheet in the annealing step. [Effects of the Invention]

[0014] According to the present invention, it is possible to provide a steel sheet that can achieve high levels of strength, formability, and fracture resistance, a method for manufacturing the same, and a method for manufacturing an intermediate steel sheet. Furthermore, according to the present invention, it is possible to provide a steel sheet that has formability suitable for use as a structural member of an automobile, etc., and has high tensile strength (e.g., 900 MPa or more). [Brief explanation of the drawings]

[0015] [Figure 1] FIG. 1 is a schematic diagram illustrating a method for measuring the Mn concentration near the interface between ferrite and martensite in this embodiment. DETAILED DESCRIPTION OF THE INVENTION

[0016] In one embodiment of the present invention, in a steel sheet (particularly in a high-strength steel sheet having a high tensile strength (for example, 900 MPa or more)), strength, formability, and fracture resistance are all achieved by controlling the grain sizes of ferrite and bainite and the Mn concentration near the interface between ferrite and martensite in addition to the area fraction of the metallographic structure. In a preferred aspect of this embodiment, the average carbon concentration in the surface layer of the steel sheet is further made lower than the average carbon concentration at a position ¼ depth from the surface of the steel sheet in the sheet thickness direction (specifically, a decarburized layer is formed in the surface layer of the steel sheet), thereby achieving a balance of strength, formability, bendability, and fracture resistance.

[0017] Hereinafter, a steel sheet according to one embodiment of the present invention will be described.

[0018] First, the metal structure of the steel sheet according to this embodiment will be described. Hereinafter, the structure fraction will be expressed as an area fraction, and therefore the unit "%" of the structure fraction means area %.

[0019] <Metal structure> (Total area ratio of ferrite and bainite: 10% or more, 60% or less) Ferrite and bainite are soft structures that are easily deformed and contribute to improving elongation. When the total area ratio of ferrite and bainite is 10% or more, sufficient elongation can be obtained, contributing to improving formability. The total area ratio of ferrite and bainite is preferably 20% or more, and more preferably 25% or more.

[0020] To ensure tensile strength, the total content of ferrite and bainite is set to 60% or less, preferably 50% or less, and more preferably 45% or less.

[0021] (Total area ratio of martensite and tempered martensite: 40% or more, 90% or less) Martensite and tempered martensite are hard structures, which contribute to improving tensile strength. By making the total of martensite and tempered martensite 40% or more, high strength can be achieved, making it easier to ensure a tensile strength of, for example, 900 MPa or more. The total is preferably 45% or more, and more preferably 50% or more.

[0022] Furthermore, if the total of martensite and tempered martensite exceeds 90%, sufficient elongation cannot be obtained and formability deteriorates, so the total is set to 90% or less, preferably 80% or less, and more preferably 75% or less.

[0023] (Total area ratio of pearlite and retained austenite: 0% or more, 10% or less) Pearlite is a structure containing hard cementite, which becomes the origin of voids after punching, thereby degrading fracture resistance. Furthermore, retained austenite is a structure that contributes to improving elongation through transformation-induced plasticity (TRIP). However, martensite, which is generated by the transformation of retained austenite, is very hard and becomes the origin of voids, thereby degrading fracture resistance. Therefore, the total area fraction of pearlite and retained austenite is set to 10% or less, preferably 5% or less. In this embodiment, pearlite and retained austenite do not necessarily have to be generated, and the total area fraction of pearlite and retained austenite may be 0%.

[0024] (The total number of grains in ferrite and bainite, N T In contrast, the area is 3 μm 2 The proportion of ferrite and bainite grains N3 is 40% or more. The total number of grains in ferrite and bainite, N T For an area of ​​3m 2 The ratio of the number N3 of ferrite and bainite grains (N3 / N T ) is an index representing the formability and fracture resistance of the steel sheet in this embodiment. 2 The ratio of the number of crystal grains (fine grains) below (N3 / N T ) is 40% or more, voids are less likely to form around the punched hole and its edge during punching or deformation after punching. In addition, the formed voids are less likely to connect with each other, leading to fracture. Therefore, in ferrite and bainite, 2 The ratio of the number of crystal grains below (N3 / N T ) to 40% or more, the fracture resistance is higher than that of automotive parts with the same level of strength. 2 The ratio of the number of crystal grains below (N3 / N TIf the ratio is less than 40%, it is difficult to obtain a sufficient effect of improving the fracture resistance. 2 The ratio of the number of crystal grains below (N3 / N T The lower limit of the content of ferrite and bainite is 40% or more, preferably 50% or more, and more preferably 55% or more. 2 There is no particular upper limit set for the proportion of the number of the following crystal grains, and it may be 100%, but from the viewpoint of suppressing the yield point elongation, it may be 90% or less.

[0025] (The total number of grains in ferrite and bainite, N T In contrast, the area is 30 μm 2 The number of ferrite and bainite grains N is greater than or equal to 30 (The percentage of The total number of grains in ferrite and bainite, N T In contrast, the area is 30 μm 2 The number of ferrite and bainite grains N is greater than or equal to 30 Proportion (N 30 / N T ) is an index representing the formability and fracture resistance of the steel sheet in this embodiment. 2 If the proportion of these coarse crystal grains is high, voids are likely to form around the punched hole and its edge during punching or deformation during punching, and the formed voids are likely to connect with each other, leading to fracture. 2 The ratio of the number of crystal grains above (N 30 / N T ) exceeds 5%, the fracture resistance is deteriorated compared to automotive parts with the same level of strength. 2 The ratio of the number of crystal grains above (N 30 / N T ) is set to 5% or less, preferably 3% or less. 2 The proportion of crystal grains above (N 30 / N T) is preferably as small as possible, so there is no particular lower limit and it may be 0%, but from the viewpoint of suppressing increases in manufacturing costs due to precise control, it may be 1% or more.

[0026] (The difference ΔMn between the Mn concentration at a position 1.0 μm from the interface between ferrite and martensite in a direction perpendicular to the interface and toward the inside of the ferrite grain, and the maximum Mn concentration in the region up to 0.5 μm is within 1.00 mass%.) The Mn concentration [Mn 1.0 ] and the maximum Mn concentration in the region up to 0.5 μm [Mn 0.5 ] and the difference ΔMn([Mn 0.5 ]-[Mn 1.0 ]) is an index that indicates the ease of void formation at the interface. Generally, when a composite structure containing ferrite and martensite is deformed, the soft ferrite deforms first. During this process, dislocations generated within the ferrite pile up at the interface, causing stress concentration at the interface. Examples of interfaces where dislocations pile up include ferrite-ferrite interfaces and ferrite-martensite interfaces. However, voids are more likely to form at the ferrite-martensite interface due to the large difference in hardness between them. Mn is an element that increases the ductile-brittle transition temperature and makes fracture more likely (reference: Tanaka et al., Iron and Steel, Vol. 100 (2014) No. 10). Therefore, if the Mn concentration is high near the interface where voids are likely to form, i.e., the ferrite-martensite interface, void formation near the interface is promoted during punching and deformation after punching. The Mn concentration at a position 1.0 μm perpendicular to the interface and toward the inside of the ferrite grain [Mn 1.0 ] corresponds to the average Mn concentration in the ferrite grains, while the maximum Mn concentration in the region up to 0.5 μm [Mn 0.5] represents the Mn concentration in the region near the interface where voids are likely to occur. Therefore, the difference ΔMn between these is an index representing the Mn concentration at the interface between ferrite and martensite. If ΔMn exceeds 1.00 mass%, void generation at the interface is significantly promoted, significantly deteriorating the fracture resistance, so ΔMn is set to 1.00 mass% or less, preferably 0.50 mass% or less. There is no particular lower limit for ΔMn, and it may be 0 mass% or more, but it may also be 0.01 mass% or more.

[0027] This distribution of Mn concentration near the interface between ferrite and martensite is formed during continuous annealing. During heating during continuous annealing, Mn is enriched at the prior austenite grain boundaries or at the ferrite / austenite interface. During the cooling process, ferrite forms from austenite. If a certain area ratio of ferrite is formed throughout the steel sheet, the amount of interfacial movement of a single grain is smaller the smaller the average grain size during heating. In other words, by reducing the average grain size before continuous annealing, the amount of interfacial movement per grain during the cooling process can be reduced, thereby shortening the distance between the enriched region at the interface before cooling and the interface after ferrite formation. In general, diffusion is faster at grain boundaries than within grains, so the closer the Mn-enriched region is to the grain boundary, the easier it is for Mn to diffuse by grain boundary diffusion. The inventors have found that by controlling the amount of interface movement during cooling in continuous annealing in this way, it is possible to diffuse the Mn-enriched areas during the cooling process and reduce the Mn concentration at the interface between ferrite and martensite, i.e., the above-mentioned ΔMn.

[0028] (area is 3 μm 2 The average aspect ratio of the ferrite and bainite grains is 1.0 or more and 2.0 or less. Area is 3 μm 2The average aspect ratio of the ferrite and bainite grains, which is equal to or less than 1.0, is an index of fracture resistance. Generally, the smaller the aspect ratio and the more equiaxed the grains, the less likely stress concentration occurs at the interface. In order to exhibit sufficient fracture resistance, it is preferable that the average aspect ratio of the ferrite and bainite grains be 1.0 or more and 2.0 or less. An average aspect ratio of 2.0 or less makes it easier to achieve this effect, and a ratio of 1.0 or more and 1.5 or less is more preferable.

[0029] In this embodiment, the aspect ratio refers to the ratio of the longest diameter (major diameter) of a ferrite grain to the longest diameter (minor diameter) of the ferrite grains perpendicular to the longest diameter (major diameter). The same applies to the aspect ratio of bainite grains.

[0030] In addition, the area is 30 μm 2 The average aspect ratio of the ferrite and bainite crystal grains is not particularly limited, but elongated grains are preferable from the viewpoint of reducing stress concentration at the interface, so the average aspect ratio may be more than 2.0 and not more than 5.0.

[0031] (The average carbon concentration at a depth of 10 μm from the steel plate surface in the plate thickness direction is 0.800 times or less of the average carbon concentration at a depth of 1 / 4 from the steel plate surface in the plate thickness direction) In this embodiment, a decarburized layer may be formed on the surface of the steel sheet. The decarburized layer formed on the surface of the steel sheet is an index representing bendability. By setting the average carbon concentration of this decarburized layer at a depth of 10 μm in the sheet thickness direction to 0.800 or less times the average carbon concentration inside the steel sheet that is not affected by decarburization, i.e., at a position 1 / 4 depth in the sheet thickness direction from the surface of the steel sheet, the bendability of the steel sheet can be improved. The average carbon concentration at a depth of 10 μm from the surface of the steel sheet being 0.800 or less times the average carbon concentration at the position 1 / 4 depth means that sufficient decarburization has occurred. Sufficient decarburization can fully enjoy the effect of improving bending properties. Therefore, the average carbon concentration at a depth of 10 μm from the surface of the steel sheet is set to 0.800 or less, preferably 0.600 or less, and more preferably 0.400 or less, times the average carbon concentration at the position 1 / 4 depth. Although there is no particular lower limit, the average carbon concentration at a depth of 10 μm from the steel sheet surface is preferably 0.001 times or more, more preferably 0.005 times or more, the average carbon concentration at a depth of 1 / 4.

[0032] When a steel sheet has a coating layer (e.g., a plating layer) on its surface, the "surface" at a depth of 10 μm from the steel sheet surface in the sheet thickness direction refers to the surface of the base steel. The reason why the average carbon concentration at a depth of 10 μm from the steel sheet surface in the sheet thickness direction is used as the standard is because the carbon concentration at this depth greatly contributes to bendability.

[0033] The average carbon concentration at each location can be measured by glow discharge optical emission spectroscopy (GDS). The concentration profile of each element is measured using GDS from the surface of the steel plate in the depth direction (plate thickness direction), and the average carbon concentration at a position 10 μm from the steel plate surface is determined. The average carbon concentration at a 1 / 4 depth position is determined by grinding the plate down to 1 / 4 of the thickness and then measuring the ground surface with GDS.

[0034] Next, the identification of ferrite, bainite, martensite, tempered martensite, pearlite, and retained austenite and the calculation of their areas and area ratios will be described.

[0035] The identification, area, and area ratio of each metal structure can be calculated by EBSD (Electron Backscattering Diffraction), X-ray measurement, corrosion using Nital reagent or Lepera solution, and by observing a 100 μm × 100 μm region of the steel sheet cross section parallel to the rolling direction and perpendicular to the sheet surface at a magnification of 1,000 to 50,000 times using a scanning electron microscope. When measuring the area ratio of any structure, measurement is made at three locations, and the average value is calculated.

[0036] The area and area fraction of ferrite can be measured using the following method. Specifically, using an EBSD attached to a scanning electron microscope, measurements are taken at 0.2 μm intervals (pitch) over a range of 1 / 8 to 3 / 8 of the thickness, centered at a position 1 / 4 of the way down from the surface of the steel sheet. The value of the grain average misorientation (GAM) is calculated from the measurement data. The area and area fraction of ferrite are then measured for regions with an average local misorientation value of less than 0.5°. Here, the average local misorientation is the misorientation between adjacent measurement points in a region surrounded by grain boundaries with a crystal misorientation of 5° or more, calculated, and averaged over all measurement points within the grain.

[0037] The area and area fraction of bainite are measured by taking a sample from a cross-section of the steel sheet parallel to the rolling direction, polishing the sample, etching it with nital, and observing the area from 1 / 8 to 3 / 8 of the thickness centered at 1 / 4 of the sheet thickness using a field emission scanning electron microscope (FE-SEM). The area fraction can be calculated using known image analysis software, such as "ImageJ." ImageJ is open-source, public domain image processing software widely used by those skilled in the art. In FE-SEM observations, the structure of an observation surface, for example, a square with sides of 30 μm, is classified as follows: Bainite is a collection of lath-shaped crystal grains that either does not contain iron-based carbides with a major axis of 20 nm or more, or contains iron-based carbides with a major axis of 20 nm or more, and the carbides belong to a single variant, i.e., a group of iron-based carbides elongated in the same direction. Here, a group of iron-based carbides elongated in the same direction refers to iron-based carbides whose elongation directions differ by within 5°. Bainite grains surrounded by grain boundaries with a misorientation of 15° or more are counted as one bainite grain.

[0038] The area fractions of martensite and tempered martensite can be calculated by etching with a Repeller solution, observing and photographing the area of ​​1 / 8 to 3 / 8 of the plate thickness, centered at 1 / 4, using an FE-SEM, and subtracting the area fraction of retained austenite measured using X-rays (details will be described later) from the area fraction of the uncorroded area.

[0039] The area fraction of retained austenite can be calculated by measuring the diffraction intensity using X-rays on a sample in which a 100 μm region in the thickness direction from the surface has been removed by electrolytic polishing or chemical polishing. Specifically, measurement is performed using MoKα radiation as characteristic X-rays, and the area fraction of retained austenite can be calculated from the integrated intensity ratio of the diffraction peaks of (200), (211) of the bcc phase and (200), (220), and (311) of the fcc phase.

[0040] The pearlite area ratio can be determined by taking a sample from a cross-section of the steel plate parallel to the rolling direction, polishing the observation surface, etching it with Nital reagent, and observing and photographing a secondary electron image of the area from 1 / 8 to 3 / 8 of the thickness, centered at 1 / 4 of the thickness from the surface of the steel plate, using a scanning electron microscope. In the secondary electron image, carbides are observed with a relatively brighter contrast than other steel structures. The area in the photographed image where plate-like carbides are arranged in rows with intervals of 0.5 μm or less is defined as pearlite, and the pearlite area ratio is calculated using the image analysis software "ImageJ" mentioned above.

[0041] The ratio of the grain size and number of ferrite and bainite crystal grains, i.e., the area of ​​3 μm 2 The number of ferrite and bainite grains with an area of ​​30 μm or less 2 The number of ferrite and bainite crystal grains and the ratio of these numbers to the total number of crystal grains in ferrite and bainite are calculated by image analysis using the above-mentioned "ImageJ." Specifically, a 100 μm × 100 μm area of ​​the steel sheet cross section parallel to the rolling direction and perpendicular to the sheet surface is observed and photographed with a scanning electron microscope at a magnification of 1,000 to 50,000 times, and the ratio is calculated using "ImageJ." 2 The number of ferrite and bainite grains that are less than 0.1 μm 2 The number of grains smaller than 0.1 μm is not included. 2 (Grains smaller than this size are excluded as noise.)

[0042] The difference ΔMn between the Mn concentration at a position 1.0 μm from the ferrite-martensite interface, perpendicular to the interface and toward the inside of the ferrite grain, and the maximum Mn concentration in the region up to 0.5 μm is calculated using an EPMA (Electron Probe Micro Analyzer). As shown in Figure 1, line analysis is performed with the EPMA in a direction perpendicular to the tangent to the ferrite-martensite interface and toward the inside of the ferrite grain. The step interval is 0.01 μm. The difference ΔMn is calculated as the difference between the Mn concentration at a position 1.0 μm from the interface and the maximum Mn concentration from the interface to a position 0.5 μm. Five ferrite-martensite interfaces located near 1 / 4 of the sheet thickness (any position between 1 / 8 and 3 / 8 of the sheet thickness depth) are selected arbitrarily, and ΔMn is calculated for each, and the average value is evaluated. Note that the measurement targets are either straight or curved interfaces to which a tangent can be drawn. Here, the "interface between ferrite and martensite" in this embodiment is defined as follows. First, a sample is taken from the cross section of the steel plate parallel to the rolling direction as the observation surface, and the carbon concentration is analyzed using an EPMA in a 100 μm x 100 μm area of ​​the observation surface in the range of 1 / 8 to 3 / 8 of the thickness, centered at 1 / 4 of the plate thickness from the surface of the steel plate. Then, based on the obtained carbon concentration mapping, the structure distribution is identified by classifying structures with a carbon concentration of less than 0.1 mass% as ferrite, structures with a carbon concentration of 0.1 to 0.8 mass% as martensite, and structures with a carbon concentration of more than 0.8 mass% as retained austenite, thereby defining the "interface between ferrite and martensite."

[0043] The average aspect ratios of ferrite and bainite grains are calculated by image analysis using the aforementioned "ImageJ."

[0044] Next, the reasons for limiting the chemical composition of the steel sheet according to this embodiment will be explained. Hereinafter, % in the chemical composition means % by mass.

[0045] (C: 0.07% or more, 0.15% or less) C is an element that ensures a predetermined amount of martensite and improves the strength of the steel sheet. If the C content is less than 0.07%, it is difficult to obtain the predetermined amount of martensite, and high tensile strength (for example, 900 MPa or more) cannot be ensured, so the C content is set to 0.07% or more. The C content is preferably 0.09% or more. On the other hand, if the C content exceeds 0.15%, the formation of ferrite is suppressed, which leads to a decrease in elongation and a deterioration in ductility of the punched end surface, so the C content is set to 0.15% or less. From the viewpoint of suppressing deterioration in bendability, a small C content is preferable. The C content is preferably 0.13% or less.

[0046] (Si: 0.01% or more, 2.0% or less) Si not only increases strength as a solid-solution strengthening element, but is also effective in obtaining a microstructure containing martensite, bainite, and even residual γ. The Si content is adjusted according to the target strength level. A content exceeding 2.0% results in poor press formability, reduced phosphatability, and poor ductility at punched edges. A low Si content is also preferable from the perspective of suppressing deterioration in bendability. Therefore, the upper limit is set to 2.0% or less. When hot-dip galvanizing is applied, problems such as reduced coating adhesion and reduced productivity due to delayed alloying reactions occur, so the Si content is preferably set to 1.2% or less. A Si content of less than 0.01% increases manufacturing costs, so 0.01% is the practical lower limit.

[0047] (Mn: 1.5% or more, 3.0% or less) Mn is an element that contributes to improving strength and also acts to suppress ferrite transformation that occurs during heat treatment in continuous annealing equipment or continuous hot-dip galvanizing equipment. If the Mn content is less than 1.5%, this effect is not fully exerted, making it difficult to obtain a sufficient amount of martensite, and therefore high tensile strength (for example, a tensile strength of 900 MPa or more) cannot be obtained. Furthermore, from the viewpoint of suppressing deterioration of bendability, it is not preferable to reduce the Mn content excessively. Therefore, the Mn content is set to 1.5% or more. Mn is preferably 1.7% or more. Mn is more preferably 1.9% or more. On the other hand, if the Mn content exceeds 3.0%, ferrite transformation is excessively suppressed, making it impossible to secure a predetermined amount of ferrite, resulting in a decrease in elongation and a deterioration in ductility at the punched end surface. Therefore, the Mn content is set to 3.0% or less. The Mn content is preferably 2.7% or less.

[0048] (P: 0% or more, 0.020% or less) P is an impurity element that segregates in the center of the steel plate thickness and impairs toughness. Furthermore, P embrittles the welded joint when the steel plate is welded. If the P content exceeds 0.020%, the strength of the weld, hole expandability, and ductility of the punched end face are significantly reduced. Therefore, the P content is set to 0.020% or less. The P content is preferably 0.010% or less. The lower the P content, the better, and there is no particular lower limit. The P content may be 0%. On the other hand, reducing the P content to less than 0.0001% in practical steel plates significantly increases manufacturing costs, resulting in economic disadvantages. Therefore, the lower limit of the P content may be set to 0.0001%.

[0049] (S: 0% or more, 0.0200% or less) S is an impurity element that impairs weldability and manufacturability during casting and hot rolling. S also forms coarse MnS, impairing hole expandability. If the S content exceeds 0.0200%, weldability, manufacturability, hole expandability, and ductility of the punched end face will be significantly reduced. Therefore, the S content is set to 0.0200% or less. The S content is preferably 0.005% or less. The lower the S content, the better, and there is no particular lower limit. The S content may be 0%. On the other hand, if the S content is reduced to less than 0.0001% in practical steel sheets, the manufacturing cost will increase significantly, which will be economically disadvantageous. Therefore, the lower limit of the S content may be set to 0.0001%.

[0050] (Al: 0.001% or more, 1.000% or less) Al is an element that acts as a deoxidizer for steel and stabilizes ferrite, and is contained as needed. When Al is contained, if the content is less than 0.001%, the effect of containing Al is not sufficiently obtained, so the lower limit is set to 0.001% or more. On the other hand, if the content exceeds 1.000%, coarse Al oxides are formed, causing a decrease in ductility of the punched end surface. Therefore, the upper limit is set to 1.000% or less. Preferably, the content is 0.001% or more and 0.50% or less.

[0051] (N: 0% or more, 0.020% or less) N is an element that forms coarse nitrides, impairs bendability and hole expandability, and causes blowholes during welding. If the N content exceeds 0.020%, coarse nitrides are formed, which reduces formability and ductility of the punched end surface, and significantly increases the occurrence of blowholes. Furthermore, from the viewpoint of suppressing deterioration of bendability, a small N content is preferable. Therefore, the N content is set to 0.020% or less. The lower the N content, the more preferable it is, and there is no particular lower limit. The N content may be 0%. On the other hand, if the N content is reduced to less than 0.0005% in practical steel sheets, the manufacturing cost increases significantly, which is economically disadvantageous. Therefore, the lower limit of the N content may be set to 0.0005%.

[0052] (Co: 0 to 0.500%) Co is an element effective in improving the strength of steel sheet. The Co content may be 0%, but to obtain the above effect, the Co content is preferably 0.001% or more, and more preferably 0.010% or more. On the other hand, if the Co content is too high, the ductility of the steel sheet may decrease, which may lead to a decrease in formability. Therefore, the Co content is preferably 0.500% or less.

[0053] (Ni: 0 to 1.000%) Ni, like Co, is an element that is effective in improving the strength of steel sheets. The Ni content may be 0%, but to obtain the above effects, the Ni content is preferably 0.001% or more, and more preferably 0.010% or more. On the other hand, if the Ni content is too high, the ductility of the steel sheet may decrease, which may lead to a decrease in formability. Therefore, the Ni content is preferably 1.000% or less.

[0054] (Mo: 0 to 1.000%) Like Mn, Mo is an element that contributes to increasing the strength of steel sheets. This effect can be achieved even with a small amount. The Mo content may be 0%, but to achieve the above effect, the Mo content is preferably 0.010% or more. On the other hand, if the Mo content exceeds 1.000%, coarse Mo carbides are formed, which may reduce the cold formability of the steel sheet. Therefore, the Mo content is preferably 1.000% or less.

[0055] (Cr:0~2.000%) Like Mn and Mo, Cr is an element that contributes to increasing the strength of steel sheets. This effect can be achieved even with a small amount. The Cr content may be 0%, but to achieve the above effect, the Cr content is preferably 0.001% or more, and more preferably 0.100% or more. On the other hand, if the Cr content exceeds 2.000%, coarse Cr nitrides are formed, which may reduce the cold formability of the steel sheet. Therefore, the Cr content is preferably 2.000% or less.

[0056] (O: 0% or more, 0.0200% or less) O is an element that forms coarse oxides, degrades formability and fracture resistance, and causes blowholes during welding. If the O content exceeds 0.0200%, the presence of coarse oxides significantly deteriorates formability and the ductility of the punched edge, and the occurrence of blowholes becomes more pronounced. Furthermore, from the viewpoint of suppressing deterioration of bendability, a small O content is preferable. Therefore, O is set to 0.0200% or less. The lower the O content, the more preferable it is, and there is no particular lower limit. The O content may be 0%. On the other hand, if the O content is reduced to less than 0.0001% in practical steel sheets, the manufacturing cost increases significantly, which is economically disadvantageous. Therefore, the lower O limit may be set to 0.0001%.

[0057] (Ti: 0 to 0.50%) Ti is an important element for controlling the morphology of carbides. Ti can promote an increase in the strength of ferrite. Ti is also an element that may form coarse Ti oxides or TiN, thereby reducing the formability of the steel sheet. Therefore, from the viewpoint of ensuring the formability of the steel sheet, the lower the Ti content, the better, and it is preferably 0.50% or less, and may even be 0%. However, since reducing the Ti content to less than 0.001% leads to an excessive increase in refining costs, the lower limit of the Ti content may be set to 0.001%.

[0058] (B: 0 to 0.0100%) B suppresses the formation of ferrite and pearlite during the cooling process from austenite and promotes the formation of low-temperature transformation structures such as bainite or martensite. Furthermore, B is a beneficial element for increasing the strength of steel sheets. This effect can be achieved even with trace amounts. Although the B content may be 0%, to achieve the above effect, it is preferable for the B content to be 0.0001% or more. However, if the B content is too high, coarse B oxides are formed, which may become the origin of voids during press forming, potentially reducing the formability of the steel sheet. For this reason, the B content is preferably 0.0100% or less. Note that meticulous care must be taken in the analysis to identify B at less than 0.0001%. If the B content is below the detection limit of an analytical instrument, the B content may be considered to be 0%.

[0059] (Nb: 0 to 0.50%) Like Ti, Nb is an element effective in controlling the morphology of carbides and is also effective in refining the structure and improving the toughness of steel sheets. This effect can be achieved even with trace amounts. The Nb content may be 0%, but to achieve the above effect, the Nb content is preferably 0.0001% or more, and more preferably 0.001% or more. However, if the Nb content is too high, a large number of fine, hard Nb carbides will precipitate, which will increase the strength of the steel sheet and significantly deteriorate the ductility, potentially reducing the formability of the steel sheet. For this reason, the Nb content is preferably 0.50% or less.

[0060] (V:0~0.500%) Like Ti and Nb, V is an element effective in controlling the morphology of carbides and is also effective in refining the structure to improve the toughness of steel sheets. The V content may be 0%, but to obtain the above effects, the V content is preferably 0.001% or more. However, if the V content is too high, a large number of fine V carbides will precipitate, which will increase the strength of the steel material and reduce its ductility, potentially reducing the formability of the steel sheet. For this reason, the V content is preferably 0.500% or less.

[0061] (Cu: 0-0.5%) Cu is an element that contributes to improving the strength of steel sheets. This effect can be achieved even with trace amounts. The Cu content may be 0%, but to achieve the above effect, the Cu content is preferably 0.001% or more. However, if the Cu content is too high, red shortness may occur, which may reduce productivity in hot rolling. For this reason, the Cu content is preferably 0.5% or less.

[0062] (W:0~0.100%) Like Nb and V, W is an element effective in controlling the morphology of carbides and improving the strength of steel sheets. The W content may be 0%, but to obtain the above effects, the W content is preferably 0.001% or more. On the other hand, if the W content is too high, a large number of fine W carbides will precipitate, which will increase the strength of the steel sheet and reduce its ductility, potentially reducing the cold workability of the steel sheet. For this reason, the W content is preferably 0.100% or less.

[0063] (Ta: 0 to 0.100%) Ta, like Nb, V, and W, is an element effective in controlling the morphology of carbides and improving the strength of steel sheets. The Ta content may be 0%, but to obtain the above effects, the Ta content is preferably 0.001% or more. On the other hand, if the Ta content is too high, a large number of fine Ta carbides will precipitate, which will increase the strength of the steel sheet and reduce its ductility, potentially reducing the cold workability of the steel sheet. Therefore, the Ta content is preferably 0.100% or less, more preferably 0.020% or less, and even more preferably 0.010% or less.

[0064] (Sn: 0 to 0.050%) Sn is an element that can be contained in steel sheets when scrap is used as a raw material for the steel sheets. Furthermore, Sn may cause a decrease in the cold formability of the steel sheets due to the embrittlement of ferrite. Therefore, the lower the Sn content, the better. The Sn content is preferably 0.050% or less, more preferably 0.040%, and may even be 0%. However, since reducing the Sn content to less than 0.001% leads to an excessive increase in refining costs, the Sn content may be 0.001% or more.

[0065] (Sb: 0 to 0.050%) Like Sn, Sb is an element that can be contained in steel sheets when scrap is used as a raw material for the steel sheets. Sb may strongly segregate at grain boundaries, causing embrittlement of the grain boundaries, reduced ductility, and reduced cold formability. Therefore, the lower the Sb content, the better. The Sb content is preferably 0.050% or less, more preferably 0.040%, and may even be 0%. However, reducing the Sb content to less than 0.001% results in an excessive increase in refining costs, so the Sb content may be 0.001% or more.

[0066] (As: 0 to 0.050%) Like Sn and Sb, As is an element that can be contained in steel sheets when scrap is used as a raw material for the steel sheets. As is an element that strongly segregates at grain boundaries and may result in a decrease in cold formability. Therefore, the lower the As content, the better. The As content is preferably 0.050% or less, more preferably 0.040% or less, and may even be 0%. However, since reducing the As content to less than 0.001% results in an excessive increase in refining costs, the As content may be 0.001% or more.

[0067] (Mg: 0 to 0.050%) Mg is an element that controls the morphology of sulfides and oxides and contributes to improving the bending formability of steel sheets. This effect can be achieved even with trace amounts. The Mg content may be 0%, but to achieve the above effect, the Mg content is preferably 0.0001% or more. However, if the Mg content is too high, there is a risk of the formation of coarse inclusions, which may cause a decrease in cold formability. For this reason, the Mg content is preferably 0.050% or less, and more preferably 0.040% or less.

[0068] (Ca: 0 to 0.050%) Like Mg, Ca is an element that can control the morphology of sulfides with a small amount. The Ca content may be 0%, but to obtain the above effects, the Ca content is preferably 0.001% or more. However, if the Ca content is too high, coarse Ca oxides are generated, and these Ca oxides can become the starting point for crack generation during cold forming. Therefore, the Ca content is preferably 0.050% or less, and more preferably 0.030% or less.

[0069] (Zr: 0 to 0.050%) Like Mg and Ca, Zr is an element that can control the morphology of sulfides with a small amount. The Zr content may be 0%, but to obtain the above effects, the Zr content is preferably 0.001% or more. However, if the Zr content is too high, coarse Zr oxides may be generated, which may reduce cold formability. Therefore, the Zr content is preferably 0.050% or less, and more preferably 0.040% or less.

[0070] (REM: 0 to 0.100%) REM stands for rare earth metal (rare earth element). Even trace amounts of REM are effective in controlling the morphology of sulfides. The REM content may be 0%, but to achieve the above-mentioned effects, the REM content is preferably 0.001% or more. However, if the REM content is too high, coarse REM oxides may be generated, which may reduce workability and fracture resistance. For this reason, the REM content is preferably 0.100% or less, and more preferably 0.050% or less. Here, REM refers collectively to two elements, scandium (Sc) and yttrium (Y), and 15 elements (lanthanoids) from lanthanum (La) to lutetium (Lu). In this embodiment, "REM" refers to one or more rare earth elements, and the "REM content" refers to the total amount of rare earth elements.

[0071] In the chemical composition of the steel sheet according to this embodiment, the balance excluding the above elements is Fe and impurities. Impurities are elements that are mixed in from the steel raw materials and / or during the steelmaking process, and whose presence is permitted to the extent that they do not impair the properties of the steel sheet according to this embodiment, and are not elements that are intentionally added to the steel sheet.

[0072] The thickness of the steel plate according to this embodiment is not limited to a particular range, but is preferably 0.3 to 6.0 mm in consideration of strength, versatility, and manufacturability.

[0073] Next, a method for manufacturing a steel sheet according to this embodiment will be described. The method for manufacturing a steel sheet according to this embodiment includes the following steps: steelmaking, hot rolling, coiling, holding, cooling, cold rolling, and annealing (continuous annealing). A pickling step may be included between the cooling and cold rolling steps. The manufacturing conditions for each step may be appropriately determined within a range that does not impair the effects of the present invention. However, from the viewpoint of controlling the grain size, average aspect ratio, and Mn concentration at the interface, it is particularly important to appropriately control the conditions for the hot rolling and continuous annealing steps. Each step and condition of the production method will be described in detail below.

[0074] First, a cast slab (hereinafter also simply referred to as a slab) having the component composition of the steel plate according to this embodiment described above is (a-1) The slab temperature in the final finishing stand of hot rolling is 900°C or less, (a-2) Hot rolling is performed with a thickness reduction rate of 30% or more at the final finishing stand of hot rolling (hot rolling process), (a-3) The hot-rolled steel sheet after hot rolling is coiled at a temperature range (coiling temperature) of 450°C or more and 650°C or less (coiling process); (a-4) holding the hot-rolled steel sheet after the coiling process in a temperature range from the coiling temperature to (coiling temperature - 50)°C for a holding time (dwell time) of 8 hours or less (holding process); (a-5) After the holding step, the hot-rolled steel sheet is cooled to 300°C at an average cooling rate of 0.10°C / second or more (cooling step), (a-6) The hot-rolled steel sheet after the cooling step is pickled to obtain an intermediate steel sheet (pickling step); (b) subjecting the intermediate steel sheet after the pickling to cold rolling at a thickness reduction rate of 20% or more and 80% or less to obtain a cold-rolled steel sheet (cold rolling step); (c-1) The cold-rolled steel sheet is heated in an atmosphere with a dew point of −80° C. or higher and 20° C. or lower. (c-2) Heat to 740°C or higher and 900°C or lower, (c-3) An annealing process is carried out in which the material is held (retained) for 60 seconds or more. In addition, (d) In the annealing steps described in (c-1) to (c-3), a coating layer forming step may be carried out to form a coating layer containing zinc, aluminum, magnesium or an alloy thereof on the front and back surfaces of the steel sheet. The (a-6) pickling step is an optional step, and when the pickling step is not performed, the steel sheet that has been subjected to the (a-5) cooling step is used as the intermediate steel sheet.

[0075] In order to enhance fracture resistance, the steel sheet according to this embodiment is cold-rolled and then annealed to produce a fine, equiaxed metallographic structure, and the Mn concentration at the ferrite-martensite interface is controlled. To achieve this, it is effective to uniformly and finely control the hot-rolled structure before cold rolling. To achieve this, it is important to control the temperature and thickness reduction rate of the final finishing stand (hereinafter simply referred to as the final stand) in hot rolling. In other words, it is important to previously achieve a uniform, fine structure by hot rolling, which allows the Mn concentration at the interface to be controlled. As a result, a steel sheet can be produced that simultaneously achieves high levels of strength, formability, and fracture resistance. In addition, the coiling temperature after rolling in the final stand and the cooling rate after the coiling process are also important. Generally, carbides are formed during the process of cooling from the start of coiling to near room temperature. The carbides formed here dissolve during the heating stage of the annealing process and become austenite nucleation sites. Therefore, to achieve a fine, equiaxed grained metal structure after annealing, it is necessary for the carbides to be uniformly dispersed at the start of the annealing process and then dissolve during the heating stage. The concentrations of alloying elements such as Mn and Cr in the carbides are important and have a significant impact on this carbide dissolution. It is generally known that these alloying elements concentrate in the carbides, and the more the alloying elements concentrate, the more difficult it is for the carbides to dissolve in the subsequent annealing process. For example, if the Mn concentration in the carbides exceeds 3 mass%, dissolution becomes significantly more difficult. The concentration of alloying elements in the carbides depends on the thermal history from the start of coiling to near room temperature, and this temperature history must also be appropriately controlled.

[0076] [Hot rolling process] (a-1) Slab temperature in the final stand of hot rolling: 900°C or less As described above, in order to refine the grain sizes of ferrite and bainite after the annealing process, it is necessary to refine the metal structure (hot-rolled structure) after hot rolling. In the hot rolling process of this embodiment, the slab is passed through multiple rolling stands in succession to perform rolling. By setting the temperature of the slab in the final stand of hot rolling to 900°C or less, a large number of nucleation sites for recrystallized grains can be dispersed during hot rolling, and the hot-rolled structure can be refined. This makes it possible to control the grain size of the structure after annealing. If the temperature of the slab in the final stand of hot rolling exceeds 900°C, the hot-rolled structure becomes coarse and mixed-grain, and the structure after the annealing process also becomes coarse, resulting in deterioration of fracture resistance. The temperature of the slab in the final stand of hot rolling is preferably less than 900°C, more preferably 890°C or less, and even more preferably 880°C or less. Although the lower limit of the slab temperature in the final stand is not particularly specified, if it is too low, productivity will decrease due to an increase in the rolling load, and there is a risk of the steel sheet breaking during rolling. Therefore, it may be set to 600°C or higher.

[0077] (a-2) Thickness reduction rate in the final stand of hot rolling: 30% or more The thickness reduction rate in the final stand of hot rolling, like the slab temperature in the final stand, contributes to the refinement of the grain size of ferrite and bainite after the annealing process. A thickness reduction rate of 30% or more allows for the uniform dispersion of nucleation sites for recrystallized grains during hot rolling, thereby refining the hot-rolled structure. This makes it possible to control the grain size of the structure after annealing. If the thickness reduction rate in the final stand is less than 30%, this effect cannot be fully achieved. The thickness reduction rate in the final stand of hot rolling is preferably 40% or more. While there is no particular upper limit for the thickness reduction rate in the final stand, if it is greater than 60%, the effect of refining the hot-rolled structure will saturate, and the rolling load will increase, resulting in an excessive increase in equipment load. Therefore, it is preferable that the thickness reduction rate in the final stand be 60% or less.

[0078] Finish rolling starting temperature: 1100°C or less In the hot rolling step of this embodiment, as described above, recrystallization is utilized to refine the hot-rolled texture. From this viewpoint, it is preferable that the hot rolling step includes a rough rolling step and a finish rolling step including rolling in a final finish stand, and that the rolling start temperature in the finish rolling step is 1100°C or less. In the hot rolling process, finish rolling is usually performed after rough rolling. In the finish rolling of this embodiment, the start temperature of finish rolling is preferably 1100°C or lower. It is generally known that if the thickness reduction rate in finish rolling (particularly the thickness reduction rate in the final stand) is excessively increased, the metal structure is elongated in the rolling direction, making it difficult to obtain a fine and equiaxed hot-rolled structure. However, by performing finish rolling at a relatively low start temperature of finish rolling, strain accumulation during rolling is promoted, making it possible to stably obtain a fine structure as the final structure. In other words, by designing the start temperature of finish rolling to be low, even when large reduction rolling with a relatively large thickness reduction rate is performed in the final stand of finish rolling, the elongation of the structure in the rolling direction can be further suppressed, and a fine and equiaxed structure can be more stably obtained. From this perspective, the start temperature of finish rolling is preferably 1100°C or lower, more preferably 1060°C or lower, and even more preferably 1030°C or lower. Although the lower limit of the start temperature of the finish rolling is not particularly specified, if it is lowered too much, the steel sheet may crack, so it may be set to 950°C or higher.

[0079] [Winding process] (a-3) Winding temperature: 450℃ or higher, 650℃ or lower The hot-rolled steel sheet (hot-rolled steel sheet) is coiled at a temperature range of 450°C or higher and 650°C or lower. If the coiling temperature exceeds 650°C, pearlite transformation progresses, and pearlite containing coarse carbides is formed non-uniformly, making it difficult for the carbides to dissolve in the annealing process, as described above. The remaining undissolved carbides do not function adequately as nucleation sites for austenite, resulting in a coarse and mixed-grain structure after annealing, resulting in a deterioration in fracture resistance. Furthermore, if the coiling temperature exceeds 650°C, the average aspect ratio of the structure after annealing may increase for the same reason as described above. On the other hand, if the coiling temperature is less than 450°C, the strength of the hot-rolled sheet becomes excessive, increasing the cold-rolling load and reducing productivity. The coiling temperature is preferably 480°C or higher and 600°C or lower.

[0080] [Holding process] (a-4) Holding time in the temperature range from the coiling temperature to (coiling temperature - 50)℃: within 8 hours The holding time (dwell time) in the temperature range from the coiling temperature to (coiling temperature - 50)°C affects the concentration of elements such as Mn in the carbides. If the holding time in this temperature range is too long, the concentration of elements such as Mn in the carbides is promoted, resulting in coarse and mixed grains in the post-annealing structure, as described above, and degraded fracture resistance. Furthermore, if the holding time exceeds 8 hours, elements such as Mn may be concentrated in the carbides, increasing the average aspect ratio of the post-annealing structure. Therefore, the holding time is set to 8 hours or less. The holding time is preferably 6 hours or less. On the other hand, from the viewpoint of ensuring time for the transformation of the structure of the hot-rolled steel sheet into a relatively soft structure such as pearlite and improving cold rollability, the holding time is preferably 30 minutes or more. The start of the holding time is the completion of winding. Examples of the holding method include a method of wrapping a heat insulating material around the wound coil, a method of covering the coil with a heat insulating box, and the like. The above-mentioned heat-retaining method (for example, wrapping a heat insulating material around the coil, covering it with a heat-retaining box, etc.) is preferably carried out immediately after the above-mentioned coiling step. If the hot-rolled coil is left in the coiled state for a long time after the coiling step, the temperature of the ends of the hot-rolled coil in the width direction may drop excessively, which may cause the temperature distribution in the hot-rolled coil to become non-uniform. Therefore, the above-mentioned heat-retaining method is preferably carried out within 20 minutes after the coiling step. More preferably, it is within 15 minutes, and even more preferably, it is within 10 minutes.

[0081] [Cooling process] (a-5) Average cooling rate from the holding step to 300°C: 0.10°C / sec or more The average cooling rate from the end of the holding step to 300°C also affects the concentration of elements such as Mn and Cr in the carbides. If the average cooling rate in the temperature range from the end of the holding step to 300°C is too small, the concentration of elements such as Mn in the carbides is promoted, resulting in coarse and mixed grains in the post-annealed structure, as described above, and degraded fracture resistance. Furthermore, if the average cooling rate is less than 0.10°C / s, elements such as Mn may be concentrated in the carbides, increasing the average aspect ratio of the post-annealed structure. Therefore, the average cooling rate is set to 0.10°C / s or more. The average cooling rate is preferably 0.50°C / s or more. On the other hand, if the steel sheet is cooled rapidly, the flatness of the steel sheet may be impaired due to the temperature difference generated in the thickness direction of the steel sheet. Therefore, the average cooling rate is preferably set to 50°C / s or less. The average cooling rate is defined as the temperature drop of the steel sheet from the end of the holding step to 300°C divided by the time required for cooling from the end of the holding step to 300°C.

[0082] [Pickling process] (a-6) After the cooling process, the hot-rolled steel sheet is pickled to produce an intermediate steel sheet. The hot-rolled steel sheet after the cooling step is pickled to obtain an intermediate steel sheet. There are no particular restrictions on the conditions for the pickling step. For example, pickling may be performed once, or may be performed multiple times as necessary. As described above, the pickling step is an optional step in this embodiment, and when the pickling step is not performed, the steel sheet that has undergone the (a-5) cooling step is used as the intermediate steel sheet.

[0083] [Cold rolling process] (b) The intermediate steel sheet is subjected to cold rolling at a thickness reduction rate of 20% to 80% to obtain a cold-rolled steel sheet (hereinafter, steel sheet). Cold rolling is performed at a thickness reduction rate of 20% or more and 80% or less. If the thickness reduction rate is less than 20%, strain accumulation in the steel sheet becomes insufficient, and the austenite nucleation sites during annealing become non-uniform. This results in coarse grain sizes after annealing or mixed grains. Furthermore, if the austenite nucleation sites become non-uniform, the aspect ratio may become large. This deteriorates the fracture resistance. If the thickness reduction rate exceeds 80%, the cold rolling load becomes excessive, and productivity deteriorates. Therefore, the thickness reduction rate is set to 20% or more and 80% or less. Preferably, it is 30% or more and 80% or less. There are no restrictions on the cold rolling method, and the number of rolling passes and the reduction rate per pass may be set appropriately.

[0084] The steel sheet obtained by cold rolling is subjected to continuous annealing. As described above, the distribution of Mn concentration near the interface between ferrite and martensite, i.e., ΔMn, is controlled during continuous annealing. During heating in continuous annealing, Mn is enriched at the prior austenite grain boundaries or the ferrite / austenite interface. During the cooling process, ferrite is formed from austenite. If ferrite is formed in a certain area ratio across the entire steel sheet, the amount of interface movement of a single grain is smaller the smaller the average grain size during heating. In other words, by reducing the average grain size before continuous annealing, the amount of interface movement per grain during the cooling process can be reduced, thereby shortening the distance between the enriched region at the interface before cooling and the interface after ferrite formation. In general, diffusion is faster at grain boundaries than within grains, so the closer the Mn-enriched region is to the grain boundary, the more easily Mn diffuses by grain boundary diffusion. The inventors have found that by controlling the amount of interface movement during cooling in continuous annealing in this way, it is possible to diffuse the Mn-enriched areas during the cooling process and reduce the Mn concentration at the interface between ferrite and martensite, i.e., the above-mentioned ΔMn.

[0085] [Annealing process (continuous annealing process)] (c-1) The cold-rolled steel sheet is subjected to a heat treatment (annealing step) in an atmosphere having a dew point of −80° C. or higher and 20° C. or lower. The dew point inside the furnace during continuous annealing is set to be between -80°C and 20°C. If the dew point is below -80°C, advanced control of the atmosphere inside the furnace is required, which reduces manufacturability and increases costs. The preferred lower limit of the dew point is -70°C or higher, -60°C or higher, -50°C or higher, or -40°C or higher. Furthermore, if the formation of a decarburized layer is not the objective, the preferred upper limit of the dew point is -15°C or lower or -20°C or lower. Furthermore, if a decarburized layer is formed on the surface of the steel sheet, the carbon concentration of the surface layer is affected by the dew point. In order to sufficiently form a decarburized layer, it is preferable to increase the dew point inside the furnace. From this perspective, the dew point when a decarburized layer is formed is preferably above -15°C. A dew point temperature above -15°C facilitates decarburization and reduces the carbon concentration in the surface layer. This improves the bendability of the steel sheet. The preferred lower limit of the dew point when a decarburized layer is formed is -10°C or higher. On the other hand, if the dew point exceeds 20° C., decarburization proceeds excessively, resulting in a decrease in the strength of the steel sheet. Therefore, the dew point is set to 20° C. or less. The upper limit of the dew point is preferably 15° C. or less, or 5° C. or less.

[0086] (c-2) Heat to a temperature range of 740°C or higher and 900°C or lower. The heating temperature in the annealing process affects the area ratio of the metal structure. If the heating temperature is less than 740°C, the amount of austenite produced during heating is small, resulting in an area ratio of martensite and tempered martensite after annealing of less than 40%, resulting in low tensile strength (e.g., less than 900 MPa). Furthermore, if the heating temperature is less than 740°C, the average aspect ratio of the structure after annealing may increase. If the heating temperature exceeds 900°C, the metal structure becomes coarse and the fracture resistance deteriorates. Therefore, the heating temperature is set to 740°C or higher and 900°C or lower. Preferably, it is set to 780°C or higher and 850°C or lower.

[0087] (c-3) Hold (stay) for 60 seconds or more. The holding time (dwell time) at the heating temperature during heating contributes to the amount of austenite produced during heating and affects the area ratio of martensite and tempered martensite after annealing. If the holding time is less than 60 seconds, austenite is not sufficiently produced, the area ratio of martensite and tempered martensite after annealing is less than 40%, and the tensile strength is low (for example, less than 900 MPa). The holding time is preferably 70 seconds or more, more preferably 80 seconds or more. There is no particular upper limit to the holding time, but from the viewpoint of productivity, it may be 1800 seconds or less.

[0088] (d) In the annealing steps described in (c-1) to (c-3) above, a coating layer forming step is carried out to form coating layers containing zinc, aluminum, magnesium or alloys thereof on the front and back surfaces of the steel sheet. In the annealing step, a step of forming a coating layer containing zinc, aluminum, magnesium, or an alloy thereof (for example, a plating layer, an alloy plating layer) on the front and back surfaces of the steel sheet may be carried out. Furthermore, after the annealing step, a coating layer may be formed by a method such as electroplating. [Example]

[0089] The present invention will be described in more detail below with reference to examples, but the present invention is not limited to these examples in any way.

[0090] Example 1 Various steel plates (plate thickness: 1.4 mm) were manufactured using various slabs having the chemical compositions listed in Tables 1A to 1C under various manufacturing conditions listed in Tables 2A to 2D. Pickling was performed between the cooling process and the cold rolling process. In Tables 1A to 1C, blank spaces indicate cases where the corresponding element was not intentionally added to the slab, or where the content of the corresponding element was 0% in significant figures (numbers down to the least significant digit) as defined in this embodiment. The units of the components of each slab are mass %, with the remainder being iron and impurities.

[0091] In the tables, values ​​outside the scope of the invention and values ​​that did not meet the pass / fail criteria are underlined. In Tables 2A to 2D, "presence or absence of plating" indicates whether hot-dip galvanizing was performed in the continuous annealing process, and "presence or absence of alloying" indicates whether alloying was performed after hot-dip galvanizing.

[0092] [Table 1A]

[0093] [Table 1B]

[0094] [Table 1C]

[0095] [Table 2A]

[0096] [Table 2B]

[0097] [Table 2C]

[0098] [Table 2D]

[0099] The metal structure (ferrite, bainite, martensite, tempered martensite, pearlite, and retained austenite (residual γ)) in the range of 1 / 8 to 3 / 8 thickness (1 / 4 part of the plate thickness) centered on the position of 1 / 4 of the plate thickness from the surface of these steel plates, and the total number N of crystal grains in ferrite and bainite in the 1 / 4 part of the plate thickness TIn contrast, the area is 3 μm 2 The ratio of the number N3 of ferrite and bainite grains (N3 / N T ), the total number of grains in ferrite and bainite, N T In contrast, the area is 30 μm 2 The number of ferrite and bainite grains N is greater than or equal to 30 Proportion (N 30 / N T ), the Mn concentration [Mn 1.0 ] and the maximum Mn concentration in the region up to 0.5 μm [Mn 0.5 ] and the difference ΔMn([Mn 0.5 ]-[Mn 1.0 ]), the area at 1 / 4 of the plate thickness is 3 μm 2 The average aspect ratios of ferrite and bainite were evaluated and shown in Tables 3A and 3B. These evaluations were carried out according to the methods described above.

[0100] Furthermore, the tensile strength (TS1), uniform elongation (uEl), local elongation (lEl) and tensile strength of punched-hole tensile test specimens (TS2) of these steel sheets were evaluated and are shown in Tables 3C and 3D. The evaluation methods are as follows.

[0101] The tensile strength (TS1) of the steel plate was evaluated by taking a JIS No. 5 test piece from the steel plate so that the longitudinal direction was perpendicular to the rolling direction of the steel plate, and conducting a tensile test in accordance with JIS Z 2241: 2011. Steel plates with a tensile strength (TS1) of 890 MPa or more (preferably 900 MPa or more) were judged to have passed the test in terms of tensile strength.

[0102] The elongation (uEl, lEl) of the steel sheets was also evaluated by taking JIS No. 5 test pieces from the steel sheets so that the longitudinal direction was perpendicular to the rolling direction of the steel sheets, and conducting tensile tests in accordance with JIS Z 2241:2011. Steel sheets with a uniform elongation (uEl) of 5.5% or more and a local elongation (lEl) of 3.2% or more were judged to pass the formability test.

[0103] The tensile strength (TS2) of a tensile test specimen with a punched hole was evaluated by taking a JIS No. 5 test specimen from the steel plate so that the longitudinal direction was perpendicular to the rolling direction of the steel plate, punching a 10 mm diameter hole at the center of the parallel part with a clearance of 10%, and conducting a tensile test in accordance with JIS Z 2241: 2011. TS2 / TS1 was calculated from the obtained TS2, and a value of 0.50 or more was judged to have passed the fracture resistance test.

[0104] [Table 3A]

[0105] [Table 3B]

[0106] [Table 3C]

[0107] [Table 3D]

[0108] The invention examples (Test Nos. A1 to A37, A56) that satisfied both the component composition and manufacturing conditions had metal structure ratios, structure characteristics, and properties that were all within the scope of the invention, and it was found that steel sheets that could achieve both high levels of formability and fracture resistance could be obtained. On the other hand, in the comparative examples (Test Nos. A38 to A55, A57 to A59) that did not meet the scope of the invention in either the component composition or the manufacturing conditions, at least one of the structure ratio and structure characteristics was outside the scope of the invention, and as a result, one of the properties was deteriorated.

[0109] Test Nos. A38 to A47 are comparative examples in which the component compositions were outside the range of the present invention, but some of the characteristics were deteriorated. Test Nos. A48 to A55 and A57 to A59 are comparative examples in which any of the conditions in the manufacturing method was outside the scope of the present invention.

[0110] Test No. A48 was N3 / N because the temperature of the final stand in hot rolling was too high. T As a result, the TS2 / TS1 and local elongation, which are indicators of the ductility of the punched edge, deteriorated. In test No. A49, the thickness reduction rate in the final stand of hot rolling was too small, so the nucleation sites for recrystallized grains could not be uniformly dispersed in large quantities during hot rolling, resulting in insufficient refinement of the structure after annealing, resulting in the N3 / N T As a result, TS2 / TS1 and local elongation were degraded. In Test No. A50, the coiling temperature was too low, which resulted in a significant increase in the strength of the hot-rolled sheet, which led to an increase in the cold-rolling load, making it impossible to perform cold rolling. Test No. A51 was N3 / N because the coiling temperature was too high. T In addition, N 30 / N T As a result, TS2 / TS1 and local elongation deteriorated. In test No. A52, the holding time in the holding step was too long, which promoted the concentration of elements such as Mn in the carbides, resulting in coarse and mixed grains in the structure after annealing. 30 / N T became high, resulting in degradation of TS2 / TS1. In test No. A53, the average cooling rate in the cooling process was too slow, which promoted the concentration of elements such as Mn in the carbides, resulting in coarse and mixed grains in the structure after annealing. 30 / N T became high, resulting in degradation of TS2 / TS1.

[0111] In test No. A54, the thickness reduction rate in the cold rolling process was too small, resulting in insufficient accumulation of strain in the steel sheet, which resulted in uneven nucleation sites for austenite during annealing, resulting in coarse and mixed grains after annealing, and the N3 / NT As a result, TS2 / TS1 and local elongation were degraded. Test No. A55 was unable to be cold rolled because the thickness reduction rate during the cold rolling process was too large.

[0112] In test No. A57, the heating temperature in the annealing process was too low, resulting in a small amount of austenite during heating and making it impossible to secure martensite and tempered martensite after annealing. As a result, the tensile strength TS1 was significantly reduced, and both the uniform elongation and local elongation were also deteriorated. In test No. A58, the heating temperature during the annealing process was too high, making it impossible to secure ferrite and bainite after annealing. As a result, the area ratio of martensite became excessively high, resulting in deterioration of TS2 / TS1, uniform elongation, and local elongation. In Test No. A59, the holding time in the annealing process was too short, so that sufficient austenite was not generated, and furthermore, undissolved carbides remained in the austenite. As a result, the area ratio of martensite and tempered martensite after annealing was less than 40% and the ductility of each martensite was reduced by the undissolved carbides, resulting in a deterioration in tensile strength TS1 and uniform elongation.

[0113] <Example 2> Next, the case of forming a decarburized layer on the surface layer of a steel sheet will be described. As in Example 1 above, various steel plates (plate thickness: 1.4 mm) were manufactured using various slabs having the chemical compositions listed in Tables 4A to 4C as materials under various manufacturing conditions listed in Tables 5A to 5D. Note that pickling was performed between the cooling process and the cold rolling process. In Tables 4A to 4C, blank cells indicate cases where the corresponding element was not intentionally added to the slab, or where the content of the corresponding element was 0% in significant figures (numbers down to the least significant digit) as defined in this embodiment. The units of the components of each slab are mass %, with the remainder being iron and impurities.

[0114] In each table, values ​​outside the scope of the invention and values ​​that did not meet the pass / fail criteria are underlined. In Tables 5A to 5D, "presence or absence of plating" indicates whether hot-dip galvanizing was performed in the continuous annealing process, and "presence or absence of alloying" indicates whether alloying treatment was performed after hot-dip galvanizing.

[0115] [Table 4A]

[0116] [Table 4B]

[0117] [Table 4C]

[0118] [Table 5A]

[0119] [Table 5B]

[0120] [Table 5C]

[0121] [Table 5D]

[0122] The metal structure (ferrite, bainite, martensite, tempered martensite, pearlite, and retained austenite (residual γ)) in the range of 1 / 8 to 3 / 8 thickness (1 / 4 part of the plate thickness) centered on the position of 1 / 4 of the plate thickness from the surface of these steel plates, and the total number N of crystal grains in ferrite and bainite in the 1 / 4 part of the plate thickness TIn contrast, the area is 3 μm 2 The ratio of the number N3 of ferrite and bainite grains (N3 / N T ), the total number of grains in ferrite and bainite, N T In contrast, the area is 30 μm 2 The number of ferrite and bainite grains N is greater than or equal to 30 Proportion (N 30 / N T ), the difference ΔMn between the Mn concentration at a position 1.0 μm from the interface between ferrite and martensite in a direction perpendicular to the interface and toward the inside of the ferrite grain and the maximum Mn concentration in the region up to 0.5 μm, and the area at 1 / 4 of the plate thickness is 3 μm 2 The average aspect ratios of ferrite and bainite, the average carbon concentration Cs (mass%) at a depth of 10 μm from the steel sheet surface in the sheet thickness direction, the average carbon concentration C4t (mass%) at a depth of 1 / 4 from the steel sheet surface in the sheet thickness direction, and Cs / C4t were evaluated and are shown in Tables 6A to 6D. These evaluations were carried out according to the methods described above.

[0123] Furthermore, the tensile strength (TS1), uniform elongation (uEl), local elongation (lEl), and the tensile strength (TS2) and bendability of the tensile test pieces with punched holes of these steel sheets were evaluated and are shown in Tables 6E and 6F. These evaluation methods are as follows.

[0124] The tensile strength (TS1) of the steel plate was evaluated by taking a JIS No. 5 test piece from the steel plate so that the longitudinal direction was perpendicular to the rolling direction of the steel plate, and conducting a tensile test in accordance with JIS Z 2241: 2011. Steel plates with a tensile strength (TS1) of 890 MPa or more (preferably 900 MPa or more) were judged to have passed the test in terms of tensile strength.

[0125] The elongation (uEl, lEl) of the steel sheets was also evaluated by taking JIS No. 5 test pieces from the steel sheets so that the longitudinal direction was perpendicular to the rolling direction of the steel sheets, and conducting tensile tests in accordance with JIS Z 2241: 2011. Steel sheets with a uniform elongation (uEl) of 5.5% or more and a local elongation (lEl) of 3.2% or more (preferably 3.3% or more) were judged to pass the formability test.

[0126] The tensile strength (TS2) of a tensile test specimen with a punched hole was evaluated by taking a JIS No. 5 test specimen from the steel plate so that the longitudinal direction was perpendicular to the rolling direction of the steel plate, punching a 10 mm diameter hole at the center of the parallel part with a clearance of 10%, and conducting a tensile test in accordance with JIS Z 2241: 2011. TS2 / TS1 was calculated from the obtained TS2, and a value of 0.50 or more was judged to have passed the fracture resistance test.

[0127] The bendability was evaluated by determining the maximum bending angle in a bending test in accordance with the VDA standard (VDA238-100) specified by the German Association of the Automotive Industry. In this example, the displacement at the maximum load obtained in the bending test was converted into an angle according to the VDA standard to determine the maximum bending angle α (°). A maximum bending angle α of 60.0° or greater was determined to be a more preferable form with excellent bendability. The test specimens used in the bending test had dimensions of 30 mm × 60 mm (the side parallel to the rolling direction was 30 mm) and a thickness of 1.4 mm (a thickness of 1.4 mm or less is specified by the VDA standard). The measurement conditions for the bending test were as follows: bend ridge: parallel to the rolling direction; roll diameter: φ30 mm; punch shape: tip R = 0.4 mm; roll spacing: 3.3 mm (the standard specifies thickness × 2 + 0.5 mm); and indentation speed: 20 mm / min.

[0128] [Table 6A]

[0129] [Table 6B]

[0130] [Table 6C]

[0131] [Table 6D]

[0132] [Table 6E]

[0133] [Table 6F]

[0134] The invention examples (Test Nos. B1 to B37, B56, and 57) that satisfied both the component composition and manufacturing conditions had metal structure ratios, structure characteristics, and surface layer characteristics and properties that were all within the scope of the invention, and it was found that steel sheets that could achieve both high levels of formability and fracture resistance could be obtained. On the other hand, in the comparative examples (Test Nos. B38 to B55, B58 to B60) that did not meet the scope of the invention in either the component composition or the manufacturing conditions, at least one of the structure ratio and structure characteristics was outside the scope of the invention, and as a result, one of the characteristics was deteriorated.

[0135] Test Nos. B38 to B47 are comparative examples whose component compositions are outside the scope of the present invention, but some of the characteristics were deteriorated. Test Nos. B48 to B55 and B58 to B60 are comparative examples in which any of the conditions in the manufacturing method was outside the scope of the present invention.

[0136] Test No. B48 was N3 / N because the temperature of the final stand in hot rolling was too high. T As a result, the TS2 / TS1 and local elongation, which are indicators of the ductility of the punched edge, deteriorated. In test No. B49, the thickness reduction rate in the final stand of hot rolling was too small, so the nucleation sites for recrystallized grains could not be uniformly dispersed in large quantities during hot rolling, resulting in insufficient refinement of the structure after annealing, resulting in the N3 / N T As a result, TS2 / TS1 and local elongation were degraded. In Test No. B50, the coiling temperature was too low, which resulted in a significant increase in the strength of the hot-rolled sheet, which led to an increase in the cold-rolling load, making it impossible to perform cold rolling. Test No. B51 was N3 / N because the coiling temperature was too high. T In addition, N 30 / N T As a result, TS2 / TS1 and local elongation deteriorated. In test No. B52, the holding time in the holding step was too long, which promoted the concentration of elements such as Mn in the carbides, resulting in coarse and mixed grains in the structure after annealing. 30 / N T became high, resulting in degradation of TS2 / TS1. In test No. B53, the average cooling rate in the cooling process was too slow, which promoted the concentration of elements such as Mn in the carbides, resulting in coarse and mixed grains in the structure after annealing. 30 / N T became high, resulting in degradation of TS2 / TS1.

[0137] In test No. B54, the thickness reduction rate during the cold rolling process was too small, resulting in insufficient strain accumulation in the steel sheet, which resulted in uneven austenite nucleation sites during annealing, resulting in coarse and mixed grains after annealing, and a N3 / N T As a result, TS2 / TS1 and local elongation were degraded. Test No. B55 was unable to be cold rolled because the thickness reduction rate during the cold rolling process was too large.

[0138] In Test No. B58, the heating temperature in the annealing process was too low, resulting in a small amount of austenite during heating and making it impossible to secure martensite and tempered martensite after annealing. As a result, the tensile strength TS1 was significantly reduced, and both the uniform elongation and local elongation were also deteriorated. In Test No. B59, the heating temperature during the annealing process was too high, making it impossible to secure ferrite and bainite after annealing. As a result, the area ratio of martensite became excessively high, resulting in deterioration of TS2 / TS1, uniform elongation, local elongation, and bendability. In Test No. B60, the holding time in the annealing process was too short, which resulted in insufficient austenite formation, and undissolved carbides remaining in the austenite, resulting in the area ratio of martensite and tempered martensite after annealing being less than 40%. The undissolved carbides also reduced the ductility of each martensite, resulting in a deterioration in tensile strength TS1 and uniform elongation. Furthermore, the holding time in the annealing process was too short, which resulted in an insufficient decarburized layer, resulting in a Cs / C4t ratio exceeding 0.80, resulting in poor bendability. [Industrial Applicability]

[0139] According to the present invention, it is possible to obtain a steel sheet that can achieve high levels of strength, formability, and fracture resistance.

Claims

1. The component composition is, in mass%, C: 0.07-0.15%, Si: 0.01-2.0%, Mn: 1.5-3.0%, P: 0-0.020%, S: 0-0.0200%, Al: 0.001-1.000%, N: 0 to 0.020%, Co: 0 to 0.500%, Ni: 0-1.000%, Mo: 0-1.000%, Cr: 0-2.000%, O: 0 to 0.0200%, Ti: 0 to 0.50%, B: 0 to 0.0100%, Nb: 0 to 0.50%, V: 0 to 0.500%, Cu: 0 to 0.5%, W: 0-0.100%, Ta: 0-0.100%, Sn: 0 to 0.050%, Sb: 0 to 0.050%, As: 0 to 0.050%, Mg: 0 to 0.050%, Ca: 0-0.050%, Zr: 0 to 0.050%, and REM: 0~0.100% and the balance being Fe and impurities, As the tissue fraction, the area ratio of ferrite and bainite is 10% or more and 60% or less in total, the area ratio of martensite and tempered martensite is 40% or more and 90% or less in total, and the area ratio of pearlite and retained austenite is 0% or more and 10% or less in total, The area of ​​the total number of crystal grains in the ferrite and the bainite is 3 μm ２ The proportion of ferrite and bainite grains is 40% or more, Area is 30 μm ２ the proportion of ferrite and bainite grains is 5% or less, a difference ΔMn between the Mn concentration at a position 1.0 μm from the interface between the ferrite and the martensite in a direction perpendicular to the interface and toward the inside of the ferrite grains and the maximum Mn concentration in a region up to 0.5 μm is 1.00 mass % or less.

2. The area is 3 μm ２ 2. The steel sheet according to claim 1, wherein the average aspect ratio of the crystal grains of ferrite and bainite is 1.0 or more and 2.0 or less.

3. 3. The steel sheet according to claim 1, wherein an average carbon concentration at a depth of 10 μm from the surface of the steel sheet in the sheet thickness direction is 0.800 times or less of an average carbon concentration at a depth of 1 / 4 from the surface of the steel sheet in the sheet thickness direction.

4. The component composition is, in mass%, Co: 0.010-0.500%, Ni: 0.010 to 1.000%, Mo: 0.010-1.000%, Cr: 0.001-2.000%, O: 0.0001-0.0200%, Ti: 0.001 to 0.50%, B: 0.0001 to 0.0100%, Nb: 0.001 to 0.50%, V: 0.001-0.500%, Cu: 0.001 to 0.5%, W: 0.001-0.100%, Ta: 0.001 to 0.100%, Sn: 0.001 to 0.050%, Sb: 0.001 to 0.050%, As: 0.001 to 0.050%, Mg: 0.0001-0.050%, Ca: 0.001-0.050%, Zr: 0.001 to 0.050%, and REM: 0.001~0.100% The steel sheet according to any one of claims 1 to 3, characterized in that it contains one or more of the following:

5. A method for manufacturing the steel plate according to claim 1, In mass%, C: 0.07-0.15%, Si: 0.01-2.0%, Mn: 1.5-3.0%, P: 0-0.020%, S: 0-0.0200%, Al: 0.001-1.000%, N: 0 to 0.020%, Co: 0 to 0.500%, Ni: 0-1.000%, Mo: 0-1.000%, Cr: 0-2.000%, O: 0 to 0.0200%, Ti: 0 to 0.50%, B: 0 to 0.0100%, Nb: 0 to 0.50%, V: 0 to 0.500%, Cu: 0 to 0.5%, W: 0-0.100%, Ta: 0-0.100%, Sn: 0 to 0.050%, Sb: 0 to 0.050%, As: 0 to 0.050%, Mg: 0 to 0.050%, Ca: 0-0.050%, Zr: 0 to 0.050%, and REM: 0~0.100% a hot rolling process in which a slab having a component composition containing Fe and impurities is hot rolled in a final finishing stand at a temperature range of 600°C or more and 900°C or less with a thickness reduction rate of 30% or more to obtain a hot rolled steel sheet; a coiling step of coiling the hot-rolled steel sheet at a coiling temperature of 650°C or less and 450°C or more after the hot rolling step; a holding step of holding the hot-rolled steel sheet after the coiling step in a temperature range from the coiling temperature to (the coiling temperature - 50) ° C. for a holding time of 30 minutes or more and 8 hours or less; a cooling step of cooling the hot-rolled steel sheet after the holding step to 300°C at an average cooling rate of 0.10°C / second or more to obtain an intermediate steel sheet; A cold rolling process in which the intermediate steel plate manufactured by the manufacturing method having the above-mentioned steps is cold-rolled at a plate thickness reduction rate of 20% or more and 80% or less to obtain a cold-rolled steel plate; An annealing process in which the cold-rolled steel sheet is annealed by holding it in a temperature range of 740°C to 900°C for 60 seconds or more in an atmosphere having a dew point of -80°C or higher and 20°C or lower; A method for manufacturing a steel sheet, comprising:

6. The method for producing a steel sheet according to claim 5, wherein the dew point is higher than -15°C and 20°C or lower.

7. In the annealing step, A coating layer forming step of forming a coating layer containing zinc, aluminum, magnesium or an alloy thereof on the front and back surfaces of the steel sheet, The method for manufacturing a steel sheet according to claim 5 or 6.

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