High-strength steel sheet, hot-dip galvanized steel sheet, and methods for producing same

Abstract

Provided are: a high-strength steel sheet having a tensile strength (TS) of 1470 MPa or higher and excellent bendability, and further having excellent delayed fracture resistance after being subjected to low-temperature aging treatment; a hot-dip galvanized steel sheet; and methods for producing the same. The present invention provides a high-strength steel sheet characterized in that: the steel sheet has a predetermined component composition and a steel structure in which, at the 1 / 4 sheet thickness position, the area ratio of quenched martensite is 60.0% or greater, the area ratio of ferrite is 30.0% or less, and the volume ratio of retained austenite is 10.0% or less; the dislocation density in the quenched martensite is 1.0 × 1014 / m2 or greater; the dislocation arrangement parameter M in the quenched martensite is 2.0 or less; the C concentration is 5.00 at% or less, the Mn concentration is 5.00 at% or less, and the P concentration is 2.00 at% or less at the prior austenite grain boundary analyzed by a three-dimensional atom probe; and the hydrogen amount in steel after being subjected to low-temperature aging treatment for 20 minutes at 170ºC is 0.20 wt.ppm or less.

Description

High-strength steel sheets, hot-dip galvanized steel sheets, and methods for manufacturing the same.

[0001] This invention relates to high-strength steel sheets, hot-dip galvanized steel sheets, and methods for manufacturing the same.

[0002] CO2 reduction due to vehicle weight reduction ２ To achieve both emission reduction and improved collision resistance through vehicle weight reduction, efforts are being made to increase the strength of automotive steel sheets. Furthermore, new legal regulations are being introduced one after another. Therefore, to increase vehicle body strength, there is an increasing trend towards using high-strength steel sheets, particularly those with a tensile strength (hereinafter simply referred to as TS) of 1180 MPa or higher, for major structural and reinforcing components that form the framework of the automobile cabin (hereinafter also referred to as automotive structural components, etc.). In addition, from the perspective of vehicle body rust prevention, hot-dip galvanized steel sheets, obtained by applying hot-dip galvanizing, are sometimes used as materials for automotive structural components, etc.

[0003] Furthermore, some structural components of automobiles, such as crash boxes, have bent sections. Therefore, from the viewpoint of formability, it is preferable to use steel sheets with high bendability for such components.

[0004] Furthermore, with the increasing strength of steel sheets, there are concerns about delayed fracture caused by the amount of hydrogen entering from the environment in which the steel sheets are used. Therefore, when applying high-strength steel sheets to structural components, there is a demand for high-strength steel sheets with excellent resistance to delayed fracture. Here, in the automobile manufacturing process, after the painting process of the white body, a baking paint process is always performed, which involves heat treatment at a low temperature. Therefore, there is a demand for high-strength steel sheets that have excellent resistance to delayed fracture not in the state of the steel sheet as it is manufactured, but after the baking paint process, i.e., after low-temperature aging treatment, assuming actual use in automobile parts.

[0005] To improve the bendability and delayed fracture resistance of such high-strength steel sheets, a high-strength steel sheet with improved bendability has been proposed by forming a soft layer on the surface. For example, Patent Document 1 proposes a high-strength steel sheet with a tensile strength of 1180 MPa or more, a soft surface layer thickness exceeding 10 μm, and hydrogen trapping sites created by dispersing carbides on the surface, thereby achieving both bendability and hydrogen embrittlement resistance. Patent Document 2 also discloses a steel sheet with improved delayed fracture resistance by controlling the amount of carbon fixed to prior austenite grain boundaries and dislocations.

[0006] International Publication No. 2021 / 186510, International Publication No. 2023 / 032424

[0007] However, Patent Document 1 requires controlling the amount of hydrogen in the steel after low-temperature aging and the carbon segregation at the prior austenite grain boundaries, which are the crack propagation paths for delayed fracture, and does not consider how to achieve both flexibility and delayed fracture resistance after low-temperature aging. Furthermore, Patent Document 2 does not disclose the absolute value of the carbon concentration at the prior austenite grain boundaries, and improving delayed fracture resistance after low-temperature aging requires controlling the segregation of grain boundaries of not only carbon but also other elements.

[0008] In view of the above issues, the present invention aims to provide high-strength steel sheets, hot-dip galvanized steel sheets, and methods for manufacturing the same, which have a tensile strength (TS) of 1470 MPa or more and excellent bendability, and which also have excellent delayed fracture resistance after low-temperature aging treatment.

[0009] The inventors of the present invention have diligently studied to solve the above problems and have obtained the following findings: (1) The base steel sheet is given a predetermined component composition and a steel structure containing a certain amount of quenched martensite. This allows for both high tensile strength and flexibility to be obtained. (2) By setting the dislocation density in the quenched martensite to a certain level or higher and cellularizing the dislocations, delayed fracture resistance characteristics can be obtained after low-temperature aging treatment. (3) Furthermore, by setting the amount of segregation of various elements at the prior austenite grain boundaries to a certain level or lower in advance, low-temperature aging treatment increases carbon segregation to the prior austenite grain boundaries, which are the crack propagation pathways for delayed fracture. This not only increases the strength of the prior austenite grain boundaries, but also reduces the amount of hydrogen in the steel that enters the prior austenite grain boundaries from the environment after low-temperature aging treatment due to the repulsive interaction energy of carbon and hydrogen, resulting in excellent delayed fracture resistance characteristics.

[0010] This invention was completed based on the above findings and further investigations. Specifically, the gist of this invention is as follows:

[0011] [1] A composition comprising, by mass%, C: 0.030% or more and 0.400%, Si: 0.01% or more and 2.50%, Mn: 0.10% or more and 5.00%, P: 0.100% or less, S: 0.0200% or less, Al: 0.100% or less, N: 0.0100% or less, and O: 0.0100% or less, with the remainder being Fe and unavoidable impurities; and a steel structure having, at a position 1 / 4 of the plate thickness, an area ratio of quenched martensite of 60.0% or more, an area ratio of ferrite of 30.0% or less, and a volume ratio of retained austenite of 10.0% or less, wherein the dislocation density in the quenched martensite is 1.0 × 10 １４ / m ２ The high-strength steel sheet is characterized in that the dislocation arrangement parameter M in the quenched martensite is 2.0 or less, the C concentration of the prior austenite grain boundaries analyzed with a three-dimensional atom probe is 5.00 at% or less, the Mn concentration is 5.00 at% or less, and the P concentration is 2.00 at% or less, and the hydrogen content in the steel after low-temperature aging treatment at 170°C for 20 minutes is 0.20 wt. ppm or less.

[0012] [2] The above component composition is further defined in mass percent as follows: Ti: 0.200% or less, Nb: 0.200% or less, V: 0.200% or less, Ta: 0.100% or less, W: 0.100% or less, B: 0.0100% or less, Ni: 1.00% or less, Cr: 1.00% or less, Mo: 1.00% or less, Cu: 1.00% or less, Co: 0.500% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.500% or less, Te: 0.100% or less The high-strength steel sheet according to [1] above, containing at least one element selected from the group consisting of Hf: 0.100% or less and Bi: 0.100% or less.

[0013] [3] A hot-dip galvanized steel sheet having a high-strength steel sheet as described in [1] or [2] above, and a hot-dip galvanized layer formed on the surface of the high-strength steel sheet.

[0014] [4] The hot-dip galvanized steel sheet according to [3] above, wherein the hot-dip galvanized layer is an alloyed hot-dip galvanized layer.

[0015] [5] A step of heating a steel slab having the component composition described in [1] or [2] above, and hot-rolling the steel slab with a finish rolling exit temperature of 750°C or more and 1000°C or less to obtain a hot-rolled steel sheet; a step of winding the hot-rolled steel sheet at 300°C or more and 750°C or less; a step of then cold-rolling the hot-rolled steel sheet with a rolling ratio of 60.0% or less to obtain a cold-rolled steel sheet; and then Ac ３ An annealing process in which the cold-rolled steel sheet is heated to a temperature above the transformation point and below 950°C, and then the cold-rolled steel sheet is subjected to Ac ３A method for manufacturing a high-strength steel sheet, comprising: a cooling step of cooling under the condition that the average cooling rate from the transformation point to 500°C is 2.0°C / s or more; and a strain application step of performing one or both of the following: (I) passing the cold-rolled steel sheet through a skin pass mill having a pair of upper and lower rolls, and (II) passing the cold-rolled steel sheet through a tension leveler having upper and lower rolls arranged alternately; wherein when (I) is performed in the strain application step, the roll diameter of the upper and lower rolls is 1000 mm or less, the temperature of the cold-rolled steel sheet at the entry side of the skin pass mill is 200°C or less, and the elongation rate is 0.01% or more; and when (II) is performed in the strain application step, the roll diameter of the upper and lower rolls is 100 mm or less, the temperature of the cold-rolled steel sheet at the entry side of the tension leveler is 100°C or less, and the value obtained by dividing the intermesh indentation amount by the sheet thickness is 1.3 or more.

[0016] [6] A method for manufacturing a hot-dip galvanized steel sheet, comprising the steps of manufacturing a high-strength steel sheet as described in [5] above, and a step of applying a hot-dip galvanizing treatment to the cold-rolled steel sheet after the cooling step and before the strain application step to obtain a hot-dip galvanized steel sheet, wherein the strain application step is performed on the hot-dip galvanized steel sheet instead of the cold-rolled steel sheet.

[0017] [7] A method for manufacturing a hot-dip galvanized steel sheet according to [6], comprising the step of applying an alloying treatment to the hot-dip galvanized steel sheet after the hot-dip galvanizing treatment and before the strain application step.

[0018] According to the present invention, it is possible to provide high-strength steel sheets, hot-dip galvanized steel sheets, and methods for manufacturing the same, which have a tensile strength (TS) of 1470 MPa or more, excellent bendability, and further excellent delayed fracture resistance after low-temperature aging treatment.

[0019] This is a schematic diagram of the four-point bending jig used in an embodiment of the present invention.

[0020] The following describes high-strength steel sheets, hot-dip galvanized steel sheets, and methods for manufacturing them according to one embodiment of the present invention. Note that the embodiments described below are examples that embody the present invention, and these specific examples do not limit the configuration of the present invention.

[0021] (High-Strength Steel Sheet) First, the component composition of a high-strength steel sheet according to one embodiment of the present invention will be described. Unless otherwise specified, the "%" representing the content of component elements means "mass%".

[0022] [C: 0.030% or more and 0.400% or less] Carbon (C) is one of the important basic components of steel, and in particular in the present invention, it is an important element that affects the area ratio of quenched martensite and ferrite, and the volume ratio of retained austenite. If the C content is less than 0.030%, the area ratio of martensite decreases, and the area ratio of ferrite increases, making it difficult to achieve a total strength (TS) of 1470 MPa or higher. Therefore, the C content should be 0.030% or more, preferably 0.080% or more, and more preferably 0.110% or more. Furthermore, if a higher TS of 1700 MPa or more is desired, the C content should preferably exceed 0.250%, more preferably 0.270% or more, and even more preferably 0.285% or more. On the other hand, if the C content exceeds 0.400%, the toughness of quenched martensite decreases excessively, and good bendability cannot be obtained. Therefore, the C content should be 0.400% or less, preferably 0.380% or less, and more preferably 0.355% or less. Also, if good ductility is desired, even if the TS is less than 1700 MPa, the C content should preferably be 0.250% or less, more preferably 0.230% or less, and even more preferably 0.220% or less.

[0023] [Si: 0.01% or more and 2.50% or less] Si is one of the important basic components of steel, and in particular in the present invention, it is an element that affects the volume fraction of retained austenite because it suppresses the formation of carbides during annealing and promotes the formation of retained austenite. Furthermore, by suppressing the formation of carbides, Si reduces the void formation starting points during bending, making it an effective element for improving bendability. From the viewpoint of obtaining these effects, the Si content should be 0.01% or more, preferably 0.20% or more, and more preferably 0.25% or more. On the other hand, if the Si content exceeds 2.50%, the amount of hydrogen in the steel after low-temperature aging treatment increases due to the increase in the volume fraction of retained austenite, making it difficult to obtain excellent delayed fracture resistance after low-temperature aging. Therefore, the Si content should be 2.50% or less, preferably 2.00% or less, and more preferably 1.50% or less.

[0024] [Mn: 0.10% or more and 5.00% or less] Mn is one of the important basic components of steel, and in particular in the present invention, it is an important element that affects the area ratio of quenched martensite, ferrite, bainite, and the volume ratio of retained austenite. Furthermore, Mn is an important element that improves the hardenability of steel and suppresses ferrite transformation during cooling after annealing, thereby increasing the area ratio of quenched martensite. From the viewpoint of obtaining these effects, the Mn content should be 0.10% or more, preferably 1.00% or more, and more preferably 2.00% or more. On the other hand, if the Mn content exceeds 5.00%, the amount of retained austenite increases, and the amount of hydrogen in the steel after low-temperature aging treatment increases. As a result, good delayed fracture resistance after low-temperature aging treatment cannot be obtained. Therefore, the Mn content should be 5.00% or less, preferably 4.00% or less, and more preferably 3.50% or less.

[0025] [P: 0.100% or less] When P is excessive, it segregates at prior austenite grain boundaries to embrittle the grain boundaries and reduce the ultimate deformability of the steel sheet, which leads to deteriorated bendability. Therefore, the P content is set to 0.100% or less, preferably 0.070% or less. On the other hand, the lower limit of the P content is not particularly limited, but since P is a solid solution strengthening element that can increase the strength of the steel sheet, the P content is preferably 0.001% or more

[0026] [S: 0.0200% or less] S is an element that exists in the form of sulfides and reduces the ultimate deformability of steel, so excessive S content leads to deteriorated bendability. Therefore, the S content is set to 0.0200% or less, preferably 0.0050% or less. On the other hand, the lower limit of the S content is not particularly limited, but in view of production process constraints, the S content is preferably 0.0001% or more.

[0027] [Al: 0.100% or less] When the Al content is excessive, A ３ transformation point rises, and a large amount of ferrite is contained in the microstructure, which makes it difficult to achieve the desired TS. Therefore, the Al content is set to 0.100% or less, preferably 0.050% or less. On the other hand, the lower limit of the Al content is not particularly limited, but since Al suppresses carbide formation during continuous annealing and promotes retained austenite formation, the Al content is preferably 0.001% or more.

[0028] [N: 0.0100% or less] N is an element that exists in the form of nitrides and reduces the ultimate deformability of the steel sheet, so excessive N content leads to deteriorated bendability. Therefore, the N content is set to 0.0100% or less, preferably 0.0050% or less. On the other hand, the lower limit of the N content is not particularly specified, but in view of production process constraints, the N content is preferably 0.0005% or more.

[0029] [O: 0.0100% or less] O exists as an oxide and is an element that reduces the ultimate deformability of steel sheets. Therefore, if O is present in excess, the bendability will decrease. Accordingly, the O content should be 0.0100% or less, and preferably 0.0050% or less. On the other hand, there is no particular lower limit for the O content, but due to constraints in production technology, it is preferable that the O content be 0.0001% or more.

[0030] The composition of the high-strength steel sheet may further contain, in mass%, at least one element selected from the group consisting of Ti: 0.200% or less, Nb: 0.200% or less, V: 0.200% or less, Ta: 0.100% or less, W: 0.100% or less, B: 0.0100% or less, Ni: 1.00% or less, Cr: 1.00% or less, Mo: 1.00% or less, Cu: 1.00% or less, Co: 0.500% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.500% or less, Te: 0.100% or less, Hf: 0.100% or less, and Bi: 0.100% or less.

[0031] [Ti: 0.200% or less] Ti is an element that generates a large amount of coarse precipitates and inclusions, reducing the ultimate deformability of the steel sheet. When the Ti content exceeds 0.200%, the bendability decreases. Therefore, when high-strength steel sheets contain Ti, the Ti content should be 0.200% or less, preferably 0.100% or less. On the other hand, by setting the Ti content to 0.001% or more, fine carbides, nitrides, or carbonitrides are formed during hot rolling or continuous annealing, thereby increasing the strength of the steel sheet and allowing for the favorable acquisition of TS. Therefore, a Ti content of 0.001% or more is preferred.

[0032] [Nb: 0.200% or less] Nb is an element that generates a large amount of coarse precipitates and inclusions, reducing the ultimate deformability of steel sheets. When the Nb content exceeds 0.200%, the bendability decreases. Therefore, when high-strength steel sheets contain Nb, the Nb content should be 0.200% or less, preferably 0.100% or less. On the other hand, by setting the Nb content to 0.001% or more, fine carbides, nitrides, or carbonitrides are formed during hot rolling or continuous annealing, thereby increasing the strength of the steel sheet and allowing for the favorable acquisition of TS. Therefore, an Nb content of 0.001% or more is preferred.

[0033] [V: 0.200% or less] V is an element that generates a large amount of coarse precipitates and inclusions, reducing the ultimate deformability of the steel sheet. If the V content exceeds 0.200%, the bendability decreases. Therefore, when a high-strength steel sheet contains V, the V content should be 0.200% or less, preferably 0.100% or less. On the other hand, by setting the V content to 0.001% or more, fine carbides, nitrides, or carbonitrides are formed during hot rolling or continuous annealing, thereby increasing the strength of the steel sheet and allowing for the favorable acquisition of TS. Therefore, a V content of 0.001% or more is preferred.

[0034] [Ta: 0.100% or less] [W: 0.100% or less] If the content of Ta or W exceeds 0.100%, a large amount of coarse precipitates and inclusions are generated, which reduces the ultimate deformability of the steel sheet and thus reduces its bendability. Therefore, when a high-strength steel sheet contains Ta or W, the content of Ta and W should be 0.100% or less, and preferably 0.080% or less. On the other hand, by setting the content of Ta or W to 0.010% or more, fine carbides, nitrides, or carbonitrides are formed during hot rolling or continuous annealing, thereby increasing the strength of the steel sheet and allowing for the suitable acquisition of TS. Therefore, the content of Ta and W should each be 0.010% or more.

[0035] [B: 0.0100% or less] If the B content is 0.0100% or less, cracks will not form inside the steel sheet during casting or hot rolling, and the ultimate deformability of the steel sheet will not decrease, and the bendability will not decrease. Therefore, when a high-strength steel sheet contains B, the B content should be 0.0100% or less, and preferably 0.0080% or less. On the other hand, B is an element that segregates at the austenite grain boundaries during annealing and improves hardenability, and from the viewpoint of obtaining this effect, the B content should preferably be 0.0003% or more.

[0036] [Ni: 1.00% or less] [Cr: 1.00% or less] [Mo: 1.00% or less] If the content of Ni, Cr, or Mo exceeds 1.00%, the amount of coarse precipitates and inclusions increases, which reduces the ultimate deformability of the steel sheet and thus reduces its bendability. Therefore, when a high-strength steel sheet contains Ni, Cr, or Mo, the content of Ni, Cr, and Mo should be 1.00% or less, and preferably 0.80% or less. On the other hand, Ni, Cr, and Mo are elements that improve hardenability, and from the viewpoint of obtaining this effect, the content of Ni, Cr, and Mo should be 0.01% or more, each.

[0037] [Cu: 1.00% or less] If the Cu content exceeds 1.00%, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet and thus decreasing its bendability. Therefore, when high-strength steel sheets contain Cu, the Cu content should be 1.00% or less, preferably 0.80% or less. On the other hand, Cu is an element that improves hardenability, and from the viewpoint of obtaining this effect, the Cu content should be 0.01% or more.

[0038] [Co: 0.500% or less] If the Co content exceeds 0.500%, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet and thus decreasing its bendability. Therefore, when high-strength steel sheets contain Co, the Co content should be 0.500% or less, preferably 0.300% or less. On the other hand, Co is an element that improves hardenability, and from the viewpoint of obtaining this effect, the Co content should be 0.001% or more.

[0039] [Sn: 0.200% or less] If the Sn content exceeds 0.200%, cracks will form inside the steel sheet during casting or hot rolling, reducing the ultimate deformability of the steel sheet and thus decreasing its bendability. Therefore, when high-strength steel sheets contain Sn, the Sn content should be 0.200% or less, preferably 0.100% or less. On the other hand, Sn is an element that improves hardenability, and from the viewpoint of obtaining this effect, the Sn content should be 0.001% or more.

[0040] [Sb: 0.200% or less] If the Sb content exceeds 0.200%, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet and thus decreasing its bendability. Therefore, when high-strength steel sheets contain Sb, the Sb content should be 0.200% or less, preferably 0.100% or less. On the other hand, Sb is an element that controls the surface softening thickness and enables strength adjustment, and from the viewpoint of obtaining this effect, the Sb content should be 0.001% or more.

[0041] [Ca: 0.0100% or less] [Mg: 0.0100% or less] [REM: 0.0100% or less] If the content of Ca, Mg, or REM exceeds 0.0100%, the amount of coarse precipitates and inclusions increases, which reduces the ultimate deformability of the steel sheet and thus reduces its bendability. Therefore, when a high-strength steel sheet contains Ca, Mg, or REM, the content of Ca, Mg, and REM should be 0.0100% or less, and preferably 0.0050% or less. On the other hand, these elements spheroidize the shape of nitrides and sulfides, improving the ultimate deformability of the steel sheet. From the viewpoint of obtaining this effect, the content of Ca, Mg, and REM should preferably be 0.0005% or more.

[0042] [Zr: 0.500% or less] If the Zr content exceeds 0.500%, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet and thus decreasing its bendability. Therefore, when high-strength steel sheets contain Zr, the Zr content should be 0.500% or less, preferably 0.300% or less. On the other hand, Zr is an element that spheroidizes the shape of nitrides and sulfides, improving the ultimate deformability of the steel sheet. From the viewpoint of obtaining this effect, a Zr content of 0.001% or more is preferable.

[0043] [Te: 0.100% or less] If the Te content exceeds 0.100%, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet and thus decreasing its bendability. Therefore, when high-strength steel sheets contain Te, the Te content should be 0.100% or less, preferably 0.080% or less. On the other hand, Te is an element that spheroidizes the shape of nitrides and sulfides, improving the ultimate deformability of the steel sheet. From the viewpoint of obtaining this effect, a Te content of 0.001% or more is preferable.

[0044] [Hf: 0.100% or less] If the Hf content exceeds 0.100%, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet and thus decreasing its bendability. Therefore, when high-strength steel sheets contain Hf, the Hf content should be 0.100% or less, preferably 0.080% or less. On the other hand, Hf is an element that spheroidizes the shape of nitrides and sulfides, improving the ultimate deformability of the steel sheet. From the viewpoint of obtaining this effect, an Hf content of 0.010% or more is preferable.

[0045] [Bi: 0.100% or less] If the Bi content exceeds 0.100%, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet and thus decreasing its bendability. Therefore, when high-strength steel sheets contain Bi, the Bi content should be 0.100% or less. On the other hand, Bi is an element that reduces segregation, and from the viewpoint of obtaining this effect, a Bi content of 0.001% or more is preferable.

[0046] High-strength steel sheets have a component composition that contains the above-mentioned elements, with the remainder being Fe and unavoidable impurities. Here, in the steel sheet according to one embodiment of the present invention, it is preferable that the steel sheet contains only the above-mentioned basic components and the remainder being Fe (iron) and unavoidable impurities. Here, examples of unavoidable impurities include Zn, Pb, and As. These impurities are permissible to be included as long as their total amount is 0.100% or less. Unavoidable impurities may be introduced from raw materials such as scrap, and their inclusion is permissible as long as it does not hinder the objective of the present invention.

[0047] Next, we will explain the microstructure of high-strength steel sheets. The microstructure of high-strength steel sheets described below refers to the microstructure at the 1 / 4 thickness point (the position corresponding to 1 / 4 of the sheet thickness in the depth direction from the surface of the high-strength steel sheet).

[0048] [Area ratio of quenched martensite: 60.0% or more] By setting the area ratio of quenched martensite to 60.0% or more, the desired TS and excellent bendability can be achieved. Therefore, the area ratio of quenched martensite in high-strength steel sheets should be 60.0% or more, preferably 65.0% or more, more preferably 70.0% or more, even more preferably 75.0% or more, even more preferably 80.0% or more, even more preferably 85.0% or more, and most preferably 90.0% or more. On the other hand, from the viewpoint of obtaining high TS and excellent bendability, the area ratio of quenched martensite should be 99.0% or less, more preferably 98.0% or less, and even more preferably 97.0% or less.

[0049] [Ferrite area ratio: 30.0% or less] By setting the ferrite area ratio to 30.0% or less, the TS increases and the bendability improves. Therefore, the ferrite area ratio in high-strength steel sheets should be 30.0% or less, preferably 25.0% or less, more preferably 15.0% or less, and even more preferably 10.0% or less. On the other hand, the lower limit of the ferrite area ratio is not particularly limited, and the ferrite area ratio may be 0.0%. In this invention, ferrite may also be defined as bainitic ferrite.

[0050] The area ratios of quenched martensite and ferrite can be measured as follows: A sample is cut out so that the thickness cross section (L section) parallel to the rolling direction of the steel sheet becomes the observation surface. After mirror polishing the observation surface with diamond paste, the microstructure is revealed by etching with 3 vol. % nital. Under conditions of an accelerating voltage of 15 kV, a scanning electron microscope (SEM) is used to observe 10 fields of view at a magnification of 2000x, with the observation position set at 1 / 4 of the thickness of the steel sheet. Using the obtained microstructure images, the area ratios of each microstructure (ferrite, quenched martensite, bainite) for 10 fields of view are calculated using Image-Pro from Media Cybernetics. The average of the obtained area ratios is taken as the area ratio of quenched martensite and ferrite in the steel sheet, respectively. Furthermore, in the above microstructure images, ferrite is a flat structure with concave areas and no carbides, while quenched martensite is a structure with convex areas and a microscopic unevenness within the structure.

[0051] [Volume fraction of retained austenite: 10.0% or less] If the volume fraction of retained austenite in the microstructure of high-strength steel sheets is 10.0% or less, the amount of hydrogen in the steel after low-temperature aging treatment of high-strength steel sheets can be reduced, and good delayed fracture resistance after low-temperature aging treatment can be obtained. Therefore, the volume fraction of retained austenite should be 10.0% or less, preferably 9.0% or less, more preferably 8.0% or less, and even more preferably 7.0% or less. On the other hand, the lower limit of the volume fraction of retained austenite is not particularly limited, and the volume fraction of retained austenite may be 0.0%.

[0052] Here, the volume fraction of retained austenite can be measured as follows. The steel sheet is ground such that the 1 / 4 thickness position of the steel sheet becomes the observation surface, and then further polished by 0.1 mm by chemical polishing. For the observation surface, using an X-ray diffractometer with a Co Kα radiation source, the integrated reflection intensities of the (200), (220), and (311) planes of fcc iron (austenite) and the (200), (211), and (220) planes of bcc iron are measured. The volume fraction of retained austenite can be determined from the intensity ratio of the integrated reflection intensity from each plane of fcc iron to the integrated reflection intensity from each plane of bcc iron.

[0053] In addition, the steel microstructure of the high-strength steel sheet may include a remainder structure in addition to quenched martensite, ferrite, and retained austenite. Examples of the remainder structure include those known as other structures of steel sheets, such as pearlite, cementite, and metastable carbides (epsilon (ε) carbide, eta (η) carbide, chi (χ) carbide, etc.). The remainder structure can be identified by observation with an SEM. The area fraction of the remainder structure is preferably 5.0% or less. On the other hand, the area fraction of the remainder structure may be 0.0%.

[0054] The area fraction of the remainder structure can be calculated by the following formula (1). [Area fraction of remainder structure (%)] = 100 - [Area fraction of quenched martensite (%)] - [Area fraction of ferrite (%)] - [Volume fraction of retained austenite (%)] ... (1)

[0055] [Dislocation density in quenched martensite: 1.0 × 10 １４ / m ２ or more] The higher the dislocation density in the quenched martensite, the more remarkable the strain concentration at the prior austenite grain boundaries when strain is applied in the shape correction step, and after the low-temperature aging treatment, the strain in the vicinity of the prior austenite grain boundaries promotes carbon segregation to the prior austenite grain boundaries. To obtain the above effects, the dislocation density in the quenched martensite is set to 1.0 × 10 １４ / m ２ or more, and 3.0 × 10 １４ / m ２The above is preferable. On the other hand, if the dislocation density in the quenched martensite is 1.0 × 10 １６ / m ２ The following conditions effectively prevent an increase in the amount of hydrogen trapped in dislocations, and favorably obtain delayed fracture resistance after low-temperature aging treatment. Therefore, the dislocation density in quenched martensite is 1.0 × 10⁻⁶. １６ / m ２ The following are preferable.

[0056] [Dislocation arrangement parameter M in quenched martensite: 2.0 or less] The dislocation arrangement parameter M is a parameter that indicates the arrangement of dislocations, and is a dimensionless number that estimates the interaction between dislocations from the magnitude of the dislocation strain field and the dislocation density. The lower the dislocation arrangement parameter M in quenched martensite, the stronger the interaction between dislocations, and the more stably the dislocations are arranged. In this case, carbon is more easily trapped in dislocations by low-temperature aging treatment, and hydrogen is prevented from being trapped in dislocations due to the repulsive interaction between carbon and hydrogen, thereby reducing the amount of hydrogen in the steel after low-temperature aging treatment. To obtain the above effect, the dislocation arrangement parameter M in quenched martensite should be 2.0 or less, and preferably 1.5 or less. On the other hand, there is no particular lower limit for the dislocation arrangement parameter M in quenched martensite, but theoretically, from the viewpoint of free energy, dislocations and strain fields always exist in real metallic materials, so this parameter M will be greater than 0.0.

[0057] The dislocation density in quenched martensite can be determined as follows. After polishing a steel plate to a surface at 1 / 4 of its thickness + 0.1 mm, the surface is further polished by chemical polishing to 0.1 mm, and the dislocation density is measured using CoKα rays with an X-ray diffractometer. The dislocation density can be calculated from the strain ε obtained from the full width at half maximum β obtained from the X-ray diffraction measurement. The Williamson-Hall method shown below can be used to extract the strain ε. First, the broadening of the full width at half maximum is influenced by the crystallite size D and the strain ε, and can be calculated as the sum of both factors using the following equation (2). Equation (2) can be rearranged to obtain the following equation (3). As shown in equation (3), by plotting βcosθ / λ against sinθ / λ, the strain ε can be calculated from the slope of the line. The diffraction lines used in the calculation are the {110}, {211}, and {220} planes of bcc iron, where θ represents the peak angle calculated by the θ-2θ method of X-ray diffraction, and λ represents the wavelength of the X-rays used in X-ray diffraction. From the strain ε obtained above, the dislocation density ρ is given by ρ = 14.4ε ２ / b ２ It can be converted as follows: Here, b is the Burgers vector of Fe(α). β = (0.9λ / (D × cosθ)) + 2ε × tanθ ... (2) βcosθ / λ = 0.9λ / D + 2ε × sinθ / λ ... (3)

[0058] The dislocation arrangement parameter M in quenched martensite can be determined from the profile shape described above by calculating the magnitude of the crystal strain field Re due to dislocations, and then using the formula M = Re × √ρ. For calculating Re, the method described in (G. Ribarik, J. Gubicza and T. Ungar: Mater. Sci. Eng. A, 387-389 (2004), 343) can be used. Note that the peak positions of retained austenite and quenched martensite are different, so they are distinguished by their peak positions and the dislocation density in austenite is excluded. Similarly, the peak positions of ferrite and quenched martensite have different lattice constants and axial ratios, so they are distinguished by their lattice constants and axial ratios and the dislocation density in ferrite is also excluded.

[0059] [C concentration of prior austenite grain boundaries analyzed with a 3D atom probe: 5.00 at% or less] The fact that the C concentration of prior austenite analyzed with a 3D atom probe is 5.00 at% or less is an important component of the present invention. By lowering the C concentration of prior austenite grain boundaries after low-temperature aging treatment, it is possible to suppress the formation of carbides at the grain boundaries and increase the amount of solid-solution carbon within the grains. This promotes carbon segregation at the prior austenite grain boundaries after low-temperature aging treatment, and significantly improves the delayed fracture resistance after low-temperature aging treatment. To obtain such effects, the C concentration of prior austenite grain boundaries analyzed with a 3D atom probe should be 5.00 at% or less, and preferably 4.50 at% or less. On the other hand, since the addition of C is necessary to obtain quenched martensite, the C concentration should preferably be 0.01 at% or more.

[0060] [Mn concentration of prior austenite grain boundaries analyzed with a 3D atom probe: 5.00 at% or less] The fact that the Mn concentration of prior austenite analyzed with a 3D atom probe is 5.00 at% or less is an important component of the present invention. Since Mn is one of the elements that embrittles prior austenite grain boundaries, if the Mn concentration of prior austenite grain boundaries exceeds 5.00 at%, even if carbon segregates to the prior austenite grain boundaries after low-temperature aging treatment, good delayed fracture resistance after low-temperature aging treatment cannot be obtained. Therefore, the Mn concentration of prior austenite grain boundaries analyzed with a 3D atom probe should be 5.00 at% or less, and preferably 4.50 at% or less. On the other hand, since the addition of Mn, a hardenable element, is necessary to obtain quenched martensite, the Mn concentration should preferably be 0.01 at% or more.

[0061] [P concentration of prior austenite grain boundaries analyzed with a 3D atom probe: 2.00 at% or less] The fact that the P concentration of prior austenite analyzed with a 3D atom probe is 2.00 at% or less is an important component of the present invention. If P segregates at the prior austenite grain boundaries, after low-temperature aging treatment, site competition with C will prevent C from segregating at the prior austenite grain boundaries, and good delayed fracture resistance after low-temperature aging treatment cannot be obtained. Therefore, the P concentration of prior austenite grain boundaries analyzed with a 3D atom probe should be 2.00 at% or less, and preferably 1.80 at% or less. On the other hand, since the addition of P, a hardenable element, is necessary to obtain quenched martensite, the P concentration should preferably be 0.0001 at% or more.

[0062] As mentioned above, the concentrations of C, Mn, and P at the prior austenite grain boundaries will be those obtained by analysis using a three-dimensional atom probe. This is because, although it is possible to measure the concentrations of elements segregated at the prior austenite grain boundaries using a transmission electron microscope, the sample becomes contaminated by electron beam irradiation, making it impossible to obtain accurate values.

[0063] The concentrations of C, Mn, and P at the prior austenite grain boundaries can be analyzed using a three-dimensional atom probe system (LEAP4000XSi, AMETK). As a sample, a sharp needle-shaped sample is prepared from a region containing the prior austenite grain boundary using the SEM-FIB method, and the measurement is performed in laser mode at a sample temperature of 80K or less. On the obtained three-dimensional atomic map, the regions with concentrated concentrations on the surface are identified as corresponding to the prior austenite grain boundaries. Line analysis is performed across the prior austenite grain boundaries, and the linear distributions of C, Mn, and P concentrations are determined from the number of detected C, Mn, and P ions and the number of other ions. The maximum concentrations of each element in the obtained results are defined as the concentrations of C, Mn, and P at the prior austenite grain boundaries, respectively.

[0064] [Hydrogen content in steel after low-temperature aging treatment at 170°C for 20 minutes: 0.20 wt. ppm or less] In order to ensure good delayed fracture resistance after low-temperature aging treatment, the hydrogen content in steel after low-temperature aging treatment at 170°C for 20 minutes should be 0.20 wt. ppm or less, preferably 0.15 wt. ppm or less. On the other hand, there is no specific lower limit for the hydrogen content in steel after low-temperature aging treatment, but due to production technology constraints, the hydrogen content in steel will generally be 0.001 wt. ppm or more.

[0065] The method for measuring the hydrogen content in steel after low-temperature aging treatment is as follows: A test piece measuring 30 mm in length and 5 mm in width is taken from the product coil. In the case of hot-dip galvanized steel sheets or alloyed hot-dip galvanized steel sheets, the hot-dip galvanized layer or alloyed hot-dip galvanized layer of the test piece is removed by grinding or alkali. Then, the sample is held at 170°C for 20 minutes to perform low-temperature aging treatment, and the amount of hydrogen released from the test piece after treatment is measured by thermal desorption spectroscopy (TDS). Specifically, the test piece after low-temperature aging treatment is continuously heated from room temperature to 300°C at a heating rate of 200°C / h, and then cooled to room temperature. The cumulative amount of hydrogen released from the test piece from room temperature to 210°C is measured and is taken as the hydrogen content in the steel after low-temperature aging treatment.

[0066] The thickness of high-strength steel plates is not particularly limited, but is usually between 0.3 mm and 2.8 mm.

[0067] (Hot-dip galvanized steel sheet) Next, a hot-dip galvanized steel sheet according to one embodiment of the present invention will be described. The hot-dip galvanized steel sheet comprises the high-strength steel sheet described above and a hot-dip galvanized layer formed on the surface of the high-strength steel sheet. The hot-dip galvanized layer may also be an alloyed hot-dip galvanized layer (a plating layer obtained by applying an alloying treatment to hot-dip galvanized plating). The hot-dip galvanized layer may also be provided on both sides of the high-strength steel sheet.

[0068] The composition of the hot-dip galvanized layer is not particularly limited and can be any general composition. For example, it may contain Fe: 20% by mass or less, Al: 0.001% by mass or more and 1.0% by mass or less, and further contain one or more selected from the group consisting of Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr, Co, Ca, Cu, Li, Ti, Be, Bi, and REM in a total amount of 0.0% by mass or more and 3.5% by mass or less, with the remainder being Zn and unavoidable impurities. In the case of a non-alloyed hot-dip galvanized layer, the Fe content in the plating layer can be less than 7% by mass. In the case of an alloyed hot-dip galvanized layer, the Fe content in the plating layer is preferably 7% by mass or more and 15% by mass or less, and more preferably 8% by mass or more and 13% by mass or less.

[0069] The amount of plating deposited on one side of the hot-dip galvanized layer is 20 g / m². ２ 80g / m or more ２ The following are preferable.

[0070] (Method for manufacturing high-strength steel sheets) Next, a method for manufacturing high-strength steel sheets according to one embodiment of the present invention will be described.

[0071] [Steel Material Melting] First, a steel material having the above-described component composition is melted to produce a steel slab. In this invention, the method of melting the steel material is not particularly limited, and any known melting method, such as a converter or electric furnace, is suitable. Furthermore, while it is preferable to manufacture the steel slab (slab) by continuous casting to prevent macrosegregation, it can also be manufactured by ingot casting or thin slab casting. In addition to the conventional method of cooling the steel slab to room temperature and then reheating it after manufacturing, energy-saving processes such as direct rolling and direct rolling, in which the slab is charged into the heating furnace while still hot without cooling to room temperature, or rolled immediately after a short period of heat retention, can also be applied without any problems.

[0072] [Heating the Steel Slab] Precipitates present during the heating stage of the steel slab exist as coarse precipitates in the final steel sheet and do not contribute to strength. Therefore, it is preferable to remelt the Ti and Nb precipitates that precipitated during casting. Accordingly, the heating temperature of the steel slab is preferably 1100°C or higher, and more preferably 1150°C or higher. Furthermore, from the viewpoint of scaling off defects such as bubbles and segregation on the slab surface, reducing cracks and irregularities on the steel sheet surface, and achieving a smooth steel sheet surface, the heating temperature of the steel slab is preferably 1100°C or higher. On the other hand, if the heating temperature of the steel slab exceeds 1300°C, scale loss increases with the increase in oxidation. Therefore, the heating temperature of the steel slab is preferably 1300°C or lower, and more preferably 1250°C or lower.

[0073] [Finish rolling exit temperature for hot rolling: 750°C or higher and 1000°C or lower] After heating, the steel slab is hot-rolled into a hot-rolled steel sheet by rough rolling and finish rolling. If the finishing temperature is below 750°C, the rolling load increases, resulting in a greater rolling load. Furthermore, the reduction ratio increases as the austenite remains in an unrecrystallized state, leading to the development of an abnormal texture, significant in-plane anisotropy in the final product, and not only is the uniformity (stability) of the material impaired, but the bendability itself also decreases. Therefore, the finish rolling exit temperature for hot rolling should be 750°C or higher, preferably 800°C or higher. On the other hand, if the finish rolling exit temperature exceeds 1000°C, the amount of oxide (scale) generated increases rapidly, the interface between the base metal and the oxide becomes rough, and the surface quality after pickling and cold rolling tends to deteriorate. In addition, if some hot-rolled scale remains after pickling, it adversely affects the bendability. Furthermore, excessively coarse grain size can occur, leading to surface roughness in the pressed product during processing. Therefore, the finish rolling exit temperature for hot rolling should be 1000°C or lower, preferably 950°C or lower.

[0074] While steel slabs are normally formed into sheet bars by rough rolling under normal conditions, if the heating temperature is set low, it is preferable to heat the sheet bars using a bar heater or the like before finish rolling to prevent problems during hot rolling.

[0075] Furthermore, rough-rolled sheets may be joined together during hot rolling and finish rolling may be performed continuously. Alternatively, the rough-rolled sheets (sheet bars) may be wound up before finish rolling. In addition, to reduce the rolling load during hot rolling, part or all of the finish rolling may be performed using lubricated rolling. Performing lubricated rolling is also effective from the viewpoint of uniformizing the shape of the steel sheets and uniformizing the material. It is preferable that the coefficient of friction during lubricated rolling be in the range of 0.10 to 0.25.

[0076] Next, the hot-rolled steel sheet may be pickled. Pickling removes oxides from the surface of the steel sheet, thereby ensuring good chemical conversion treatment properties and plating quality in the final high-strength steel sheet product. Pickling may be performed only once or in multiple stages.

[0077] [Winding temperature: 300°C or higher and 750°C or lower] Next, the hot-rolled steel sheet is wound up. If the winding temperature is below 300°C, the strength of the hot-rolled steel sheet increases, increasing the rolling load in cold rolling and causing defects in the sheet shape, thus reducing productivity. Therefore, the winding temperature after hot rolling should be 300°C or higher, preferably 400°C or higher. On the other hand, if the winding temperature after hot rolling exceeds 750°C, scale will form on the surface of the hot-rolled steel sheet, and during the subsequent heat treatment process, excessive surface decarburization will reduce the surface hardness, resulting in the inability to obtain the desired TS. Therefore, the winding temperature after hot rolling should be 750°C or lower, preferably 650°C or lower.

[0078] [Cold rolling reduction ratio: 60.0% or less] Hot-rolled steel sheets are wound up, pickled as necessary, and then cold-rolled. If the cold rolling reduction ratio exceeds 60.0%, the austenite grain size formed in the subsequent annealing process becomes finer, improving the stability of the austenite and resulting in an excess of retained austenite. Therefore, the cold rolling reduction ratio should be 60.0% or less, preferably 55.0% or less, and more preferably 50.0% or less. On the other hand, if the cold rolling reduction ratio is low, the frequency of austenite nucleation during annealing decreases, and the area of ​​prior austenite grain boundaries increases due to austenite grain coarsening, which increases the segregation of C, P, and Mn, making it impossible to obtain good delayed fracture resistance after low-temperature aging treatment. Therefore, the cold rolling reduction ratio should preferably be 5.0% or more, and preferably 10.0% or more. Furthermore, since strain is introduced uniformly and efficiently, resulting in a uniform structure, it is preferable to perform cold rolling using multi-pass rolling that requires two or more passes, such as tandem multi-stand rolling or reverse rolling.

[0079] [Heating temperature for annealing process: Ac ３ [Above transformation point and below 950°C] Next, the cold-rolled steel sheet is Ac ３ An annealing process is performed, in which the material is heated to a temperature above the transformation point but below 950°C. The heat treatment method for the annealing process can be either continuous annealing or batch annealing. The heating temperature for the annealing process is Ac ３ Below the transformation point, annealing occurs in a two-phase region of ferrite and austenite, resulting in the inclusion of a large amount of ferrite after annealing, making it difficult to achieve the desired TS and high bendability. Therefore, the heating temperature in the annealing process should be Ac ３ Assuming it is above the transformation point, (Ac ３ (Transformation point + 20°C) or higher is preferred, (Ac ３A temperature of 950°C (transformation point + 50°C) or higher is more preferable. On the other hand, if the heating temperature in the annealing process exceeds 950°C, the austenite grain size becomes coarser during annealing, reducing the grain boundary area. This increases the carbon concentration at the prior austenite grain boundaries, preventing the acquisition of good delayed fracture resistance after low-temperature aging treatment. Therefore, the heating temperature in the annealing process should be 950°C or lower, preferably 930°C or lower, and more preferably 910°C or lower. The heating temperature is measured based on the temperature of the steel sheet surface. Furthermore, the heat retention time at the heating temperature is preferably 10 s to 600 s. ３ The transformation point (°C) is determined by the following formula (4): Ac ３ Transformation point = 910 - 203√(%C) + 45 × (%Si) - 30 × (%Mn) - 20 × (%Cu) - 15 × (%Ni) + 11 × (%Cr) + 32 × (%Mo) + 104 × (%V) + 400 × (%Ti) + 200 × (%Al) ... (4) Here, (%C), (%Si), (%Mn), (%Ni), (%Cu), (%Cr), (%Mo), (%V), (%Ti), (%W), and (%Al) are the mass %) content of each element in the composition of the high-strength steel sheet, and are set to zero if not contained.

[0080] [Ac ３ [Average cooling rate from the transformation point to 500°C is 2.0°C / s or higher] Next, the cold-rolled steel sheet is subjected to Ac ３ A cooling process is performed under the condition that the average cooling rate from the transformation point to 500°C is 2.0°C / s or higher. ３ If the average cooling rate from below the transformation point to 500°C is less than 2.0°C / s, Mn segregates at the prior austenite grain boundaries, and good delayed fracture resistance after low-temperature aging treatment cannot be obtained. Therefore, the average cooling rate should be 2.0°C / s or higher, and preferably 5.0°C / s or higher. On the other hand, there is no particular upper limit to the average cooling rate, but it is generally 50°C / s or lower.

[0081] As described above, the cold-rolled steel sheet is cooled to 500°C, and then cooled to room temperature. However, the cooling method and cooling rate are not particularly limited, and furnace cooling, air cooling, gas jet cooling, mist cooling, water cooling, etc., may be used in batch annealing. In addition, pickling treatment by conventional methods may be performed.

[0082] [Strain application process] Next, a strain application process is performed which involves either or both of the following steps: (I) passing the cold-rolled steel sheet through a skin pass mill having a pair of upper and lower rolls, and (II) passing the cold-rolled steel sheet through a tension leveler having upper and lower rolls arranged alternately.

[0083] [Skin pass mill: Upper and lower roll diameter 1000 mm or less, entry temperature 200°C or less, elongation rate 0.01% or more] When (I) is performed in the strain application process, the upper and lower roll diameters should be 1000 mm or less, the temperature of the cold-rolled steel sheet at the entry side of the skin pass mill should be 200°C or less, and the elongation rate should be 0.01% or more. If the roll diameters of the upper and lower rolls of the skin pass mill exceed 1000 mm, sufficient processing will not be introduced during skin pass, and the desired dislocation arrangement parameter M cannot be obtained. Therefore, the roll diameters of the upper and lower rolls of the skin pass mill should be 1000 mm or less, preferably 800 mm or less. On the other hand, there is no particular lower limit to the roll diameter of the upper and lower rolls of the skin pass mill, but it should generally be 200 mm or more.

[0084] If the temperature of the cold-rolled steel sheet at the entry side of the skin pass mill exceeds 200°C, P segregation to the prior austenite grain boundaries is promoted, and good delayed fracture resistance after low-temperature aging treatment cannot be obtained. Therefore, the temperature of the cold-rolled steel sheet at the entry side of the skin pass mill (entry side temperature) should be 200°C or lower. On the other hand, if the entry side temperature is too low, the toughness of the steel sheet decreases, and cracking may occur due to skin passing. Therefore, it is preferable that the temperature of the cold-rolled steel sheet at the entry side of the skin pass mill be 0°C or higher. Note that the temperature of the cold-rolled steel sheet may rise while passing through the skin pass mill, and the temperature of the cold-rolled steel sheet at the exit side of the skin pass mill may become higher than the temperature at the entry side, but the effects of the present invention can be obtained as long as the entry side temperature is within the above range.

[0085] If the elongation rate in the skin pass mill is less than 0.01%, sufficient processing is not introduced, and the desired dislocation arrangement parameter M cannot be obtained. Therefore, the elongation rate in the skin pass mill should be 0.01% or higher, and preferably 0.03% or higher. On the other hand, there is no particular upper limit to the elongation rate in the skin pass mill, but it is generally 0.5% or lower, and preferably 0.1% or lower.

[0086] [Tension leveler: Upper and lower roll diameters 100 mm or less, entry temperature 100°C or less, intermesh indentation amount divided by plate thickness 1.3 or more] When (II) is performed in the strain application process, the upper and lower roll diameters shall be 100 mm or less, the temperature of the cold-rolled steel sheet at the entry side of the tension leveler shall be 100°C or less, and the value obtained by dividing the intermesh indentation amount by plate thickness shall be 1.3 or more. If the roll diameters of the upper and lower rolls of the tension leveler exceed 100 mm, sufficient processing by the tension leveler will not be introduced, and the desired dislocation density cannot be obtained. Therefore, the roll diameters of the upper and lower rolls of the tension leveler shall be 100 mm or less. On the other hand, there is no particular lower limit to the roll diameters of the upper and lower rolls of the tension leveler, but it shall generally be 20 mm or more.

[0087] If the temperature of the cold-rolled steel sheet at the entry side of the tension leveler exceeds 100°C, P segregation to the prior austenite grain boundaries is promoted, and good delayed fracture resistance after low-temperature aging treatment cannot be obtained. Therefore, the temperature of the cold-rolled steel sheet at the entry side of the tension leveler should be 100°C or lower. On the other hand, if the temperature of the cold-rolled steel sheet at the entry side of the tension leveler is too low, the toughness of the steel sheet decreases, and cracking may occur due to shape correction by the tension leveler. Therefore, it is preferable that the temperature of the cold-rolled steel sheet at the entry side of the tension leveler be 0°C or higher. Note that the temperature of the cold-rolled steel sheet may rise while passing through the tension leveler, and the temperature of the cold-rolled steel sheet at the exit side of the tension leveler may become higher than the temperature at the entry side, but the effects of the present invention can be obtained as long as the entry side temperature is within the above range.

[0088] In a tension leveler, if the value obtained by dividing the intermesh indentation amount by the plate thickness is less than 1.3, sufficient processing is not introduced during the bending-unbending process by the tension leveler. As a result, the desired dislocation arrangement parameter M cannot be obtained, and good delayed fracture resistance after low-temperature aging treatment cannot be obtained. Therefore, in a tension leveler, the value obtained by dividing the intermesh indentation amount by the plate thickness should be 1.3 or more, preferably 1.5 or more. On the other hand, if the processing amount is too high, cracks may occur inside the steel plate, and the bendability may deteriorate, so this value should preferably be 8.0 or less.

[0089] Furthermore, if both (I) and (II) are performed in the strain application process, (II) should be performed after (I). In this case, the conditions for both the skin pass mill and tension leveler described above must be met.

[0090] (Method for manufacturing hot-dip galvanized steel sheet) Next, a method for manufacturing a hot-dip galvanized steel sheet according to one embodiment of the present invention will be described. The method for manufacturing a hot-dip galvanized steel sheet comprises the steps of the above-described method for manufacturing a high-strength steel sheet, and a step of applying a hot-dip galvanizing treatment to a cold-rolled steel sheet after the cooling step and before the strain application step to obtain a hot-dip galvanized steel sheet, wherein the strain application step is performed on the hot-dip galvanized steel sheet instead of the cold-rolled steel sheet. Furthermore, the method may also include a step of applying an alloying treatment to the hot-dip galvanized steel sheet after the hot-dip galvanizing treatment and before the strain application step.

[0091] [Hot-dip zinc plating process] After the cooling process and before the strain application process described above, the cold-rolled steel sheet is subjected to hot-dip zinc plating. Alternatively, an alloying treatment may be performed after the hot-dip zinc plating. In addition, annealing, cooling, and plating may be performed continuously on a single line (CGL (Continuous Galvanizing Line)). For example, after annealing, the cold-rolled steel sheet is cooled to a temperature range of about 500°C. Next, the cold-rolled steel sheet is passed through the steel strip exit side of the cooling strip and moved into the hot-dip zinc plating bath via a snout whose tip is immersed in the hot-dip zinc plating bath, while being further cooled. The time from the end of cooling of the cold-rolled steel sheet until the cold-rolled steel sheet enters the hot-dip zinc plating bath is not particularly limited, but from the viewpoint of controlling the area ratio of martensite and ferrite, it is preferable to set it to 1 s or more and 300 s or less. Furthermore, a roll is provided immediately before the connection between the cooling zone and the snout to change the direction of travel of the cold-rolled steel sheet and allow it to enter the snout. The cold-rolled steel sheet passes through this roll before entering the snout. Next, the cold-rolled steel sheet, guided through the snout to a molten zinc plating bath, is immersed in the molten zinc plating bath and subjected to molten zinc plating to become a molten zinc plated steel sheet.

[0092] In the hot-dip zinc plating process, for example, cold-rolled steel sheets are immersed in a hot-dip zinc plating bath at a temperature of 480°C to 500°C. It is also preferable to use a hot-dip zinc plating bath with an Al content of 0.10% to 0.23% by mass, with the remainder being Zn and unavoidable impurities.

[0093] Furthermore, an alloying treatment may be performed after the hot-dip zinc plating treatment described above. When the alloying treatment temperature is 460°C or higher, a suitable Zn-Fe alloying rate can be obtained, and productivity can be favorably achieved. Therefore, the alloying treatment temperature is preferably 460°C or higher, and more preferably 470°C or higher. On the other hand, when the alloying treatment temperature is 600°C or lower, the transformation of untransformed austenite into pearlite is favorably prevented, and a suitable area ratio of quenched martensite can be obtained. Therefore, the alloying treatment temperature is preferably 600°C or lower, and more preferably 560°C or lower.

[0094] Furthermore, the plating adhesion amount is 20 g / m² per side. ２ 80g / m or more ２ The following (double-sided plating) is preferred. The amount of plating can be adjusted by performing gas wiping or the like after the hot-dip zinc plating treatment.

[0095] Although the above explanation mainly focused on hot-dip zinc plating and alloyed hot-dip zinc plating, plating layers such as Zn plating, Zn-Ni alloy plating, or Al plating may also be formed by electroplating. In one example, the plating layer can be an electroplated zinc layer. When forming an electroplated zinc layer, for example, a plating bath can be used that contains Ni: 9% to 25% by mass, with the remainder being Zn and unavoidable impurities. Furthermore, when forming an electroplated zinc layer, it is preferable to use a plating bath at room temperature or above 100°C. In the case of electroplating, the plating adhesion amount is 15 g / m² per side. ２ More than 100g / m ２ The following (double-sided plating) is preferable.

[0096] For processes and conditions not described in this specification, conventional methods may be used.

[0097] The operation and effects of the present invention will be described below using examples. However, the present invention is not limited to the following examples.

[0098] Steel slabs (steel material) having the component composition shown in Table 1, with the remainder being Fe and unavoidable impurities, were melted in a converter, and steel slabs were obtained by continuous casting. The obtained steel slabs were heated to 1250°C and roughly rolled to obtain sheet bars. Next, the obtained sheet bars were subjected to finish rolling to obtain hot-rolled steel sheets, which were then coiled. After pickling the obtained hot-rolled steel sheets, they were cold-rolled to obtain cold-rolled steel sheets. Next, the obtained cold-rolled steel sheets were annealed. In some cases, the cold-rolled steel sheets were subjected to hot-dip galvanizing to obtain hot-dip galvanized steel sheets. The manufacturing conditions are shown in Table 2. Note that "plate thickness" in Table 2 refers to the plate thickness after cold rolling.

[0099] In the "Plating Type" column of Table 2, "CR" indicates cases where no plating treatment was performed, "GI" indicates cases where only hot-dip zinc plating treatment was performed (resulting in a hot-dip zinc plated steel sheet without alloying treatment), and "GA" indicates cases where alloying treatment was performed after hot-dip zinc plating treatment (resulting in an alloyed hot-dip zinc plated steel sheet).

[0100] In the GI model, a molten zinc plating bath containing 0.20% by mass of Al, with the remainder being Zn and unavoidable impurities, was used. In the GA model, a molten zinc plating bath containing 0.14% by mass of Al, with the remainder being Zn and unavoidable impurities, was used. The plating bath temperature was 470°C in both cases. The plating deposition amount for the GI model was 45 to 72 g / m² per side. ２ (Double-sided plating) is used, and for GA, the rate is 45 g / m² per side. ２ The plating process was limited to (double-sided plating). In addition, for GA, the alloying treatment temperature was set to approximately 550°C.

[0101] Furthermore, the composition of the molten zinc plating layer of GI was Fe: 0.1 to 1.0 mass%, Al: 0.2 to 1.0 mass%, with the remainder being Zn and unavoidable impurities. The composition of the alloyed molten zinc plating layer of GA was Fe: 7 to 15 mass%, Al: 0.1 to 1.0 mass%, with the remainder being Zn and unavoidable impurities.

[0102]

[0103]

[0104] The steel sheets thus obtained were subjected to the following measurements using the method described above: the microstructure at 1 / 4 of the sheet thickness, the dislocation density and dislocation arrangement parameter M in the quenched martensite, the C, Mn, and P concentrations at the prior austenite grain boundaries, and the hydrogen content in the steel after low-temperature aging. The Burgers vector of Fe(α) used to determine the dislocation density was set to 0.25 nm. The measurement results are shown in Table 3.

[0105] For each example, the tensile properties, bendability, and delayed fracture resistance after low-temperature aging treatment were evaluated using the following test methods. The evaluation results are shown in Table 3. The component composition of the high-strength steel plates in each example was substantially the same as that of the steel slabs. The suitable steels were all within the range of component composition described above, while the comparative steels were all outside the range of component composition described above.

[0106] [Tensile Test] The tensile test was conducted in accordance with JIS Z 2241:2011. From the obtained steel plate, a JIS No. 5 test specimen was taken so that the tensile direction was perpendicular to the rolling direction of the steel plate, and the crosshead speed was 1.67 × 10⁻⁶. －１ Tensile tests were conducted under mm / s conditions, and the TS was measured. A TS of 1470 MPa or higher was considered acceptable. For steel sheets with a TS of 1470 to less than 1700 MPa, EL was also measured as an indicator of ductility. Steel sheets with a TS of 1470 to less than 1700 MPa are sometimes used for structural components of automobiles with complex shapes, and therefore ductility may also be required from the standpoint of formability, and an EL of 6.0% or higher is preferable.

[0107] [Bending Test] The bending test was conducted in accordance with JIS Z 2248:2022. From the steel plate 24 hours after manufacturing, a strip-shaped test piece with a width of 30 mm and a length of 100 mm was taken so that the direction parallel to the rolling direction of the steel plate was the axial direction of the bending test. Then, a 90° V bending test was performed under the conditions of a pressing load of 100 kN and a pressing holding time of 5 seconds. After that, the presence or absence of cracks was determined on the outer side of the bend using a stereomicroscope, and the smallest bending radius at which no cracks occurred was defined as the limiting bending radius R. In this invention, the bending properties of the steel plate were judged to be good if the limiting bending radius R / t ≤ 4.00 (t: thickness of the steel plate) in the 90° V bending test was satisfied.

[0108] [Evaluation Test of Delayed Fracture Resistance After Low-Temperature Aging] The evaluation test of delayed fracture resistance after low-temperature aging was performed as follows. A steel plate cut to 20 mm x 75 mm was ground down by 2 mm on each end to prepare a 16 mm x 75 mm strip-shaped sample. Then, a 3% tensile strain was applied, and after unloading, low-temperature aging treatment was performed at 170°C for 20 minutes. The sample after low-temperature aging treatment was mounted on the four-point bending jig shown in Figure 1. The four-point bending jig 100 in Figure 1 has a structure in which the sample 10 is fixed by four glass rods 20 and stress is applied by pressing the center. A stress of 1000 MPa was applied using the dial gauge method with the four-point bending jig. Three samples were prepared with stress applied in the same manner, and the specific liquid volume was 15 mL / cm³. ２ The samples were immersed in hydrochloric acid at pH 3 for 96 hours. If no cracks occurred in any of the three samples after immersion, it was determined that they had superior delayed fracture resistance after low-temperature aging treatment.

[0109]

[0110] As shown in Table 3, the present invention example exhibits superior TS, bendability, and delayed fracture resistance after low-temperature aging treatment. On the other hand, the comparative example is inferior in one or more of the following: TS, bendability, and delayed fracture resistance after low-temperature aging treatment.

[0111] Although embodiments of the present invention have been described above, the present invention is not limited by the descriptions that constitute part of the disclosure of the present invention according to these embodiments. That is, other embodiments, examples, and operational techniques made by those skilled in the art based on these embodiments are all included in the scope of the present invention. For example, in the series of heat treatments in the manufacturing method described above, the equipment for heat-treating the steel plate is not particularly limited, as long as the thermal history conditions are satisfied.

[0112] According to the present invention, it is possible to provide high-strength steel sheets, hot-dip galvanized steel sheets, and methods for manufacturing the same, which have a tensile strength (TS) of 1470 MPa or more and excellent bendability, as well as excellent delayed fracture resistance after low-temperature aging treatment. By applying these high-strength steel sheets or hot-dip galvanized steel sheets to automotive structural members, it is possible to improve fuel efficiency by reducing the weight of the vehicle body, and thus have great industrial value.

[0113] 100 4-point bending jig 10 samples 20 glass rods

Claims

1. A steel structure having the following composition by mass%, containing C: 0.030% to 0.400%, Si: 0.01% to 2.50%, Mn: 0.10% to 5.00%, P: 0.100% or less, S: 0.0200% or less, Al: 0.100% or less, N: 0.0100% or less, and O: 0.0100% or less, with the remainder being Fe and unavoidable impurities; and a steel structure having the area ratio of quenched martensite being 60.0% or more, the area ratio of ferrite being 30.0% or less, and the volume ratio of retained austenite being 10.0% or less, with a dislocation density in the quenched martensite being 1.0 × 10⁻⁶ １４ / m ２ The high-strength steel sheet is characterized in that the dislocation arrangement parameter M in the quenched martensite is 2.0 or less, the C concentration of the prior austenite grain boundaries analyzed with a three-dimensional atom probe is 5.00 at% or less, the Mn concentration is 5.00 at% or less, and the P concentration is 2.00 at% or less, and the hydrogen content in the steel after low-temperature aging treatment at 170°C for 20 minutes is 0.20 wt. ppm or less.

2. The above component composition is further defined in mass percent as follows: Ti: 0.200% or less, Nb: 0.200% or less, V: 0.200% or less, Ta: 0.100% or less, W: 0.100% or less, B: 0.0100% or less, Ni: 1.00% or less, Cr: 1.00% or less, Mo: 1.00% or less, Cu: 1.00% or less, Co: 0.500% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.500% or less, Te: 0.100% or less. The high-strength steel sheet according to claim 1, containing at least one element selected from the group consisting of Hf: 0.100% or less and Bi: 0.100% or less.

3. A hot-dip galvanized steel sheet having a high-strength steel sheet according to claim 1 or 2, and a hot-dip galvanized layer formed on the surface of the high-strength steel sheet.

4. The hot-dip galvanized steel sheet according to claim 3, wherein the hot-dip galvanized layer is an alloyed hot-dip galvanized layer.

5. A step of heating a steel slab having the component composition described in claim 1 or 2, and hot-rolling the steel slab with a finish rolling exit temperature of 750°C or more and 1000°C or less to obtain a hot-rolled steel sheet; a step of winding the hot-rolled steel sheet at 300°C or more and 750°C or less; a step of cold-rolling the hot-rolled steel sheet with a rolling ratio of 60.0% or less to obtain a cold-rolled steel sheet; and a step of Ac ３ An annealing process in which the cold-rolled steel sheet is heated to a temperature above the transformation point and below 950°C, and then the cold-rolled steel sheet is subjected to Ac ３ A method for manufacturing a high-strength steel sheet, comprising: a cooling step of cooling under the condition that the average cooling rate from the transformation point to 500°C is 2.0°C / s or more; and a strain application step of performing one or both of the following: (I) passing the cold-rolled steel sheet through a skin pass mill having a pair of upper and lower rolls, and (II) passing the cold-rolled steel sheet through a tension leveler having upper and lower rolls arranged alternately; wherein when (I) is performed in the strain application step, the roll diameter of the upper and lower rolls is 1000 mm or less, the temperature of the cold-rolled steel sheet at the entry side of the skin pass mill is 200°C or less, and the elongation rate is 0.01% or more; and when (II) is performed in the strain application step, the roll diameter of the upper and lower rolls is 100 mm or less, the temperature of the cold-rolled steel sheet at the entry side of the tension leveler is 100°C or less, and the value obtained by dividing the intermesh indentation amount by the sheet thickness is 1.3 or more.

6. A method for manufacturing a hot-dip galvanized steel sheet, comprising the steps of the method for manufacturing a high-strength steel sheet described in claim 5, and a step of applying a hot-dip galvanizing treatment to the cold-rolled steel sheet after the cooling step and before the strain application step to obtain a hot-dip galvanized steel sheet, wherein the strain application step is performed on the hot-dip galvanized steel sheet instead of the cold-rolled steel sheet.

7. A method for manufacturing a hot-dip galvanized steel sheet according to claim 6, comprising the step of applying an alloying treatment to the hot-dip galvanized steel sheet after the hot-dip galvanizing treatment and before the strain application step.

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