WO2023008002A1

WO2023008002A1 - Steel sheet, member, and methods for producing said steel sheet and said member

Info

Publication number: WO2023008002A1
Application number: PCT/JP2022/024963
Authority: WO
Inventors: 大洋浅川; 真平吉岡; 真次郎金子
Original assignee: Ｊｆｅスチール株式会社
Priority date: 2021-07-28
Filing date: 2022-06-22
Publication date: 2023-02-02
Also published as: JPWO2023008002A1; EP4350015A1; KR20240024946A; CN117677725A; JP7226673B1

Abstract

Provided is a steel sheet which has high strength and excellent delayed fracture resistance. Also provided is a method for producing this steel sheet.　The steel sheet includes, in terms of mass%, C: 0.15% to 0.45%; Si: 1.5% or less; Mn: 1.7% or less; P: 0.03% or less; S: less than 0.0020%; sol. Al: 0.20% or less; N: 0.005% or less; B: 0.0015% to 0.0100%; Nb and/or Ti: 0.005% to 0.080% in total, and the remainder being Fe and unavoidable impurities. The area ratio of martensite with respect to the entire structure is 95% to 100%; the prior γ particle size is less than 11.0 µm; and the number density A of the precipitate having an equivalent circular diameter of 500 nm or greater satisfies the mathematical expression of A (pcs./mm2) ≤ 8.5 × 105 × [B].

Description

Steel plate, member and manufacturing method thereof

The present invention relates to steel sheets such as high-strength steel sheets for cold press forming that are used in automobiles and the like through cold press forming, members using the steel sheets, and manufacturing methods thereof.

In recent years, steel sheets with a tensile strength TS of 1310 MPa class or higher have been increasingly applied to automobile frame parts for the purpose of weight reduction and collision safety of automobiles. In addition, steel sheets with a tensile strength TS of 1470 MPa class or higher are being applied to bumpers, impact beam parts, and the like.

When a high-strength steel sheet with a tensile strength TS of 1470 MPa class or higher is formed into a part by cold pressing, there is a risk of delayed fracture due to an increase in residual stress in the part and deterioration of the delayed fracture resistance due to the steel itself. There is

Here, delayed fracture means that when a part is placed in a hydrogen penetration environment with high stress applied to the part, hydrogen penetrates into the steel plate that constitutes the part, reducing the interatomic bonding strength. It is a phenomenon in which microcracks are generated by causing local deformation and breakage occurs as the microcracks propagate.

As a technique for improving such delayed fracture characteristics, for example, based on the knowledge that delayed fracture resistance is improved by reducing coarse precipitates that are the starting point of delayed fracture fracture, Patent Document 1 discloses mass % , C: 0.13% or more and 0.40% or less, Si: 0.02% or more and 1.5% or less, Mn: 0.4% or more and 1.7% or less, P: 0.030% or less, S : 0.0002% or more and less than 0.0010%, sol. Al: 0.01% or more and 0.20% or less, N: 0.0055% or less, O: 0.0025% or less, Nb: 0.002% or more and 0.035% or less, and Ti: 0.002% or more and 0 .040% or less so as to satisfy the formulas (1) and (2), the balance being Fe and unavoidable impurities, and the total area ratio of martensite and bainite to the entire structure is 95% or more. 100% or less, the balance consists of one or two types of ferrite and retained austenite, the average grain size of prior austenite grains exceeds 5 μm, the following conditions are satisfied, and the length of the major axis is 20 to 80 μm and a steel structure in which inclusion groups are present at 5/mm ² or less, and a tensile strength of 1,320 MPa or more.
[% Ti] + [% Nb] > 0.007 (1)
[% Ti] × [% Nb] ² ≤ 7.5 × 10 ^-6 (2)
Here, [%Nb] and [%Ti] represent the contents (%) of Nb and Ti.

In addition, in Patent Document 2, in mass%, C: 0.05% to 0.30%, Si: 2.0% or less (including 0%), Mn: more than 0.1% and 2.8% or less , P: 0.1% or less, S: 0.005% or less, N: 0.01% or less, Al: 0.01 to 0.50% or less, and one or two of Nb, Ti and Zr at least 0.01% in total, and [%C]-[%Nb]/92.9×12-[%Ti]/47.9×12-[%Zr]/91.2 × 12 > 0.03, the balance has a component composition consisting of iron and unavoidable impurities, tempered martensite contains 50% or more (including 100%) in terms of area ratio, and the balance is ferrite The distribution of precipitates in tempered martensite is 20 or more per 1 μm ² of tempered martensite with an equivalent circle diameter of 1 to 10 nm, and the precipitates have an equivalent circle diameter of 20 nm or more. There are 10 or less precipitates per 1 μm ² of tempered martensite containing one or more of Nb, Ti and Zr, and the ferrite surrounded by large-angle grain boundaries with a crystal orientation difference of 15 ° or more. A high-strength cold-rolled steel sheet having an average grain size of 5 μm or less and having excellent hydrogen embrittlement resistance and workability is disclosed.

Japanese Patent No. 6388085 Japanese Patent No. 4712882

However, it could not be said that the conventional technology is sufficient as a technology that secures a tensile strength TS of 1470 MPa or more and has excellent delayed fracture resistance.

The present invention has been made to solve such problems, and provides a steel sheet, a member, and a method for producing the same having a tensile strength of 1470 MPa or more (TS≧1470 MPa) and excellent delayed fracture resistance. intended to

The term "excellent delayed fracture resistance" means judging to have excellent delayed fracture resistance by the following evaluation.
(1) First, a strip test piece of 100 mm in the direction perpendicular to the rolling direction and 30 mm in the rolling direction is taken from the 1/4 position of the coil width from the end of the obtained steel plate (coil) in the width direction.
(2) Cut out the end face of the long side with a length of 100 mm by shearing, and bend it so that the burr is on the outer peripheral side while being sheared (without machining to remove burrs). After processing, the test piece is fixed with bolts while maintaining the shape of the test piece at the time of bending.
The shearing clearance was 13% and the rake angle was 1°. Bending is performed with a tip bending radius of 10 mm and an angle inside the bending apex of 90 degrees (V bending).
The punch has a U-shaped punch whose tip radius is the same as the tip bending radius R (the tip R is semicircular and the thickness of the punch barrel is 2R), and the die has a corner radius of 30 mm. Use Then, the depth to which the punch pushes the steel plate is adjusted, and the steel plate is formed so that the bending angle of the tip (the angle inside the bending apex) is 90 degrees (V shape).
Sandwich the test piece with a hydraulic jack so that the distance between the flange ends of the straight piece when bending is the same as when bending (to cancel out the opening of the straight piece due to springback). and tighten the bolts in that state. The bolt is passed through an elliptical hole (minor axis: 10 mm, major axis: 15 mm) previously provided 10 mm inward from the short side edge of the strip test piece and fixed.
(3) The resulting bolted test piece was immersed in a solution prepared by mixing 0.1% by mass of ammonium thiocyanate aqueous solution and McIlvaine buffer solution at a ratio of 1:1 and adjusting the pH to 8.0. A delayed fracture resistance evaluation test is carried out. At this time, the temperature of the solution shall be 20° C., and the volume of the solution per 1 cm ³ of the surface area of the test piece shall be 20 ml.
(4) After 24 hours have passed, the presence or absence of cracks at a level that can be visually confirmed (1 mm or more in length) is confirmed, and if no cracks are observed, it is judged that the delayed fracture resistance is excellent.

The inventors of the present invention conducted extensive studies to solve the above problems, and found that the delayed fracture resistance can be greatly improved by satisfying all of the following conditions.
i) The area ratio of martensite is 95% or more.
ii) The average grain size (prior γ grain size) of prior austenite grains is less than 11.0 µm.
iii) The number density A of precipitates having an equivalent circle diameter of 500 nm or more satisfies the following conditions.
A (pieces/mm ² ) ≤ 8.5 × 10 ⁵ × [B]
Here, [B] represents the content of B (% by mass).

The present invention was completed through further studies based on the above findings, and the gist thereof is as follows.
[1] % by mass,
C: 0.15% or more and 0.45% or less,
Si: 1.5% or less,
Mn: 1.7% or less,
P: 0.03% or less,
S: less than 0.0020%,
sol. Al: 0.20% or less,
N: 0.005% or less,
B: 0.0015% or more and 0.0100% or less,
Containing 0.005% or more and 0.080% or less in total of one or more of Nb and Ti,
Having a component composition in which the balance is Fe and unavoidable impurities,
Having a structure in which martensite accounts for 95% or more and 100% or less of the area of the entire structure,
The average grain size of the prior austenite grains is less than 11.0 μm,
A steel sheet in which the number density A of precipitates having an equivalent circle diameter of 500 nm or more satisfies the following formula (1).
A (number/mm ² )≦8.5×10 ⁵ ×[B] Expression (1)
Here, [B] represents the content of B (% by mass).
[2] The steel sheet according to [1], which further contains, as the chemical composition, one or two selected from Cu: 1.0% or less and Ni: 1.0% or less in mass%.
[3] As the component composition, in terms of mass %, Cr: 1.0% or less, Mo: less than 0.3%, V: 0.5% or less, Zr: 0.2% or less, and W: 0.2 The steel sheet according to [1] or [2], containing one or more selected from % or less.
[4] As the component composition, in mass%, Ca: 0.0030% or less, Ce: 0.0030% or less, La: 0.0030% or less, REM (excluding Ce and La): 0.0030% The steel sheet according to any one of [1] to [3], containing one or more selected from the following and Mg: 0.0030% or less.
[5] The component composition of [1] to [4] further contains one or two selected from Sb: 0.1% or less and Sn: 0.1% or less in mass%. The steel plate according to any one of the above.
[6] The steel sheet according to any one of [1] to [5], which has a plating layer on the surface of the steel sheet.
[7] A member using the steel plate according to any one of [1] to [6].
[8] A steel slab having the chemical composition according to any one of [1] to [5] is heated at a slab surface temperature from 1000 ° C. to a heating and holding temperature of 1250 ° C. or higher at an average heating rate of 10 ° C./min or less. After heating and holding at the heating and holding temperature for 30 minutes or more,
Hot finish rolling is performed under the conditions of a residence time of 20 seconds or more and 150 seconds or less at 900 to 1000 ° C. and a finish rolling temperature of 850 ° C. or more,
Cooling is performed at an average cooling rate of 40 ° C./sec or more in the range from the finish rolling temperature to 650 ° C.,
After that, it is coiled at a coiling temperature of 650 ° C. or less to form a hot rolled steel sheet,
A cold-rolled steel sheet is obtained by cold-rolling the hot-rolled steel sheet at a rolling reduction of 40% or more,
The annealing temperature is 830 to 950° C., and the cold-rolled steel sheet is heated from 400° C. to the annealing temperature at an average heating rate of 1.0° C./second or more,
Hold at the annealing temperature for 600 seconds or less,
Cooling from a cooling start temperature of 680 ° C. or higher to a cooling stop temperature of 260 ° C. or lower at an average cooling rate of 70 ° C./sec or higher,
A method for manufacturing a steel sheet, followed by continuous annealing at a holding temperature of 150 to 260° C. for 20 to 1500 seconds.
[9] The method for producing a steel sheet according to [8], wherein the surface of the steel sheet is plated after the continuous annealing.
[10] A method for manufacturing a member, comprising the step of subjecting the steel plate according to any one of [1] to [6] to at least one of forming and joining to form a member.

　According to the present invention, a high-strength steel plate, a member, and a method for manufacturing the same are provided, which are excellent in delayed fracture resistance.

Embodiments of the present invention will be described below.

The steel sheet of the present invention, in mass%, C: 0.15% or more and 0.45% or less, Si: 1.5% or less, Mn: 1.7% or less, P: 0.03% or less, S: 0 less than .0020%, sol. Al: 0.20% or less, N: 0.005% or less, B: 0.0015% or more and 0.0100% or less, 0.005% or more and 0.080% or less in total of one or more of Nb and Ti with the balance being Fe and unavoidable impurities, having a structure in which the area ratio of martensite to the entire structure is 95% or more and 100% or less, and prior austenite grains (hereinafter also referred to as prior γ grains ) has an average grain size (prior γ grain size) of less than 11.0 μm, and the number density A of precipitates having an equivalent circle diameter of 500 nm or more satisfies the following formula (1).
A (pieces/mm ² )≦8.5×10 ⁵ ×[B] Expression (1)
Here, [B] represents the content of B (% by mass).

Composition The reason for limiting the range of the composition of the steel sheet of the present invention will be described below. In addition, % regarding component content is "mass %."

C: 0.15% or more and 0.45% or less C is contained in order to improve the hardenability to obtain a martensite steel structure and to increase the strength of martensite. In order to ensure a tensile strength of 1470 MPa or more (hereinafter also referred to as TS≧1470 MPa), the C content is made 0.15% or more. The C content is preferably 0.20% or more, more preferably 0.27% or more, from the viewpoint of reducing the weight of automobile frame parts by increasing the tensile strength. On the other hand, excessively added C causes deterioration of delayed fracture resistance due to formation of iron carbide and segregation to grain boundaries. From these points of view, the C content is limited to a range of 0.45% or less. The C content is preferably 0.40% or less, more preferably 0.37% or less.

Si: 1.5% or less Si is used as a strengthening element by solid solution strengthening, and from the viewpoint of suppressing the formation of film-like carbides when tempering in a temperature range of 200 ° C. or higher to improve delayed fracture resistance. contains. In addition, Si is contained from the viewpoint of reducing Mn segregation at the central portion of the sheet thickness and suppressing the formation of MnS. Furthermore, Si is contained in order to suppress decarburization and deboronation due to oxidation of the surface layer during annealing in a continuous annealing line (CAL). Although the lower limit of the Si content is not specified, it is desirable to contain 0.02% or more of Si from the viewpoint of obtaining the above effect. The Si content is preferably 0.10% or more, more preferably 0.20% or more. On the other hand, if the Si content is too high, a significant increase in rolling load during hot rolling and cold rolling and a drop in toughness are caused. From these points of view, the Si content should be 1.5% or less (including 0%). The Si content is preferably 1.2% or less, more preferably 1.0% or less.

Mn: 1.7% or less Mn is contained in order to improve the hardenability of steel and obtain the desired strength, and to keep the area ratio of martensite within a predetermined range. On the other hand, if the Mn content is too high, the segregation of Mn increases, degrading workability and weldability. Therefore, Mn should be 1.7% or less. The Mn content is preferably 1.5% or less, more preferably 1.3% or less. Although the lower limit of the Mn content is not specified, it is preferable to contain 0.2% or more of Mn in order to industrially stably secure a predetermined area ratio of martensite.

P: 0.03% or less P is an element that strengthens steel, but when its content is high, it segregates at grain boundaries and lowers grain boundary strength, resulting in significant deterioration in delayed fracture resistance and spot weldability. Invite. From the above point of view, the P content should be 0.03% or less. The P content is preferably 0.02% or less, more preferably 0.01% or less. Although the lower limit of the P content is not specified, it is set to 0.002% as the lower limit currently industrially practicable.

S: less than 0.0020% S forms coarse MnS and acts as a starting point for delayed fracture, thereby greatly deteriorating the resistance to delayed fracture. Therefore, the S content should be at least less than 0.0020% to reduce MnS. From the viewpoint of improving delayed fracture resistance, the S content is preferably less than 0.0010%, more preferably 0.0008% or less, and even more preferably 0.0006% or less. Although the lower limit is not defined, it is set to 0.0002% as the lower limit currently industrially practicable.

sol. Al: 0.20% or less Al is contained in order to sufficiently deoxidize and reduce inclusions in the steel. sol. Although the lower limit of Al is not specified, in order to stably deoxidize, sol. It is desirable to set the Al content to 0.005% or more. sol. The Al content is more preferably 0.01% or more, still more preferably 0.02% or more. On the other hand, sol. If the Al content exceeds 0.20%, the cementite generated during winding becomes less likely to form a solid solution during the annealing process, and the delayed fracture resistance deteriorates. Therefore, sol. Al content is 0.20% or less. sol. The Al content is preferably 0.10% or less, more preferably 0.05% or less.

N: 0.005% or less N forms precipitates such as TiN, (Nb, Ti) (C, N) in the steel. , (Nb,Ti)C. These hinder adjustment to the steel structure required by the present invention, and adversely affect the delayed fracture resistance. In order to reduce such adverse effects, the N content is made 0.005% or less. The N content is preferably 0.0040% or less. Although the lower limit is not defined, it is set to 0.0006% as the lower limit currently industrially practicable.

B: 0.0015% or more and 0.0100% or less B is an element that improves the hardenability of steel, and has the advantage of forming martensite with a predetermined area ratio even with a small Mn content. In addition, B increases the cohesive force of the grain boundary by segregating at the grain boundary, and suppresses the segregation of P, which reduces the grain boundary strength, thereby improving the delayed fracture resistance. On the other hand, the addition of excessive B increases Fe ₂₃ (C, B) ₆ and BN, and it becomes a starting point of delayed fracture, resulting in a decrease in the resistance to delayed fracture. Therefore, in order to obtain the effect of improving the delayed fracture resistance by adding B, it is necessary to achieve both an increase in grain boundary solid solution B and a suppression of B-based precipitates. In steel with a prior γ grain size of 10 µm or less, the B content is set to 0.0015% or more in order to obtain a sufficient grain boundary solid solution B amount. The B content is preferably 0.0025% or more, more preferably 0.0040% or more. On the other hand, when the content of B exceeds 0.0100%, it becomes difficult to reduce B-based precipitates even by controlling the hot rolling conditions and annealing conditions. Therefore, the B content is set to 0.0100% or less. The B content is preferably 0.0090% or less, more preferably 0.0080% or less.

0.005% or more and 0.080% or less in total of one or more of Nb and Ti Improves delayed fracture resistance. From this point of view, one or more of Nb and Ti are contained in a total amount of 0.005% or more. The total content of Nb and Ti is preferably 0.010% or more, more preferably 0.020% or more. On the other hand, if the total content of one or more of Nb and Ti exceeds 0.080%, Nb and Ti do not completely dissolve in the slab reheating, and TiN, Ti(C,N), NbN, Nb(C, Precipitates having an equivalent circle diameter of 500 nm or more, such as N), (Nb, Ti), (C, N), increase and act as starting points for delayed fracture, rather degrading the delayed fracture resistance. Therefore, the upper limit of the total content of Nb and Ti is 0.080%. The total content of Nb and Ti (Ti+Nb) is preferably 0.07% or less, more preferably 0.06% or less.

The chemical composition of the steel sheet in the present invention contains the above constituent elements as basic components, and the balance includes iron (Fe) and inevitable impurities. Here, the steel sheet of the present invention preferably has a chemical composition containing the above-described basic components, with the balance being iron (Fe) and unavoidable impurities.

In the present invention, one or more selected from the following (A) to (D) may be contained as the component composition.
(A) in mass%, one or two selected from Cu: 1.0% or less and Ni: 1.0% or less,
(B) in mass%, selected from Cr: 1.0% or less, Mo: less than 0.3%, V: 0.5% or less, Zr: 0.2% or less, and W: 0.2% or less 1 or 2 or more
(C) In mass %, Ca: 0.0030% or less, Ce: 0.0030% or less, La: 0.0030% or less, REM (excluding Ce and La): 0.0030% or less, and Mg: 0.0030% or less. 1 or 2 or more selected from 0030% or less,
(D) In mass%, one or two selected from Sb: 0.1% or less and Sn: 0.1% or less

Cu: 1.0% or less Cu improves corrosion resistance in the use environment of automobiles. In addition, the inclusion of Cu has the effect of suppressing penetration of hydrogen into the steel sheet by coating the surface of the steel sheet with corrosion products. Moreover, Cu is an element that is mixed when scrap is used as a raw material, and by allowing Cu to be mixed, recycled materials can be used as raw materials, and manufacturing costs can be reduced. From the above point of view, the content of Cu is preferably 0.01% or more, and from the point of view of improving the delayed fracture resistance, the content of Cu is preferably 0.05% or more. Cu content is more preferably 0.10% or more. However, if the Cu content is too high, surface defects may occur, so the Cu content is preferably 1.0% or less. From the above, when Cu is contained, the Cu content is set to 1.0% or less. The Cu content is more preferably 0.50% or less, still more preferably 0.30% or less.

Ni: 1.0% or less Ni is also an element that acts to improve corrosion resistance. In addition, Ni has the effect of reducing surface defects that tend to occur when Cu is contained. Therefore, from the above point of view, it is desirable to contain 0.01% or more of Ni. The Ni content is more preferably 0.05% or more, still more preferably 0.10% or more. However, if the Ni content is too high, scale formation in the heating furnace becomes non-uniform, causing surface defects and a significant increase in cost. Therefore, when Ni is contained, the Ni content is set to 1.0% or less. The Ni content is more preferably 0.50% or less, still more preferably 0.30% or less.

Cr: 1.0% or less Cr can be added to obtain the effect of improving the hardenability of steel. In order to obtain the effect, it is preferable to contain 0.01% or more of Cr. The Cr content is more preferably 0.05% or more, and still more preferably 0.10% or more. However, when the Cr content exceeds 1.0%, the cementite dissolution rate during annealing is delayed, and undissolved cementite remains, thereby deteriorating the delayed fracture resistance of the sheared end face. Moreover, pitting corrosion resistance is also deteriorated. Furthermore, it also degrades chemical convertibility. Therefore, when Cr is contained, the Cr content is made 1.0% or less. Delayed fracture resistance, pitting corrosion resistance, and chemical conversion treatability all tend to start to deteriorate when the Cr content exceeds 0.2%, so from the viewpoint of preventing these, the Cr content is 0.2% or less. is more preferable.

Mo: less than 0.3% Mo has the effect of improving the hardenability of steel, the effect of generating fine carbides containing Mo that serve as hydrogen trap sites, and the improvement of delayed fracture resistance by refining martensite. can be added for the purpose of obtaining the effect of If a large amount of Nb or Ti is added, these coarse precipitates are formed and the delayed fracture resistance deteriorates, but the solid solubility limit of Mo is larger than that of Nb and Ti. When added in combination with Nb and Ti, it forms fine precipitates in which these and Mo are combined, and has the effect of refining the structure. Therefore, by adding Mo in addition to small amounts of Nb and Ti, it is possible to refine the structure without leaving coarse precipitates and disperse a large amount of fine carbides, thereby improving delayed fracture resistance. can be improved. In order to obtain the effect, it is desirable to contain Mo at 0.01% or more. The Mo content is more preferably 0.03% or more, still more preferably 0.05% or more. However, when Mo is contained in an amount of 0.3% or more, the chemical conversion treatability deteriorates. Therefore, when Mo is contained, the Mo content should be less than 0.3%. The Mo content is preferably 0.2% or less.

V: 0.5% or less V has the effect of improving the hardenability of steel, the effect of forming fine carbides containing V that serve as hydrogen trap sites, and the improvement of delayed fracture resistance by refining martensite. It can be added for the purpose of obtaining an effect. In order to obtain the effect, it is desirable to set the V content to 0.003% or more. The V content is more preferably 0.03% or more, still more preferably 0.05% or more. However, if the V content exceeds 0.5%, the castability is remarkably deteriorated. Therefore, when V is contained, the V content should be 0.5% or less. The V content is more preferably 0.3% or less, still more preferably 0.2% or less. Further, the V content is preferably 0.1% or less.

Zr: 0.2% or less Zr contributes to high strength through refinement of prior γ grains and the resulting refinement of the internal structure of martensite, and improves delayed fracture resistance. In addition, through the formation of fine Zr-based carbides/carbonitrides that serve as hydrogen trapping sites, the strength is increased and the delayed fracture resistance is improved. Zr also improves castability. From this point of view, the Zr content is desirably 0.005% or more. The Zr content is more preferably 0.010% or more, preferably 0.015% or more. However, if a large amount of Zr is added, coarse precipitates of ZrN and ZrS that remain undissolved when the slab is heated in the hot rolling process increase, degrading the delayed fracture resistance of the sheared edge. Therefore, when Zr is contained, the Zr content should be 0.2% or less. The Zr content is more preferably 0.1% or less, still more preferably 0.04% or less.

W: 0.2% or less W contributes to increasing the strength and improving the delayed fracture resistance through the formation of fine W-based carbides and carbonitrides that serve as hydrogen trap sites. From this point of view, it is desirable to contain W at 0.005% or more. The W content is more preferably 0.010% or more, still more preferably 0.030% or more. However, if a large amount of W is contained, coarse precipitates remaining undissolved when the slab is heated in the hot rolling process increase, and the delayed fracture resistance of the sheared end surface deteriorates. Therefore, when W is contained, the W content should be 0.2% or less. The W content is more preferably 0.1% or less.

Ca: 0.0030% or less Ca fixes S as CaS and improves delayed fracture resistance. In order to obtain this effect, it is desirable to contain 0.0002% or more of Ca. The Ca content is more preferably 0.0005% or more, still more preferably 0.0010% or more. However, since adding a large amount of Ca deteriorates the surface quality and bendability, the Ca content is preferably 0.0030% or less. As mentioned above, when containing Ca, Ca content shall be 0.0030% or less. The Ca content is more preferably 0.0025% or less, still more preferably 0.0020% or less.

Ce: 0.0030% or less Ce also fixes S and improves delayed fracture resistance. In order to obtain this effect, it is desirable to contain Ce at 0.0002% or more. The Ce content is more preferably 0.0003% or more, and still more preferably 0.0005% or more. However, since adding a large amount of Ce deteriorates the surface quality and bendability, the Ce content is preferably 0.0030% or less. From the above, when Ce is contained, the Ce content is set to 0.0030% or less. The Ce content is more preferably 0.0020% or less, still more preferably 0.0015% or less.

La: 0.0030% or less La also fixes S and improves delayed fracture resistance. In order to obtain this effect, it is desirable to contain 0.0002% or more of La. The La content is more preferably 0.0005% or more, still more preferably 0.0010% or more. However, since adding a large amount of La deteriorates the surface quality and bendability, the La content is preferably 0.0030% or less. From the above, when La is contained, the La content shall be 0.0030% or less. The La content is more preferably 0.0020% or less, still more preferably 0.0015% or less.

REM: 0.0030% or less REM also fixes S and improves delayed fracture resistance. In order to obtain this effect, it is desirable to contain 0.0002% or more of REM. The REM content is more preferably 0.0003% or more, still more preferably 0.0005% or more. However, since addition of a large amount of REM degrades the surface quality and bendability, the REM content is preferably 0.0030% or less. From the above, when REM is contained, the REM content is set to 0.0030% or less. The REM content is more preferably 0.0020% or less, still more preferably 0.0015% or less.
In the present invention, REM includes scandium (Sc) with atomic number 21, yttrium (Y) with atomic number 39, and lanthanum (La) with atomic number 57 to lutetium (Lu) with atomic number 71. Among the lanthanoids, it refers to elements other than Ce and La. The REM concentration in the present invention is the total content of one or more elements selected from the above REMs.

Mg: 0.0030% or less Mg fixes O as MgO and improves delayed fracture resistance. In order to obtain this effect, it is desirable to contain 0.0002% or more of Mg. The Mg content is more preferably 0.0005% or more, still more preferably 0.0010% or more. However, since adding a large amount of Mg deteriorates the surface quality and bendability, the Mg content is preferably 0.0030% or less. From the above, when Mg is contained, the Mg content should be 0.0030% or less. The Mg content is more preferably 0.0020% or less, still more preferably 0.0015% or less.

Sb: 0.1% or less Sb suppresses the oxidation and nitridation of the surface layer, thereby suppressing the reduction of C and B. Suppressing the reduction of C and B suppresses the formation of ferrite in the surface layer, contributing to higher strength and improved delayed fracture resistance. From this point of view, the Sb content is desirably 0.002% or more. The Sb content is more preferably 0.004% or more, still more preferably 0.006% or more. However, if the Sb content exceeds 0.1%, the castability deteriorates, and Sb segregates at the prior γ grain boundaries to deteriorate the delayed fracture resistance of the sheared edge. Therefore, the Sb content is desirably 0.1% or less. From the above, when Sb is contained, the Sb content is set to 0.1% or less. The Sb content is more preferably 0.05% or less, still more preferably 0.02% or less.

Sn: 0.1% or less Sn suppresses oxidation and nitridation of the surface layer, thereby suppressing a decrease in the content of C and B in the surface layer. Suppressing the reduction of C and B suppresses the formation of ferrite in the surface layer, contributing to higher strength and improved delayed fracture resistance. From this point of view, the Sn content is desirably 0.002% or more. The Sn content is preferably 0.003% or more. However, if the Sn content exceeds 0.1%, the castability deteriorates, and Sn segregates at the prior γ grain boundaries, resulting in deterioration of the delayed fracture resistance of the sheared edges. Therefore, when Sn is contained, the Sn content is set to 0.1% or less. The Sn content is more preferably 0.05% or less, still more preferably 0.01% or less.

When the content of the arbitrary element is less than the preferred lower limit, the arbitrary element is included as an unavoidable impurity.

Steel Structure The steel structure of the steel sheet of the present invention has the following structure.
(Configuration 1) The area ratio of martensite to the entire structure is 95% or more and 100% or less.
(Configuration 2) The average grain size of prior austenite grains is less than 11.0 μm.
(Structure 3) The number density A of precipitates having an equivalent circle diameter of 500 nm or more satisfies the following formula (1).
A (number/mm ² )≦8.5×10 ⁵ ×[B] Expression (1)
Here, [B] represents the content of B (% by mass).

Each configuration will be described below.

(Configuration 1) The area ratio of martensite to the entire structure is 95% or more and 100% or less.
In order to achieve both high strength of TS≧1470 MPa and excellent delayed fracture resistance, the area ratio of martensite in the steel structure is set to 95% or more. More preferably 99% or more, still more preferably 100%. In addition, when other than martensite is included, the balance includes bainite, ferrite, and retained austenite (retained γ). Other than these structures are trace amounts of carbides, sulfides, nitrides and oxides. The residual tissue is 5% or less, preferably 1% or less. In addition, martensite includes self-tempering during continuous cooling and martensite that is not tempered by staying at about 150° C. or higher for a certain period of time. Note that the area ratio of martensite may be 100% without including the remainder.

(Configuration 2) The average grain size of prior austenite grains is less than 11.0 μm.
In steels with a martensite area ratio of 95% or more in the steel structure, the delayed fracture surface often exhibits intergranular fracture surface, and the initiation point of delayed fracture and the crack propagation path at the initial stage of delayed fracture are on the prior austenite grain boundary. Conceivable. Refinement of prior-austenite grains is effective for suppressing intergranular fracture, and refinement of prior-austenite grains significantly improves delayed fracture resistance. The mechanism is thought to be that the area ratio of the prior austenite grain boundaries increases due to the refinement of the prior austenite grains, and the concentration of impurity elements such as P, which is a grain boundary embrittlement element, on the prior austenite grain boundaries decreases. Further, the refinement of the prior austenite grains also contributes to the improvement of the tensile strength. From the viewpoint of delayed fracture resistance and strength, the average grain size of prior austenite grains (prior γ grain size) is less than 11.0 μm. This average particle size is preferably 10 μm or less, more preferably 7.0 μm or less, and still more preferably 5.0 μm or less.

(Structure 3) The number density A of precipitates having an equivalent circle diameter of 500 nm or more satisfies the following formula.
A (pieces/mm ² )≦8.5×10 ⁵ ×[B] Expression (1)
Here, [B] represents the content of B (% by mass).
In order to suppress intergranular fracture in steel having a high strength of TS≧1470 MPa, in addition to refining prior austenite grains, it is effective to segregate B at the grain boundaries to strengthen the grain boundaries. However, a mere increase in the amount of B added increases not only the grain boundary segregation B but also the B-based precipitates mainly composed of Fe ₂₃ (C, B) ₆ which is the starting point of delayed fracture, so the delayed fracture resistance is rather improved. Lower. The present inventors have found that by controlling the hot rolling conditions, the number density A of precipitates with an equivalent circle diameter of 500 nm or more is reduced, and by satisfying the following conditions, the delayed fracture resistance is improved by strengthening the grain boundaries of B and the origin of the precipitates It was found that it is possible to simultaneously suppress the destruction of the
A (pieces/mm ² ) ≤ 8.5 × 10 ⁵ × [B]
Preferably, A (number/mm ² )≦5.0×10 ⁵ ×[B], more preferably A (number/mm ² )≦2.0×10 ⁵ ×[B].

A method of measuring each component in the above steel structure will be described.
The area ratios of martensite, bainite, and ferrite are obtained by polishing the L section of the steel sheet (the section parallel to the rolling direction and perpendicular to the surface of the steel sheet (hereinafter also referred to as the vertical section parallel to the rolling direction)) with Nital. Corroded and observed at 1/4 thickness position from the surface of the steel plate with SEM at a magnification of 2000 times, 4 fields of view are observed, and the photographed structure photograph is image-analyzed and measured. Here, martensite and bainite indicate gray or white structures in SEM. On the other hand, ferrite is a region exhibiting black contrast in SEM. Martensite and bainite contain trace amounts of carbides, nitrides, sulfides, and oxides, but since it is difficult to exclude them, the area ratio of the region including these is used as the area ratio.

Here, bainite has the following characteristics. That is, it has an aspect ratio of 2.5 or more, exhibits a plate-like form, and is a slightly blacker structure than martensite. The width of the plate is 0.3-1.7 μm. The distribution density of carbides with a diameter of 10 to 200 nm inside bainite is 0 to 3 pieces/μm ² .

　The retained austenite (retained γ) is measured by chemically polishing the 200 μm surface layer of the steel plate with oxalic acid, and using the X-ray diffraction intensity method for the plate surface. It is calculated from the integrated intensities of (200)α, (211)α, (220)α, (200)γ, (220)γ, and (311)γ diffraction plane peaks measured by Mo-Kα rays.

The average grain size of prior austenite grains (prior γ grain size) is measured by polishing the L cross section (vertical cross section parallel to the rolling direction) of the steel sheet, and then applying a chemical solution that corrodes the prior γ grain boundaries (such as a saturated picric acid aqueous solution or It was corroded by adding ferric chloride to this), and observed at 1/4 thickness position from the steel plate surface with an optical microscope at a magnification of 500 times. 15 lines are drawn in each direction at intervals of 10 μm or more in actual length, and the number of intersections between grain boundaries and lines is counted. Furthermore, the prior γ grain size (average grain size of prior austenite grains) can be measured by multiplying the value obtained by dividing the total line length by the number of intersections by 1.13.

The number density A of precipitates with an equivalent circle diameter of 500 nm or more is obtained by polishing the L cross section (vertical cross section parallel to the rolling direction) of the steel sheet, and then the area from the 1/5 position to the 4/5 position of the steel plate thickness, that is, from the steel plate surface. In the region from the 1/5th position to the 4/5th position across the center of the plate thickness, an area of 2 mm ² was continuously photographed with an SEM, and from the photographed SEM photographs, such precipitates It was obtained by counting the number of Also, the magnification for photographing is 2000 times. Moreover, when performing the component analysis of each inclusion particle, each inclusion particle is magnified 10000 times, and the said precipitate is analyzed. Here, the precipitate having an equivalent circle diameter of 500 nm or more is a precipitate containing B such as Fe ₂₃ (C, B) ₆ , and an elemental analysis by energy dispersive X-ray spectroscopy (EDS) at an acceleration voltage of 3 kV The presence or absence of a B peak was examined, and when there was a B peak, it was evaluated that the above precipitates were present.
If the slab is not sufficiently reheated, precipitates containing Nb and Ti also increase, and these precipitates also adversely affect the delayed fracture properties.
The equivalent circle diameter refers to the diameter of a perfect circle having an area of each precipitate calculated from the SEM photograph.

Tensile strength (TS): 1470 MPa or more Deterioration of delayed fracture resistance becomes conspicuous when the steel sheet has a tensile strength of 1470 MPa or more. One of the characteristics of the present invention is that the delayed fracture resistance is good even if the tensile strength is 1470 MPa or more. Therefore, in the present invention, the tensile strength must be 1470 MPa or more. From the viewpoint of reducing the weight of automobile frame parts, it is preferably 1700 MPa or more. The tensile strength of the steel sheet of the present invention may be 2100 MPa or less.

The tensile strength can be measured by cutting out a JIS No. 5 tensile test piece so that the direction perpendicular to the rolling direction is the longitudinal direction at the position of 1/4 of the coil width, and conducting a tensile test based on JIS Z2241.

The steel sheet of the present invention described above may be a steel sheet having a plating layer on its surface. The plating layer may be Zn plating or plating of other metals. Further, it may be either a hot-dip plated layer or an electroplated layer.

Next, the method for manufacturing the steel sheet of the present invention will be described.
In the method for producing a steel plate of the present invention, a steel slab having the above chemical composition is heated at a slab surface temperature from 1000 ° C. to a heating and holding temperature of 1250 ° C. or higher at an average heating rate of 10 ° C./min or less, and this heating and holding After holding at the temperature for 30 minutes or more, hot finish rolling is performed under the conditions of a residence time of 20 seconds or more and 150 seconds or less at 900 to 1000 ° C. and a finish rolling temperature of 850 ° C. or more, and the finish rolling temperature is 650 ° C. Cooling is performed at an average cooling rate of 40 ° C./sec or more in the range of up to and then coiling at a coiling temperature of 650 ° C. or less to form a hot-rolled steel sheet, and the hot-rolled steel sheet is rolled at a reduction rate of 40% or more. A cold-rolled steel sheet is obtained by cold rolling, the annealing temperature is set to 830 to 950 ° C., the cold-rolled steel plate is heated from 400 ° C. to the annealing temperature at an average heating rate of 1.0 ° C./sec or more, and the annealing is performed. Hold the temperature for 600 seconds or less, cool from the cooling start temperature of 680 ° C. or more to the cooling stop temperature of 260 ° C. or less at an average cooling rate of 70 ° C./sec or more, and then keep the temperature at 150 to 260 ° C. for 20 to 1500. A steel sheet manufacturing method in which continuous annealing is performed for seconds.

Hot rolling In the slab heating before hot rolling, the average heating rate is set at 10°C/min or less from 1000°C to the heating and holding temperature of 1250°C or higher to promote the dissolution of sulfides and the formation of inclusions. Reduction in size and number is achieved. Since Nb and Ti have high dissolution temperatures, the heating and holding temperature at the slab surface temperature is set to 1250° C. or higher, and the holding time is set to 30 minutes or longer to promote solid dissolution of Nb and Ti. reduction is achieved. The heating and holding temperature is preferably 1300° C. or higher. More preferably, it is 1350° C. or higher.
Here, the average heating rate is "(Temperature at the completion of slab heating (heating holding temperature) (°C) - Temperature at the start of slab heating (°C) (1000°C)) / Heating time from the start of heating to the completion of heating. (minutes)”.

After that, the slab is held at 900-1000°C for 20 seconds or more and 150 seconds or less. An increase in the residence time in the temperature range of 900 to 1000° C. produces and coarsens precipitates mainly composed of BN. Precipitates generated in these temperature ranges are difficult to form a solid solution by annealing heating, and reduce the amount of solid solution B after annealing. Therefore, if the residence time exceeds 150 seconds, it is not possible to obtain a solid solution B amount that is effective in suppressing delayed fracture. Therefore, the residence time is 150 seconds or less, preferably 120 seconds or less, and more preferably 100 seconds or less. On the other hand, if the residence time is less than 20 seconds, the tissue may become non-uniform. Therefore, the residence time is 20 seconds or longer, preferably 30 seconds or longer, and more preferably 40 seconds or longer.

　In the hot finish rolling, the finish rolling temperature (FT) is set to 850°C or higher in order to suppress the precipitation of Nb, Ti, B, etc. Preferably, the finish rolling temperature is 930° C. or less.

Further, cooling after hot finish rolling is performed at an average cooling rate of 40°C/second or more in the range from the finish rolling temperature to 650°C. When the average cooling rate is less than 40° C./sec, the number of carbonitrides having an equivalent circle diameter of 1.0 μm or more increases due to the coarsening of Nb carbonitrides and Ti carbonitrides, and the desired delayed fracture resistance is obtained. can't Preferably, the average cooling rate is 250° C./s or less, more preferably 200° C./s or less.
The average cooling rate in the hot rolling process is "(temperature at the start of cooling (finish rolling temperature) (°C) - temperature at the completion of cooling (°C) (650°C)) / from the start of cooling to the completion of cooling. cooling time (seconds).

After cooling to 650°C, it is cooled and wound as necessary. At this time, if the coiling temperature exceeds 650° C., only the Nb and Ti precipitates precipitated in the fine austenite region are coarsened, so the coarse precipitates increase and the delayed fracture characteristics deteriorate. Therefore, the winding temperature should be 650° C. or lower. Preferably, the winding temperature is 500°C or higher.

Cold Rolling In cold rolling, if the rolling reduction (cold rolling rate) is 40% or more, recrystallization behavior and texture orientation in the subsequent continuous annealing can be stabilized. If the content is less than 40%, some of the austenite grains during annealing may become coarse and the strength may decrease. Also, the cold rolling rate is preferably 80% or less.

Continuous Annealing After cold rolling, the steel sheet is subjected to annealing and, if necessary, tempering and temper rolling in a continuous annealing line (CAL).
Since Fe ₂₃ (C, B) ₆ is generated in the ferrite region during annealing heating and coarsens, in order to reduce Fe ₂₃ (C, B) ₆ and sufficiently obtain the effect of grain boundary strengthening by B, It is very important to increase the average heating rate above 400°C. Also, from the viewpoint of refining the prior γ grain size to less than 11.0 μm, it is necessary to increase the heating rate. From the above point of view, the average heating rate at 400° C. or higher is 1.0° C./second or higher. Also, the average heating rate at 400° C. or higher is preferably 1.5° C./second or higher, more preferably 3.0° C./second or higher.
Also, preferably, the average heating rate is 10° C./sec or less.
Here, the average heating rate is defined as "annealing temperature (° C.)−400 (° C.) described later)/heating time (minutes) from 400° C. to the annealing temperature".
In order to sufficiently reduce precipitates such as Fe ₂₃ (C, B) ₆ remaining undissolved after annealing, annealing is performed at a high temperature for a long time. Specifically, the annealing temperature must be 830° C. or higher.
On the other hand, if the annealing temperature exceeds 950°C, the prior γ grain size becomes coarse and the desired structure cannot be obtained. Moreover, since BN may precipitate at grain boundaries and the delayed fracture resistance may deteriorate when annealing is performed at a temperature exceeding 900° C., the annealing temperature is preferably 900° C. or lower. Even if the soaking time (holding time) at the annealing temperature is prolonged, the prior γ grain size becomes too coarse, so the soaking time is set to 600 seconds or less. Preferably, this soaking time is 10 seconds or longer.

After that, in order to reduce ferrite and retained austenite and increase the area ratio of martensite to 95% or more, the average cooling rate is 70 ° C./sec or more from the cooling start temperature of 680 ° C. or higher to the cooling stop temperature of 260 ° C. or lower. Cool with

Here, the average cooling rate is defined as "cooling start temperature of 680 ° C. or higher (° C.) - cooling stop temperature of 260 ° C. or lower (° C.) / cooling from the cooling start temperature of 680 ° C. or higher to the cooling stop temperature of 260 ° C. or lower. time (seconds)".

If the cooling start temperature is less than 680°C, the area ratio of martensite cannot be increased to 95% or more. Therefore, the cooling start temperature is set to 680° C. or higher. Preferably, this cooling start temperature is 800° C. or lower.

If the average cooling rate is less than 70°C/sec, a large amount of ferrite and bainite will be generated and sufficient strength cannot be obtained. Therefore, the average cooling rate is set to 70° C./second or more. The average cooling rate is preferably at least 700°C/sec.

In addition, when the cooling stop temperature exceeds 260°C, there is a problem that upper and lower bainite are generated, and retained austenite and fresh martensite increase. Therefore, the cooling stop temperature is set to 260° C. or lower.

The carbides distributed inside the martensite are the carbides that are generated during holding in a low temperature range after quenching. In order to ensure excellent delayed fracture resistance and tensile strength of 1470 MPa or more (TS≧1470 MPa), it is necessary to appropriately control the formation of the carbides.
For this purpose, it is necessary to control the holding temperature to 150 to 260° C. and the holding time to 20 to 1500 seconds.
If the holding temperature is lower than the lower limit of 150° C. or if the holding time is short, the distribution density of carbides inside the transformation phase becomes insufficient and the delayed fracture resistance deteriorates. On the other hand, if the holding temperature is higher than the upper limit of 260° C., coarsening of carbides in grains and block grain boundaries becomes significant, possibly deteriorating the delayed fracture resistance. On the other hand, when the holding time exceeds 1500 seconds, coarsening of carbides in grains and block grain boundaries becomes significant, and there is a possibility that delayed fracture resistance deteriorates. Therefore, in the present invention, the continuous annealing is carried out at a holding temperature of 150 to 260° C. for 20 to 1500 seconds.

The steel sheet thus obtained can be subjected to skin-pass rolling from the viewpoint of stabilizing press formability, such as adjusting the surface roughness and flattening the plate shape. In that case, the skin pass elongation rate is preferably 0.1% or more. Also, the skin pass elongation rate is preferably 0.6% or less. In this case, dull rolls are used as the skin pass rolls, and it is preferable to adjust the roughness Ra of the steel sheet to 0.8 μm or more from the viewpoint of flattening the shape. Further, it is preferable to adjust the roughness Ra of the steel sheet to 1.8 μm or less.

Also, the obtained steel sheet may be subjected to a plating treatment. That is, after continuous annealing, the surface of the steel sheet may be plated. A steel sheet having a plating layer on its surface can be obtained by plating.

As described above, according to the present invention, the delayed fracture resistance of high-strength cold-rolled steel sheets is greatly improved, and the application of high-strength steel sheets contributes to the improvement of part strength and weight reduction. The steel sheet of the present invention preferably has a thickness of 0.5 mm or more. Also, the plate thickness is preferably 2.0 mm or less.

Next, the member of the present invention and its manufacturing method will be described.

The member of the present invention is obtained by subjecting the steel plate of the present invention to at least one of forming and joining. Further, the method for manufacturing a member of the present invention includes a step of subjecting the steel plate of the present invention to at least one of forming and joining to form a member.

The steel sheet of the present invention has a tensile strength of 1470 MPa or more and excellent delayed fracture resistance. Therefore, members obtained using the steel sheet of the present invention also have high strength and are superior in delayed fracture resistance to conventional high-strength members. Moreover, if the member of the present invention is used, the weight can be reduced. Therefore, the member of the present invention can be suitably used for, for example, vehicle body frame parts.

General processing methods such as press processing can be used without restrictions for molding. In addition, as the joining process, general welding such as spot welding and arc welding, riveting, caulking, and the like can be used without limitation.

Examples of the present invention will be described below.
Steels having chemical compositions shown in Table 1 were melted and then cast into slabs.
This slab was subjected to the heat treatment and rolling shown in Table 2 to obtain a steel plate having a thickness of 1.4 mm.

Specifically, a slab having each component composition was heated at an average heating rate of 6° C./min up to the heating and holding temperature shown in Table 2 at the slab surface temperature, and held for the heating and holding time shown in Table 2. After that, the slab is retained for the residence time at 900 to 1000 ° C. shown in Table 2, and hot finish rolling is performed at a finish rolling temperature of 870 ° C. The average cooling rate in the range from the finish rolling temperature to 650 ° C. is 50 ° C. / sec.
After that, it was cooled and coiled at a coiling temperature of 550°C to obtain a hot-rolled steel sheet, and the hot-rolled steel sheet was cold-rolled at a reduction rate (cold rolling reduction rate) of 50% to obtain a cold-rolled steel sheet. .
After that, the cold-rolled steel sheet was heated from 400° C. to the annealing temperature shown in Table 2 at the average heating rate shown in Table 2, and soaked at the annealing temperature for the soaking time shown in Table 2.
After that, from the cooling start temperature shown in Table 2 to the cooling stop temperature shown in Table 2, it is cooled at the average cooling rate shown in Table 2 and reheated as necessary. Continuous annealing was performed for the indicated holding times.

Also, No. For No. 13, the obtained steel sheet was subjected to electroplating to obtain a steel sheet having a Zn plating layer formed thereon.

The metal structure of the obtained steel sheet was quantified by the method described above, and then a tensile test and a delayed fracture resistance evaluation test were performed.
Specifically, the tissue measurement method was as follows.
The area ratio of martensite, bainite, and ferrite was determined by polishing the L cross section (vertical cross section parallel to the rolling direction) of the steel sheet, corroding it with nital, and examining the 1/4 thickness position from the surface of the steel sheet with an SEM at a magnification of 2000 times. Visual field observation was performed, and the photographed tissue photograph was image-analyzed and measured. Here, martensite and bainite refer to gray or white structures in SEM. Here, bainite has the following characteristics. That is, it has an aspect ratio of 2.5 or more, exhibits a plate-like form, and is a slightly blacker structure than martensite. The width of the plate is 0.3-1.7 μm. The distribution density of carbides with a diameter of 10 to 200 nm inside the bainite is 0 to 3 pieces/μm ² . On the other hand, ferrite is a region exhibiting black contrast in SEM. Note that the martensite and bainite contain trace amounts of carbides, nitrides, sulfides, and oxides, but since it is difficult to exclude them, the area ratio of the region including these was used as the area ratio.

　The retained austenite (retained γ) was measured by chemically polishing a 200 μm surface layer of the steel plate with oxalic acid and using the X-ray diffraction intensity method for the plate surface. It was calculated from the integrated intensities of (200)α, (211)α, (220)α, (200)γ, (220)γ, and (311)γ diffraction surface peaks measured by Mo-Kα radiation.

The average grain size of prior austenite grains (prior γ grain size) is measured by polishing the L cross section (vertical cross section parallel to the rolling direction) of the steel sheet, and then applying a chemical solution that corrodes the prior γ grain boundaries (such as a saturated picric acid aqueous solution or It was corroded by adding ferric chloride to this), and observed at 1/4 thickness position from the steel plate surface with an optical microscope at a magnification of 500 times. Fifteen lines in each direction were drawn at intervals of 10 μm or more in actual length, and the number of intersections between the grain boundaries and the lines was counted. The prior-γ grain size was obtained by multiplying the value obtained by dividing the total line length by the number of intersections by 1.13.

The number density A of precipitates having an equivalent circle diameter of 500 nm or more is obtained by polishing the L cross section (vertical cross section parallel to the rolling direction) of the steel sheet, and then the area from the 1/5 position to the 4/5 position of the steel plate thickness, that is, from the steel plate surface. In the region from the 1/5th position to the 4/5th position across the center of the plate thickness, a 2 mm ² region was continuously photographed with an SEM, and from the photographed SEM photographs, such precipitates It was obtained by counting the number of Also, the magnification for photographing is 2000 times. Moreover, when performing the component analysis of each inclusion particle, each inclusion particle was magnified 10000 times, and the said precipitate was analyzed. Here, the precipitate having an equivalent circle diameter of 500 nm or more is a precipitate containing B such as Fe ₂₃ (C, B) ₆ , and an elemental analysis by energy dispersive X-ray spectroscopy (EDS) at an acceleration voltage of 3 kV The presence or absence of a B peak was examined, and when there was a B peak, it was evaluated that the above precipitate was present.

For the tensile test, a JIS No. 5 tensile test piece was cut out so that the direction perpendicular to the rolling direction was the longitudinal direction at the position of 1/4 of the coil width, and a tensile test (based on JIS Z2241) was performed to evaluate YP, TS, and El.

Evaluation of delayed fracture resistance was performed as follows.
A strip test piece of 100 mm in the direction perpendicular to the rolling direction and 30 mm in the rolling direction was taken from the 1/4 position of the coil width in the width direction of the obtained steel plate (coil). The end face of the long side with a length of 100 mm is cut out by shearing, and in the state of shearing (without machining to remove burrs), bending is performed so that the burrs are on the outer peripheral side of the bend. , the test piece was fixed with bolts while maintaining the shape of the test piece at the time of bending. The shearing clearance was 13% and the rake angle was 1°. Bending was performed with a tip bending radius of 10 mm and an angle inside the bending apex of 90 degrees (V bending). The punch has a U-shaped punch whose tip radius is the same as the tip bending radius R (the tip R is semicircular and the thickness of the punch barrel is 2R), and the die has a corner radius of 30 mm. was used. The depth to which the punch pushes the steel plate was adjusted, and the steel plate was formed so that the bending angle of the tip (the angle inside the bending apex) was 90 degrees (V shape). Sandwich the test piece with a hydraulic jack so that the distance between the flange ends of the straight piece when bending is the same as when bending (to cancel out the opening of the straight piece due to springback). and bolted in that state. The bolt was passed through an elliptical hole (minor axis: 10 mm, major axis: 15 mm) previously provided 10 mm inward from the short side edge of the strip test piece and fixed. The obtained test piece after bolting was mixed with 0.1% by mass of ammonium thiocyanate aqueous solution and McIlvaine buffer solution at a ratio of 1:1 and immersed in a solution adjusted to pH 8.0 for delayed fracture resistance. A characterization test was performed. At this time, the temperature of the solution was set at 20° C., and the amount of the solution per 1 cm ³ of surface area of the test piece was set at 20 ml. After 24 hours had elapsed, the presence or absence of cracks at a visually recognizable level (1 mm or more in length) was confirmed, and those in which no cracks were observed were judged to have excellent delayed fracture resistance.

Table 3 shows the structure and properties of the obtained steel sheets.

The steel sheets within the scope of the present invention had high strength and excellent delayed fracture resistance.
On the other hand, No. In No. 14 (Steel N), the C content was less than the lower limit of the specified value of the present invention, and the TS was insufficient.
No. In No. 15 (steel O), the C content exceeded the upper limit of the specified value of the present invention, and sufficient delayed fracture resistance was not obtained.
No. In No. 16 (Steel P), the P content exceeded the upper limit of the specified value of the present invention, and sufficient delayed fracture resistance was not obtained.
No. In No. 17 (steel Q), the S content exceeded the upper limit of the specified value of the present invention, and sufficient delayed fracture resistance was not obtained.
No. 18 (steel R) is sol. The Al content exceeded the upper limit of the value specified in the present invention, and sufficient delayed fracture resistance was not obtained.
No. In No. 19 (steel S), the N content exceeded the upper limit of the specified value of the present invention, and sufficient delayed fracture resistance was not obtained.
No. In No. 20 (Steel T), the Nb and Ti contents were less than the lower limits of the specified values of the present invention, the prior γ grain size was large, and sufficient delayed fracture resistance was not obtained.
No. In No. 21 (steel U), the Nb and Ti contents exceeded the upper limits of the values specified in the present invention, and sufficient delayed fracture resistance was not obtained.
No. In No. 22 (steel V), the B content exceeded the upper limit of the value specified in the present invention, and there were many coarse precipitates, and sufficient delayed fracture resistance was not obtained.
No. In No. 23 (steel W), the B content was less than the lower limit of the value specified in the present invention, and sufficient delayed fracture resistance was not obtained.
No. In No. 24 (Steel A), the heating temperature (slab surface temperature (SRT)) is less than the lower limit of the specified value of the present invention, the prior γ grain size is large, the number density A of precipitates is excessive, and sufficient delayed fracture resistance is obtained. I couldn't.
No. In No. 25 (steel A), the slab heating and holding time was less than the lower limit of the value specified in the present invention, the prior γ grain size was large, the number density A of precipitates was excessive, and sufficient delayed fracture resistance was not obtained.
No. In No. 26 (steel A), the residence time at 900 to 1000° C. exceeded the upper limit of the specified value of the present invention, the number density A of precipitates was excessive, and sufficient delayed fracture resistance was not obtained.
No. In No. 27 (Steel A), the average heating rate during annealing was less than the lower limit of the specified value of the present invention, the prior γ grain size was large, the number density A of precipitates was excessive, and sufficient delayed fracture resistance was not obtained.
No. In No. 28 (Steel A), the soaking time during annealing exceeded the upper limit of the specified value of the present invention, the prior γ grain size was large, and sufficient delayed fracture resistance was not obtained.
No. In No. 29 (Steel A), the cooling start temperature during annealing was less than the lower limit of the specified value of the present invention, martensite was not sufficiently formed, and sufficient delayed fracture resistance was not obtained.
No. In No. 30 (Steel A), the average cooling rate during annealing was less than the lower limit of the value specified in the present invention, and the formation of martensite was insufficient, and sufficient delayed fracture resistance was not obtained.

In addition, the steel plate of the present invention example has high strength and excellent delayed fracture resistance in the members obtained by molding and the members obtained by joining processing using the steel plate of the example of the present invention. Therefore, it was found that the steel sheets have high strength and excellent delayed fracture resistance, like the steel sheets of the invention examples.

Claims

in % by mass,
C: 0.15% or more and 0.45% or less,
Si: 1.5% or less,
Mn: 1.7% or less,
P: 0.03% or less,
S: less than 0.0020%,
sol. Al: 0.20% or less,
N: 0.005% or less,
B: 0.0015% or more and 0.0100% or less,
Containing 0.005% or more and 0.080% or less in total of one or more of Nb and Ti,
Having a component composition in which the balance is Fe and unavoidable impurities,
Having a structure in which the area ratio of martensite to the entire structure is 95% or more and 100% or less,
The average grain size of the prior austenite grains is less than 11.0 μm,
A steel sheet in which the number density A of precipitates having an equivalent circle diameter of 500 nm or more satisfies the following formula (1).
A (number/mm 2 )≦8.5×10 5 ×[B] Expression (1)
Here, [B] represents the content of B (% by mass).
The steel sheet according to claim 1, further comprising, as the chemical composition, one or two selected from Cu: 1.0% or less and Ni: 1.0% or less in mass%.
As the component composition, in mass%, Cr: 1.0% or less, Mo: less than 0.3%, V: 0.5% or less, Zr: 0.2% or less and W: 0.2% or less The steel sheet according to claim 1 or 2, containing one or more selected from among them.
As the component composition, in mass%, Ca: 0.0030% or less, Ce: 0.0030% or less, La: 0.0030% or less, REM (excluding Ce and La): 0.0030% or less and Mg The steel sheet according to any one of claims 1 to 3, containing one or more selected from: 0.0030% or less.
The composition according to any one of claims 1 to 4, further comprising one or two selected from Sb: 0.1% or less and Sn: 0.1% or less in mass%. steel plate.
The steel sheet according to any one of claims 1 to 5, which has a plating layer on the surface of the steel sheet.
A member using the steel plate according to any one of claims 1 to 6.
A steel slab having the chemical composition according to any one of claims 1 to 5 is heated at a slab surface temperature of 1000 ° C. to a heating and holding temperature of 1250 ° C. or higher at an average heating rate of 10 ° C./min or less, and the heating After holding at the holding temperature for 30 minutes or more,
Hot finish rolling is performed under the conditions of a residence time of 20 seconds or more and 150 seconds or less at 900 to 1000 ° C. and a finish rolling temperature of 850 ° C. or more,
Cooling is performed at an average cooling rate of 40 ° C./sec or more in the range from the finish rolling temperature to 650 ° C.,
After that, it is coiled at a coiling temperature of 650 ° C. or less to form a hot rolled steel sheet,
A cold-rolled steel sheet is obtained by cold-rolling the hot-rolled steel sheet at a rolling reduction of 40% or more,
The annealing temperature is 830 to 950° C., and the cold-rolled steel sheet is heated from 400° C. to the annealing temperature at an average heating rate of 1.0° C./second or more,
Hold at the annealing temperature for 600 seconds or less,
Cooling from a cooling start temperature of 680 ° C. or higher to a cooling stop temperature of 260 ° C. or lower at an average cooling rate of 70 ° C./sec or higher,
A method for manufacturing a steel sheet, followed by continuous annealing at a holding temperature of 150 to 260° C. for 20 to 1500 seconds.
The steel sheet manufacturing method according to claim 8, wherein the surface of the steel sheet is plated after the continuous annealing.
A method of manufacturing a member, comprising the step of subjecting the steel plate according to any one of claims 1 to 6 to at least one of forming and joining to form a member.