WO2024122037A1 - High-strength steel sheet, member formed using high-strength steel sheet, automobile framework structure component or automobile reinforcing component composed of member, and production methods for high-strength steel sheet and member - Google Patents

High-strength steel sheet, member formed using high-strength steel sheet, automobile framework structure component or automobile reinforcing component composed of member, and production methods for high-strength steel sheet and member Download PDF

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WO2024122037A1
WO2024122037A1 PCT/JP2022/045373 JP2022045373W WO2024122037A1 WO 2024122037 A1 WO2024122037 A1 WO 2024122037A1 JP 2022045373 W JP2022045373 W JP 2022045373W WO 2024122037 A1 WO2024122037 A1 WO 2024122037A1
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strength steel
sheet
steel sheet
steel plate
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PCT/JP2022/045373
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French (fr)
Japanese (ja)
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悠佑 和田
秀和 南
勇樹 田路
由康 川崎
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Ｊｆｅスチール株式会社
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Priority to JP2023519413A priority Critical patent/JP7367893B1/en
Priority to PCT/JP2022/045373 priority patent/WO2024122037A1/en
Publication of WO2024122037A1 publication Critical patent/WO2024122037A1/en

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  • the present invention relates to high-strength steel plates, components made of high-strength steel plates, automobile structural components or automobile reinforcing parts made of components, and methods for manufacturing high-strength steel plates and components.
  • Patent Document 1 proposes a high-strength steel plate having a specified component composition and microstructure, and in which the grain size of the iron carbide contained in the low-temperature transformation phase is 500 nm or less.
  • parts such as crash boxes have punched end faces and bent parts, so the steel plates used in these parts are required to have good ductility, stretch flangeability, and bendability of the shear end faces.
  • parts using high-strength steel plates with a tensile strength of 780 MPa or more are used in low-temperature environments, there is a risk that their toughness will deteriorate and cracks will occur during a collision. Therefore, automotive steel plates are required to have excellent low-temperature toughness to prevent cracks during a collision when used in low-temperature environments.
  • high strength steel plate refers to a steel plate having a tensile strength (TS) of 780 MPa or more as determined by a tensile test described later.
  • Excellent component strength means that the yield ratio (YR) determined by the tensile test described below is 55% or more.
  • excellent ductility means that the total elongation (El) determined by the tensile test described below is 10% or more.
  • excellent stretch flangeability means that the hole expansion ratio ( ⁇ ) determined by the hole expansion test described below is 20% or more.
  • Excellent bendability at the sheared end surface means that the ratio (Rs/Rg) of the limit bending radius (Rs/t) determined in a bending test of a sample having a sheared end surface (described later) to the limit bending radius (Rg/t) determined in a bending test of a sample having a ground end surface is 1.50 or less.
  • excellent low-temperature toughness means that the low-temperature toughness parameter (P) is 3000 or more in the Charpy impact test described below.
  • the gist of the present invention is as follows. (1) In mass%, C: 0.030% or more and 0.500% or less, Si: 0.01% or more and 2.50% or less, Mn: 0.10% or more and 5.00% or less, P: 0.100% or less, S: 0.0200% or less, Al: 1.000% or less, A composition comprising N: 0.0100% or less and O: 0.0100% or less, with the balance being Fe and unavoidable impurities; At the 1/4 plate thickness position, The area ratio of martensite is 10% or more and 80% or less, The area ratio of bainite is 2% or more and 70% or less, The area ratio of ferrite is 80% or less, a steel structure in which the area ratio of retained austenite is 15% or less and the ratio of the number of martensite blocks in which metastable carbides are present to the number of mar
  • the composition further includes, in mass%, Ti: 0.200% or less, Nb: 0.200% or less, V: 0.200% or less, Ta: 0.10% or less, W: 0.10% or less, B: 0.0100% or less, Cr: 1.00% or less, Mo: 1.00% or less, Ni: 1.00% or less, Co: 0.010% or less, Cu: 1.00% 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: 0.100% or less, Te: 0.100% or less,
  • the high-strength steel plate has a Vickers hardness of 85% or less of the Vickers hardness at a 1/4 position in the plate thickness direction of the high-strength steel plate, and has a surface soft layer which is a region within 200 ⁇ m from the surface of the high-strength steel plate in the plate thickness direction;
  • the ratio of the number of measurements in which the nano hardness of the sheet surface at a position of 1/4 of the sheet thickness direction depth of the soft surface layer from the surface of the high strength steel sheet is 7.0 GPa or more to the total number of measurements is 0.10 or less,
  • the high-strength steel plate according to any one of (1) to (5) above having a metal plating layer containing at least one of zinc and aluminum in a total amount of 50% by mass or more on one or both outermost layers of the high-strength steel plate.
  • a first cooling step in which the first cooling rate in a temperature range from T2 to 750°C is 2.0°C/s or more;
  • a second heating step is performed under conditions of a temperature X (° C.) and a holding time Y (s) that satisfy the following formula 2.
  • Formula 2 7000 ⁇ (273+X)(20+log(Y/3600)) ⁇ 13000 (10)
  • a metal plating containing more than 50 mass% of one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi
  • any of the manufacturing methods according to (9) to (12) above comprising a step of applying metal plating containing 50 mass% or more in total of at least one of zinc and aluminum to the steel sheet subjected to the first heating and the second heating steps.
  • a method for manufacturing a component comprising the step of subjecting any one of the high-strength steel plates according to (1) to (6) above to at least one of forming and joining to form a component.
  • the present invention it is possible to provide a high-strength steel sheet excellent in part strength, ductility, stretch flangeability, bendability of a sheared end surface, and low-temperature toughness. Also, it is possible to provide a member made using the high-strength steel sheet. Furthermore, according to the present invention, there can be provided a method for manufacturing the above-mentioned high-strength steel plate and a method for manufacturing a member using the high-strength steel plate. In addition, according to the present invention, it is possible to provide an automobile frame structural part or an automobile reinforcing part made of the above-mentioned member.
  • FIGS. 2A and 2B are schematic diagrams showing the preparation of samples for V-bending and orthogonal VDA bending tests in the examples, in which Fig. 2A shows V-bending (primary bending) and Fig. 2B shows orthogonal VDA bending (secondary bending).
  • 3(a) is a front view of a test member
  • FIG. 3(b) is a front view of a test member
  • FIG. 3(c) is a schematic diagram showing an axial crush test.
  • the high-strength steel plate of the present invention (hereinafter, for convenience, also referred to as "steel plate”) has a chemical composition and a steel structure described below.
  • composition of the high strength steel plate of the present invention (hereinafter, for convenience, also referred to as the “composition of the present invention") will be described.
  • “%” in the composition of the present invention means “mass%” unless otherwise specified.
  • C is one of the important basic components of steel, and in particular in the present invention, it affects the area ratio of martensite. If the C content is too low, the area ratio of martensite decreases, making it difficult to achieve a TS of 780 MPa or more. For this reason, the C content is 0.030% or more, preferably 0.040% or more, and more preferably 0.050% or more. On the other hand, if the C content is too high, the amount of retained austenite increases excessively, and the hardness of martensite generated from the retained austenite during punching increases significantly. As a result, crack propagation during hole expansion is promoted, the hole expansion ratio decreases, and the stretch flangeability decreases.
  • the residual austenite undergoes stress-induced transformation, which reduces the YR and reduces the strength of the part.
  • the C content is 0.500% or less, preferably 0.400% or less, and more preferably 0.300% or less.
  • Silicon is a component that increases the strength of a steel sheet by suppressing the precipitation of cementite in martensite and by solid solution strengthening.
  • the silicon content is 0.01% or more, preferably 0.05% or more, and more preferably 0.10% or more.
  • the Si content is too high, the carbide precipitation during bainite transformation is significantly suppressed, the residual austenite increases excessively, and the hardness of martensite generated from the residual austenite during punching increases significantly. As a result, crack growth during hole expansion is promoted, the hole expansion ratio decreases, and the stretch flangeability decreases.
  • the residual austenite undergoes stress-induced transformation, which reduces the YR and reduces the part strength.
  • the Si content is 2.50% or less, preferably 2.00% or less, and more preferably 1.50% or less.
  • Mn is one of the important basic components of steel, and particularly in the present invention, it affects the area ratio of martensite. If the Mn content is too low, the area ratio of martensite decreases, making it difficult to achieve a TS of 780 MPa or more. Therefore, the Mn content is 0.10% or more, preferably 0.90% or more, and more preferably 1.80% or more. On the other hand, if the Mn content is too high, the austenite is stabilized, the residual austenite is excessively increased, and the hardness of the martensite generated from the residual austenite during punching is greatly increased.
  • the Mn content is 5.00% or less, preferably 4.20% or less, and more preferably 3.60% or less.
  • P is a component that segregates at prior austenite grain boundaries to embrittle the grain boundaries, thereby reducing the ultimate deformability of the steel sheet, thereby reducing ⁇ and reducing bendability.
  • the P content is 0.100% or less, and preferably 0.070% or less.
  • the lower limit of the P content is not particularly limited, but since P is a solid solution strengthening element and can increase the strength of the steel sheet, it is preferable to set the lower limit to 0.001% or more.
  • ⁇ S 0.0200% or less ⁇ S exists as a sulfide and is a component that reduces the ultimate deformability of the steel sheet, thereby reducing ⁇ and reducing bendability. Therefore, the S content is 0.0200% or less, preferably 0.0050% or less.
  • the lower limit of the S content is not particularly limited, but due to constraints on production technology, it is preferably set to 0.0001% or more.
  • Al is an effective component for sufficient deoxidation and reducing inclusions in steel, but if the Al content is too high, a large amount of ferrite is generated, the hole expansion ratio decreases, and the stretch flangeability may decrease. Therefore, the Al content is 1.000% or less, preferably 0.500% or less, and more preferably 0.100% or less. On the other hand, in order to perform stable deoxidation, the Al content is preferably 0.010% or more, more preferably 0.015% or more, and even more preferably 0.020% or more.
  • N exists as a nitride and is a component that reduces the ultimate deformability of the steel sheet, thereby reducing ⁇ and reducing bendability. Therefore, the N content is 0.0100% or less, preferably 0.0050% or less. Although there is no particular lower limit for the N content, due to constraints on production technology, the N content is preferably 0.0001% or more.
  • O exists as an oxide and is a component that reduces the ultimate deformability of the steel sheet, thereby reducing ⁇ and reducing bendability. Therefore, the O content is 0.0100% or less, and preferably 0.0050% or less. Although there is no particular lower limit for the O content, due to constraints on production technology, the O content is preferably 0.0001% or more.
  • the high strength steel plate of the present invention further comprises, in addition to the above-mentioned composition, in mass%: Ti: 0.200% or less, Nb: 0.200% or less, V: 0.200% or less, Ta: 0.10% or less, W: 0.10% or less, B: 0.0100% or less, Cr: 1.00% or less, Mo: 1.00% or less, Ni: 1.00% or less, Co: 0.010% or less, Cu: 1.00% 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: 0.020% or less, Te: 0.020% or less, At least one element selected from the group consisting of Hf: 0.10% or less and Bi: 0.200% or less may be contained. These elements may be contained alone or in combination of two or more kinds.
  • the Ti, Nb or V content is preferably 0.200% or less, and more preferably 0.100% or less.
  • the Ti, Nb or V content is 0.001% or more.
  • the Ta or W content is preferably 0.10% or less, and more preferably 0.08% or less.
  • the Ta or W content is preferably 0.01% or more.
  • the B content is preferably 0.0100% or less, and more preferably 0.0003% or more.
  • the B content is preferably 0.0003% or more.
  • the Cr, Mo or Ni content is 1.00% or less, and more preferably 0.80% or less.
  • the Cr, Mo or Ni content is 0.01% or more.
  • the Co content is preferably 0.010% or less, and more preferably 0.008% or less.
  • the Co content is 0.001% or more.
  • the Cu content is preferably 1.00% or less, and more preferably 0.80% or less.
  • the Cu content is 0.01% or more.
  • the Sn content is preferably 0.200% or less, and more preferably 0.100% or less. There is no particular lower limit for the Sn content, but since Sn is an element that improves hardenability, the Sn content is preferably 0.001% or more.
  • the Sb content is preferably 0.200% or less, and more preferably 0.100% or less.
  • the Sb content is preferably 0.001% or more.
  • the Ca, Mg or REM content is 0.0100% or less, and more preferably 0.0050% or less.
  • the Ca, Mg or REM content is 0.0001% or more.
  • the Zr or Te content is 0.100% or less, and more preferably 0.080% or less.
  • the Zr or Te content is 0.001% or more.
  • the Hf content is preferably 0.10% or less, and more preferably 0.08% or less.
  • the Hf content is preferably 0.01% or more.
  • the Bi content is preferably 0.200% or less, and more preferably 0.100% or less.
  • the Bi content is 0.001% or more.
  • the high-strength steel plate according to one embodiment of the present invention has a composition containing the above essential components and optional components, with the balance being Fe and unavoidable impurities.
  • the unavoidable impurities include Zn, Pb, As, Ge, Sr, and Cs. These unavoidable impurities are permitted to be contained in an amount of 0.100% or less in total.
  • ⁇ Area ratio of martensite 10% to 80%>
  • the area ratio of martensite is 10% or more, preferably 15% or more, and more preferably 20% or more.
  • the area ratio of martensite is 80% or less, preferably 75% or less, and more preferably 70% or less.
  • martensite includes lower bainite, martensite that has undergone self-tempering during cooling performed in the annealing step described later, and martensite that has been tempered in the second heating step described later. As described later, the martensite was observed at a position corresponding to 1/4 of the thickness of the steel plate.
  • ⁇ Area ratio of bainite 2% to 70%>
  • the hardness difference between the structures is reduced and ⁇ is increased.
  • the crack propagation at the interface is suppressed, and thus low-temperature toughness is improved.
  • the area fraction of bainite is 2% or more, preferably 3% or more, and more preferably 4% or more.
  • the area ratio of bainite is 70% or less, preferably 60% or less, and more preferably 50% or less.
  • bainite is a mixed structure of angular bainitic ferrite, iron-based carbides, and retained austenite that is formed in a temperature range of Ms or higher and 700° C. or lower.
  • the observation position of bainite is a quarter position of the sheet thickness of the steel sheet, as described later.
  • the area ratio of ferrite 80% or less
  • the desired strength can be easily obtained.
  • the effect of the present invention can be obtained even if the area ratio of ferrite is 0%.
  • the area ratio of ferrite is 80% or less, preferably 75% or less, and more preferably 70% or less.
  • the area ratio of ferrite is preferably 10% or more, and more preferably 15% or more.
  • ferrite is soft BCC iron formed at relatively high temperatures, and includes allotriomorph ferrite and idiomorph ferrite.
  • the observation position of ferrite was a quarter position of the sheet thickness of the steel sheet, as described later.
  • the method for measuring the area ratios of martensite, bainite, and ferrite is as follows. First, a sample is cut out from a steel sheet so that a thickness cross section (L cross section at 1/4 of the sheet thickness) parallel to the rolling direction becomes an observation surface. The observation surface of the sample is mirror-polished with diamond paste, then finish-polished with colloidal silica, and further etched with 1 volume % nital to reveal the structure. Next, the observation surface of the sample is observed at a magnification of 3000 times using a scanning electron microscope (SEM) under the condition of an acceleration voltage of 10 kV, and SEM images of three visual fields (one visual field is 40 ⁇ m ⁇ 30 ⁇ m) are obtained.
  • SEM scanning electron microscope
  • the area ratio of each structure is calculated using Adobe Photoshop (manufactured by Adobe Systems). Specifically, the value obtained by dividing the area of each structure by the measured area is regarded as the area ratio of each structure. The area ratio of each structure is calculated for three fields of view, and the average value thereof is regarded as the area ratio of each structure.
  • Bainite is a mixed structure region consisting of gray angular bainitic ferrite, white contrasting iron carbides, and needle-shaped retained austenite. Martensite has a hierarchical structure with minute internal irregularities. These can be distinguished from one another.
  • the area ratio of the retained austenite is 15% or less, and preferably 10% or less. There is no particular lower limit, and this effect can be obtained even if the area ratio of the retained austenite is 0%.
  • the method for measuring the area ratio of retained austenite is as follows. First, the steel plate is ground so that the 1/4 position of the plate thickness (the position corresponding to 1/4 of the plate thickness in the depth direction from the surface of the steel plate) becomes the measurement surface, and then the plate is further polished by 0.1 mm by chemical polishing to obtain a sample. For the measurement surface of the sample, an X-ray diffractometer is used to measure 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, using a Co K ⁇ radiation source.
  • the intensity ratio of the integrated reflection intensity of each surface of the fcc iron to the integrated reflection intensity of each surface of the bcc iron is calculated.
  • the average value of the nine intensity ratios is taken as the volume fraction of the retained austenite.
  • This volume fraction of the retained austenite is considered to be three-dimensionally uniform, and is taken as the area fraction of the retained austenite at the 1/4 position of the plate thickness of the steel plate.
  • the steel structure of the present invention may have a structure (remaining structure) other than the above-mentioned martensite, bainite, ferrite, and retained austenite.
  • the remaining structure is a structure other than martensite, bainite, ferrite, and retained austenite, and may be any structure known as a steel sheet structure, such as pearlite and alloy carbonitrides precipitated in ferrite. It should be noted that iron-based carbides present in bainite, metastable carbides precipitated in martensite, and iron-based carbides such as cementite precipitated in martensite are not included in the remaining structure.
  • the area ratio of the remaining structure is preferably 3% or less so as not to impair the effects of the present invention.
  • the metastable carbides precipitated in the martensite block improve low-temperature toughness while maintaining excellent part strength, ductility, shear edge bendability, and stretch flangeability. This is believed to be because the metastable carbides precipitated in the martensite block suppress the initiation and propagation of cracks at low temperatures.
  • ratio p the ratio of the number of martensite blocks in which metastable carbides exist to the number of martensite blocks.
  • ratio p is 2% or more, preferably 5% or more, more preferably 10% or more, even more preferably 20% or more, and particularly preferably 30% or more.
  • the upper limit of ratio p is not particularly limited and may be 100%.
  • metastable carbides are metastable carbides that precipitate during the tempering process of martensite.
  • Metastable carbides are, for example, Fe carbides (iron-based carbides) other than cementite, and include at least one type of carbide selected from the group consisting of epsilon ( ⁇ ) carbides, eta ( ⁇ ) carbides, and chi ( ⁇ ) carbides.
  • the method for measuring the ratio (ratio p) of the number of martensite blocks containing metastable carbides to the number of martensite blocks is as follows. First, a steel sheet is ground so that a 1/4 position of the sheet thickness (a position corresponding to 1/4 of the sheet thickness in the depth direction from the surface of the steel sheet) becomes an observation surface, and then electrolytic polishing is performed to prepare a sample. The observation surface of the prepared sample is observed using a transmission electron microscope (TEM) at an acceleration voltage of 200 kV. When an electron beam is incident on the martensite block from the [100] direction, an electron diffraction pattern of the parent martensite is obtained.
  • TEM transmission electron microscope
  • Adjacent martensite blocks have different crystal orientations across the block boundaries, and therefore can be distinguished from each other by the different contrast in the bright-field image. Martensite can be distinguished from ferrite and bainite by the high density of dislocations observed in martensite and the relatively low dislocation density in ferrite and bainite.
  • FIG. 1 is an example of an electron diffraction pattern of martensite in which carbides are present.
  • the electron diffraction pattern of the carbides is obtained in addition to the electron diffraction pattern of the parent martensite ( ⁇ ), as shown in FIG.
  • black circles indicate electron diffraction spots of the martensite parent phase when the electron beam is incident from the [100] direction
  • white circles indicate electron diffraction spots of carbides.
  • the martensite block is defined as a martensite block containing metastable carbides.
  • the ratio of the number of martensite blocks containing metastable carbides to the number of martensite blocks can be rephrased as "the ratio of the number of martensite blocks with a ratio dc/dm of 1.020 or more and 1.150 or less to the number of martensite blocks.”
  • Metastable carbides may be present inside the martensite blocks or at boundary portions such as block boundaries, but are preferably present inside the martensite blocks.
  • Average number density of metastable carbides in martensite blocks containing metastable carbides 1 x 10 6 /mm 2 or more
  • a high number density of metastable carbides in the martensite blocks is preferred for better low temperature toughness reasons, as it is believed that a high number density of metastable carbides provides greater resistance to crack propagation in the martensite at low temperatures.
  • the average number density of metastable carbides in a martensite block in which metastable carbides are present (hereinafter also referred to as "number density n”) is preferably 1 x 10 pcs/mm 2 or more, more preferably 10 x 10 pcs/mm 2 or more, and even more preferably 100 x 10 pcs/mm 2 or more.
  • the upper limit of the number density n is not particularly limited, and the number density n can be, for example, 10,000,000 ⁇ 10 6 pieces/mm 2 or less, preferably 1,000,000 ⁇ 10 6 pieces/mm 2 or less, more preferably 100,000 ⁇ 10 6 pieces/mm 2 or less, and even more preferably 10,000 ⁇ 10 6 pieces/mm 2 or less.
  • the method for measuring the average number density (number density n) of metastable carbides in a martensite block in which metastable carbides are present is as follows.
  • a selected area electron diffraction pattern is obtained for a single martensite block in which metastable carbides are present, and a dark-field image is obtained using the electron diffraction spots obtained from the metastable carbides.
  • the metastable carbides show white contrast.
  • An area of 300 nm ⁇ 300 nm is photographed within a single martensite block, and the number of metastable carbides is counted. Note that adjacent martensite blocks may exist across a block boundary within the 300 nm ⁇ 300 nm area.
  • the area of a martensite block with metastable carbides is defined as the area of a single martensite block for which a selected-area electron diffraction pattern was obtained. Adjacent martensite blocks are distinguished from each other by their contrast in bright-field images due to different crystal orientations across the block boundaries.
  • the average of these values is regarded as the average number density (number density n) of the metastable carbides in the martensite block in which the metastable carbides are present.
  • Average circle equivalent diameter of metastable carbides 20 nm or less
  • the average circle-equivalent diameter of the metastable carbides in the martensite blocks is preferably 20 nm or less, and more preferably 5 nm or less.
  • the method for measuring the average value of the circle equivalent diameter of metastable carbides in a martensite block is as follows.
  • a selected area electron diffraction pattern is obtained for a single martensite block in which metastable carbides are present, and a dark-field image is obtained using the electron diffraction spots obtained from the metastable carbides.
  • the metastable carbides show white contrast.
  • a dark field image of a 300 nm x 300 nm area within a single martensite block is taken and image processing is performed to obtain a binary image in which metastable carbides can be distinguished.
  • the binary image is subjected to particle analysis to determine the circle equivalent diameter for each metastable carbide particle. If metastable carbides overlap in the dark field image, the binary image is segmented using the Watershed method. The circle equivalent diameter is determined for each of all metastable carbides present in the 300 nm ⁇ 300 nm region (three visual fields). The average of the circle equivalent diameters for the three visual fields is determined and this is set as the average circle equivalent diameter of the metastable carbides in the martensite block.
  • the standard deviation ⁇ n of the nanohardness is 0.60 ⁇ [H n ] ave or less, preferably 0.50 ⁇ [H n ] ave or less, where [H n ] ave is the average value of the nanohardness.
  • the lower limit of the standard deviation ⁇ n of the nanohardness is not particularly limited and may be 0.
  • the average nano-hardness [Hn] ave is preferably 3.0 GPa or more and 9.0 GPa or less, and more preferably 3.5 GPa or more and 8.5 GPa or less.
  • a method for measuring the standard deviation ⁇ n of nanohardness will be described.
  • a nanoindentation device equipped with a Berkovich indenter is used. After cutting out a sample so that the plate thickness cross section (L cross section) parallel to the rolling direction of the steel plate is the observation surface, the observation surface is mirror-polished using diamond paste, and then finish-polished using colloidal silica.
  • the nanohardness of 225 or more points is measured for the sample under load control conditions of a loading rate and unloading rate of 50 ⁇ N/s, a maximum load of 500 ⁇ N, and a data collection pitch of 5 msec.
  • the measurement position is set to 1/4 the plate thickness from the surface of the high-strength steel plate, and the distance between the indentations is set to 2 ⁇ m or more.
  • a histogram is created from the nanohardness measurement results obtained at 225 or more points, and the standard deviation is calculated and the result is defined as the standard deviation of the nanohardness ⁇ n .
  • the average value of the nanohardness measurement results obtained at 225 or more points is defined as [H n ] ave .
  • the high-strength steel sheet has a soft surface layer formed on the surface of the base steel sheet.
  • the soft surface layer contributes to suppressing the propagation of bending cracks during press forming and vehicle body collision, thereby improving bending fracture resistance.
  • the base steel sheet refers to a high-strength steel sheet that is the base (undercoat) for the various platings in the case of a steel sheet that has been subjected to a plating treatment, such as a hot-dip galvanized steel sheet, a galvannealed steel sheet, an electrolytic galvanized steel sheet, or a steel sheet that has been plated with other metals, and refers to a high-strength steel sheet in the case of a steel sheet that has not been plated.
  • the surface layer refers to a region corresponding to a thickness of 200 ⁇ m from the surface of the base steel sheet to a depth of 200 ⁇ m in the sheet thickness direction.
  • the soft layer refers to a region having a Vickers hardness of 85% or less of the Vickers hardness of a cross section (a plane parallel to the steel sheet surface) at 1/4 of the sheet thickness of the base steel sheet.
  • the soft layer includes a decarburized layer in the surface layer of the base steel sheet.
  • the surface soft layer refers to a soft layer included in the surface layer, and may be a soft layer in its entirety or a part of it.
  • the surface soft layer may be a region corresponding to a thickness of 200 ⁇ m or less from the surface of the base steel sheet in the sheet thickness direction.
  • a region having a Vickers hardness of 85% or less on a cross section (plane parallel to the steel plate surface) at 1/4 of the plate thickness of the base steel plate is formed at a predetermined depth from the surface of the base steel plate in the plate thickness direction, when the predetermined depth is within 200 ⁇ m in the plate thickness direction, the region corresponding to the thickness from the surface to the predetermined depth in the plate thickness direction is the surface soft layer, and when the predetermined depth is more than 200 ⁇ m in the plate thickness direction, the region corresponding to a thickness of 200 ⁇ m from the surface of the base steel plate to a depth of 200 ⁇ m in the plate thickness direction is the surface soft layer.
  • the lower limit of the thickness of the soft surface layer is not particularly limited, but is preferably 8 ⁇ m or more, and more preferably more than 17 ⁇ m.
  • the Vickers hardness is measured based on JIS Z 2244-1 (2020) at a load of 10 gf.
  • the proportion of nano hardness of 7.0 GPa or more should be 0.10 or less.
  • the proportion of nano hardness of 7.0 GPa or more is 0.10 or less, it means that the proportion of hard structures (martensite, etc.), inclusions, etc. is small, and it is possible to further suppress the generation and connection of voids and crack growth in hard structures (martensite, etc.) and inclusions during press forming and collision, and to obtain excellent R/t and SF max .
  • the standard deviation ⁇ of the nanohardness of the sheet surface at a position 1/4 of the sheet thickness direction depth of the soft surface layer from the surface of the base steel sheet is 1.8 GPa or less
  • the standard deviation ⁇ of the nanohardness of the sheet surface at a position 1/2 of the sheet thickness direction depth of the soft surface layer from the surface of the base steel sheet is 2.2 GPa or less
  • a more preferred range for the standard deviation ⁇ of the nanohardness of the sheet surface at a position 1/4 of the way from the base steel sheet surface to the soft surface layer in the sheet thickness direction is 1.7 GPa or less.
  • a more preferred range for the standard deviation ⁇ of the nanohardness of the sheet surface at a position 1/2 of the way from the base steel sheet surface to the soft surface layer in the sheet thickness direction is 2.1 GPa or less.
  • the nano-hardness of the plate surface at the 1/4 and 1/2 positions in the plate thickness direction depth is a hardness measured by the following method. First, if a plating layer is formed, after peeling off the plating layer, mechanical polishing is performed from the surface of the base steel sheet to a position 1/4 of the depth in the sheet thickness direction of the soft surface layer, buff polishing with diamond and alumina is performed, and further colloidal silica polishing is performed.
  • the nano hardness is measured using a Berkovich-shaped diamond indenter under the following conditions: load: 500 ⁇ N, measurement area: 50 ⁇ m ⁇ 50 ⁇ m, and impact spacing: 2 ⁇ m.
  • the soft surface layer is mechanically polished to a position halfway down the thickness direction, buffed with diamond and alumina, and then polished with colloidal silica.Then, the nano-hardness is measured with a Berkovich diamond indenter under the following conditions: load: 500 ⁇ N, measurement area: 50 ⁇ m ⁇ 50 ⁇ m, and impact spacing: 2 ⁇ m.
  • the thickness of the surface soft layer can be measured by the following method. After smoothing the thickness cross section (L cross section) of the base steel sheet parallel to the rolling direction by wet polishing, measurements were taken at 1 ⁇ m intervals using a Vickers hardness tester with a load of 10 gf from a position 1 ⁇ m from the surface of the base steel sheet in the thickness direction to a position 100 ⁇ m in the thickness direction. After that, measurements were taken at 20 ⁇ m intervals up to the center of the thickness. The area where the hardness has decreased to 85% or less compared to the hardness at the 1/4 position in the thickness direction was defined as the soft layer (surface soft layer), and the thickness in the thickness direction of this area was taken as the thickness of the soft layer.
  • the high-strength steel sheet according to an embodiment of the present invention preferably has a first plating layer, which is a metal plating layer, on one or both surfaces of the base steel sheet.
  • the first plating layer is formed directly on the surface of the base steel sheet and is a metal plating layer containing more than 50 mass% in total of one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb and Bi, and is not a hot-dip galvanized layer, an alloyed hot-dip galvanized layer, a zinc plating layer of an electrogalvanized layer, or a hot-dip aluminum plating layer.
  • the first plating layer is preferably a metal electroplated layer, and the following description will be given taking a metal electroplated layer as an example.
  • the outermost metal electroplating layer helps to prevent bending cracks during press forming and vehicle collisions, further improving bending fracture resistance.
  • the metal type of the metal electroplating layer may be any of Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi, but Fe is more preferable.
  • the coating weight of the Fe-based electroplating layer is more than 0 g/ m2 , and preferably 2.0 g/ m2 or more. There is no particular upper limit to the coating weight of the Fe-based electroplating layer per side, but from the viewpoint of cost, the coating weight of the Fe-based electroplating layer per side is preferably 60 g/ m2 or less.
  • the coating weight of the Fe-based electroplating layer is preferably 50 g/ m2 or less, more preferably 40 g/ m2 or less, and even more preferably 30 g/ m2 or less.
  • the adhesion weight of the Fe-based electroplating layer is measured as follows. A 10 x 15 mm sample is taken from the Fe-based electroplated steel sheet and embedded in resin to create a cross-section embedded sample. Three random locations on the cross section are observed using a scanning electron microscope (SEM) at an accelerating voltage of 15 kV and a magnification of 2,000 to 10,000 times depending on the thickness of the Fe-based plating layer, and the average thickness of the three fields of view is multiplied by the specific gravity of iron to convert it into the adhesion weight of the Fe-based plating layer per side.
  • SEM scanning electron microscope
  • the Fe-based electroplating layer in addition to pure Fe, alloy plating layers such as Fe-B alloy, Fe-C alloy, Fe-P alloy, Fe-N alloy, Fe-O alloy, Fe-Ni alloy, Fe-Mn alloy, Fe-Mo alloy, and Fe-W alloy can be used.
  • the composition of the Fe-based electroplating layer is not particularly limited, but it is preferable that the composition contains one or more elements selected from the group consisting of B, C, P, N, O, Ni, Mn, Mo, Zn, W, Pb, Sn, Cr, V, and Co in a total amount of 10 mass% or less, with the remainder consisting of Fe and unavoidable impurities.
  • the C content is preferably 0.08 mass% or less.
  • the high-strength steel sheet according to one embodiment of the present invention may have a second plating layer, which is a metal plating layer, as an outermost layer on one or both sides of the high-strength steel sheet.
  • the second plating layer contains at least one of zinc and aluminum in a total amount of 50 mass % or more, and may be a hot-dip galvanized layer, a galvannealed layer, an electrolytic galvanized layer, a hot-dip aluminum plating layer, or the like.
  • the second plating layer may be formed directly on one or both surfaces of the base steel sheet surface, or it may be formed on the first plating layer.
  • the hot-dip galvanized layer, alloyed hot-dip galvanized layer, and electrolytic galvanized layer refer to a plating layer containing Zn (zinc) as a main component (Zn content of 50.0 mass % or more).
  • the plating layer of an aluminum-plated steel sheet refers to a plating layer containing Al (aluminum) as a main component (Al content is 50.0 mass % or more).
  • the hot-dip galvanized layer is preferably composed of, for example, Zn, 20.0 mass% or less Fe, and 0.001 mass% or more and 1.0 mass% or less Al.
  • the hot-dip galvanized layer may optionally contain one or more elements 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 more than 0.0 mass% and 3.5 mass% or less.
  • the Fe content of the hot-dip galvanized layer is more preferably less than 7.0 mass%. The remainder other than the above elements is unavoidable impurities.
  • the galvannealed layer is preferably composed of, for example, 20% by mass or less Fe and 0.001% by mass or more and 1.0% by mass or less Al.
  • the galvannealed layer may optionally contain one or more elements 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 more than 0% by mass and 3.5% by mass or less.
  • the Fe content of the galvannealed layer is more preferably 7.0% by mass or more, and even more preferably 8.0% by mass or more.
  • the Fe content of the galvannealed layer is more preferably 15.0% by mass or less, and even more preferably 12.0% by mass or less. The remainder other than the above elements is unavoidable impurities.
  • the plating weight of the zinc plating layer per side is not particularly limited, but is preferably 20 g/m 2 or more and 80 g/m 2 or less.
  • the coating weight of the zinc plating layer is measured as follows.
  • a treatment solution is prepared by adding 0.6 g of a corrosion inhibitor for Fe (Ivit 700BK (registered trademark) manufactured by Asahi Chemical Industry Co., Ltd.) to 1 L of a 10 mass % aqueous hydrochloric acid solution.
  • a sample of a steel sheet having a zinc plating layer is then immersed in the treatment solution to dissolve the zinc plating layer.
  • the mass loss of the sample before and after dissolution is then measured, and the value is divided by the surface area of the base steel sheet (the surface area of the part that was covered with plating) to calculate the coating weight (g/ m2 ).
  • the thickness of the high-strength steel plate is not particularly limited, and can be 0.3 mm or more and 3.0 mm or less.
  • the production method of the present invention is also a method for producing the high-strength steel plate according to the present invention described above.
  • the temperature in the production method is based on the surface temperature of the steel slab or steel plate, unless otherwise specified.
  • ⁇ Hot rolling, pickling and cold rolling> In the manufacturing method of the present invention, first, a steel slab having the above-mentioned component composition of the present invention is subjected to hot rolling, pickling and cold rolling to obtain a cold-rolled sheet.
  • the steel slab may be, for example, a molten steel having the composition of the present invention obtained by melting a steel material and solidifying the molten steel.
  • the method for melting steel is not particularly limited, and known melting methods such as converter molten steel and electric furnace molten steel can be used.
  • the method for producing a steel slab from molten steel is not particularly limited, and known methods such as continuous casting, ingot casting, and thin slab casting can be used. From the viewpoint of preventing macrosegregation, it is preferable to produce the steel slab by the continuous casting method.
  • the produced steel slab is, for example, cooled to room temperature once, then heated again and hot rolled (rough rolling and finish rolling), and then coiled. In this way, a hot-rolled sheet is obtained.
  • the produced steel slab may be charged into a heating furnace as a hot piece without being cooled to room temperature, or may be roughly rolled immediately after being slightly kept at room temperature.
  • the temperature at which the steel slab is heated is preferably 1100° C. or higher from the viewpoint of dissolving carbides and reducing the rolling load.
  • the slab heating temperature is preferably 1300° C. or lower. The steel slab heated to the slab heating temperature is subjected to rough rolling.
  • the standard deviation of nanohardness can be reduced to 0.60 ⁇ [H n ] ave or less. This is thought to be because solute atoms such as Si and Mn rapidly diffuse through dislocations and grain boundaries of recrystallized grains during the plastic deformation and dynamic recrystallization of austenite grains during rough rolling, resulting in an appropriate distribution and uniform plastic deformation resistance in local regions.
  • the average strain rate during rough rolling is defined as the total rolling reduction ⁇ (-) from the first mill to the final mill in rough rolling divided by the time t R (s) required from the start of rolling at the first mill to the completion of rolling at the final mill in rough rolling ( ⁇ /t R ).
  • the average strain rate during rough rolling is less than 1 ⁇ 10 ⁇ 4 /s, the recovery of dislocations in austenite grains is promoted, the driving force for recrystallization is reduced, and dynamic recrystallization is suppressed, resulting in insufficient diffusion of solute atoms such as Si and Mn, and the standard deviation of nanohardness exceeds 0.60 ⁇ [H n ] ave . Therefore, rough rolling is performed under the conditions of an average strain rate of 1 ⁇ 10 ⁇ 4 /s to 1 ⁇ 10 ⁇ 1 /s and a total reduction of 50% or more.
  • the average strain rate during rough rolling is preferably 1 ⁇ 10 ⁇ 3 /s to 1 ⁇ 10 ⁇ 2 /s.
  • the total reduction during rough rolling is preferably 60% or more.
  • the rough rolling end temperature is preferably set to 950° C. or more.
  • the rough rolling end temperature can be set to, for example, 1250° C. or less.
  • the slab heating temperature is set to a low temperature, it is preferable to heat the roughly rolled sheet using a bar heater or the like before finish rolling in order to prevent problems during hot rolling.
  • the temperature when performing the finish rolling is preferably equal to or higher than the Ar3 transformation point. This reduces the rolling load. Furthermore, the rolling reduction in the non-recrystallized state of austenite is reduced, and the development of abnormal structures elongated in the rolling direction is suppressed, resulting in excellent workability.
  • Finish rolling may be performed continuously by joining the rough rolled sheets together.
  • the rough rolled sheets may be wound up once before finishing rolling is performed.
  • Lubricated rolling is also preferred from the viewpoint of making the steel sheet shape and material uniform.
  • the friction coefficient during lubricated rolling is preferably in the range of 0.10 to 0.25.
  • the coiling temperature after the hot rolling is preferably 300° C. or more and 700° C. or less from the viewpoint of improving the sheet passing property during the cold rolling and annealing described later.
  • the hot-rolled sheet obtained by hot rolling is pickled.
  • oxides on the surface of the hot-rolled sheet are removed, and excellent chemical conversion treatability and quality of the plating layer can be obtained in the final product, a high-strength steel sheet.
  • Pickling may be performed once or multiple times.
  • the hot-rolled sheet is optionally subjected to a softening heat treatment and then cold-rolled. In this way, a cold-rolled sheet is obtained.
  • the total reduction ratio of the cold rolling is preferably 20% or more and 75% or less.
  • the number of rolling passes and the reduction ratio of each pass are no particular limitations on the number of rolling passes and the reduction ratio of each pass.
  • a first plating step may be included in which a first plating layer, which is a metal plating layer, is formed on one or both sides of the steel sheet after the hot rolling step (after the cold rolling step if cold rolling is performed) and before the annealing step.
  • the first plating step is preferably a metal electroplating step.
  • a metal plating process such as metal electroplating may be applied to the surface of the cold-rolled sheet obtained as described above to produce a pre-annealed metal-plated steel sheet having a pre-annealed metal plating layer formed on at least one side.
  • the metal plating layer referred to here may be the above-mentioned first plating layer.
  • the pre-annealed metal-plated steel sheet is preferably a pre-annealed metal electroplated steel sheet provided with a pre-annealed metal electroplating layer.
  • the metal electroplated steel sheet before annealing means that the metal electroplating layer has not been subjected to an annealing process, and does not exclude the case where the hot-rolled sheet before the metal electroplating process, the pickled sheet after hot rolling, or the cold-rolled sheet has been annealed in advance.
  • the metal species of the electroplating layer can be any of Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi, but since Fe is more preferable, the manufacturing method of Fe-based electroplating is described below.
  • the Fe ion content in the Fe-based electroplating bath before the start of energization is preferably 0.5 mol/L or more in terms of Fe 2+ . If the Fe ion content in the Fe-based electroplating bath is 0.5 mol/L or more in terms of Fe 2+ , a sufficient Fe deposition amount can be obtained. In addition, in order to obtain a sufficient Fe deposition amount, the Fe ion content in the Fe-based electroplating bath before the start of energization is preferably 2.0 mol/L or less.
  • the Fe-based electroplating bath may contain at least one element selected from the group consisting of B, C, P, N, O, Ni, Mn, Mo, Zn, W, Pb, Sn, Cr, V, and Co in addition to Fe ions.
  • the total content of these elements in the Fe-based electroplating bath is preferably such that the total content of these elements in the Fe-based electroplating layer before annealing is 10 mass% or less.
  • the metal elements may be contained as metal ions, and the nonmetal elements may be contained as part of boric acid, phosphoric acid, nitric acid, organic acid, etc.
  • the iron sulfate plating solution may also contain a conductivity aid such as sodium sulfate or potassium sulfate, a chelating agent, or a pH buffer.
  • the temperature of the Fe-based electroplating solution is preferably 30° C. or higher from the viewpoint of constant temperature retention, and is preferably 85° C. or lower.
  • the pH of the Fe-based electroplating bath is also not particularly limited, but is preferably 1.0 or higher from the viewpoint of preventing a decrease in current efficiency due to hydrogen generation, and is preferably 3.0 or lower from the viewpoint of the electrical conductivity of the Fe-based electroplating bath.
  • the current density is preferably 10 A/dm 2 or higher from the viewpoint of productivity, and is preferably 150 A/dm 2 or lower from the viewpoint of facilitating control of the deposition amount of the Fe-based electroplating layer.
  • the sheet passing speed is preferably 5 mpm or higher from the viewpoint of productivity, and is preferably 150 mpm or lower from the viewpoint of stably controlling the deposition amount.
  • a degreasing treatment and water washing for cleaning the surface of the cold-rolled sheet, and further, a pickling treatment and water washing for activating the surface of the cold-rolled sheet can be performed.
  • the Fe-based electroplating treatment is performed following these pretreatments.
  • the method of the degreasing treatment and water washing is not particularly limited, and a conventional method can be used.
  • various acids such as sulfuric acid, hydrochloric acid, nitric acid, and mixtures thereof can be used. Among them, sulfuric acid, hydrochloric acid, and mixtures thereof are preferred.
  • the concentration of the acid is not particularly limited, but is preferably 1% by mass or more and 20% by mass or less from the viewpoints of the ability to remove the oxide film and the prevention of roughness (surface defects) due to over-pickling.
  • the pickling solution may also contain an antifoaming agent, a pickling promoter, a pickling inhibitor, and the like.
  • the obtained cold-rolled sheet is subjected to a first heating at a temperature of 750° C. or more.
  • the cold-rolled sheet may be one that has been subjected to an electroplating treatment, but may not be one that has been subjected to said treatment.
  • the first heating temperature is 750° C. or higher, and preferably 770° C. or higher.
  • the upper limit of the heating temperature is not particularly limited, but is preferably 950° C. or less from the viewpoint of operability and the like.
  • heating time is not particularly limited, but if the time is too short, the reverse transformation to austenite may not proceed sufficiently, so the time is preferably 30 seconds or more, and more preferably 60 seconds or more.
  • the upper limit of the heating time is not particularly limited, and can be, for example, 6000 s or less, and preferably 3000 s or less, where "s" means seconds.
  • the dew point of the annealing atmosphere in the first heating is preferably -30°C or higher.
  • the annealing atmosphere in the annealing process is more preferably -15°C or higher, and even more preferably -5°C or higher.
  • the dew point of the annealing atmosphere in the annealing process is preferably 30°C or lower.
  • the first average cooling rate v1 is 2.0° C./s or more, preferably 3.0° C./s or more, and more preferably 5.0° C./s or more.
  • the upper limit of the first average cooling rate v1 is not particularly limited, but from the viewpoint of reducing the capital investment burden, it is preferably 60.0° C./s or less.
  • the cooling in the temperature region T1 is preferably continuous cooling.
  • the cooling rate from the first heating temperature to 750° C. is not particularly limited.
  • the cold-rolled sheet that has passed through the temperature range T1 is then subjected to a residence time step of holding at a residence temperature T2 of 350°C or more and 550°C or less.
  • the cold-rolled sheet that has passed through the first cooling step is then held at a residence temperature T2 of 350°C or more and 550°C or less for a residence time t2 (s) that satisfies F defined by formula 1 of 0.20 or more and 0.90 or less, thereby causing bainite transformation.
  • this residence time t2 (s) from an expansion curve obtained from a Formaster test, the bainite transformation and the associated precipitation of iron carbides and the distribution of C into untransformed austenite are optimized. Since the expansion curve depends on the steel composition and the thermal history up to the first cooling, it is necessary to determine an expansion curve for each steel composition and thermal history from the first heating temperature to T2 °C, and select an appropriate residence time t2 .
  • the holding time t2 is the time t or more at which F becomes 0.20, and preferably the time t or more at which F becomes 0.30.
  • the bainite transformation proceeds excessively, the amount of martensite decreases, and TS decreases.
  • the holding time t2 is equal to or shorter than the time t at which F becomes 0.90, and preferably equal to or shorter than the time t at which F becomes 0.80.
  • the relationship between F and t in formula 1 is calculated as follows.
  • the steel slab is subjected to a process up to the first cooling, and then is retained at a retention temperature T2 of 350° C. to 550° C.
  • the process up to the first cooling includes a hot rolling process, a pickling and cold rolling process, a first heating process, and a process up to the first cooling.
  • an expansion curve during retention at retention temperature T2 is obtained. Retention at T2 is continued until expansion stops.
  • the expansion amount at the start of retention at retention temperature T2 is set to 0, and the expansion amount at retention is set to 1.
  • the expansion curve is fitted with Equation 1 to calculate constants k and n. This determines the relationship between F and t at retention temperature T2 .
  • Formula 1: F 1-exp(-kt n ) t: residence time (s) k, n: constants obtained from the expansion curve of the Formaster test
  • the cooling stop temperature is Ms-20°C or less. This allows the martensitic transformation to proceed sufficiently. If the cooling stop temperature exceeds Ms-20°C, the untransformed austenite does not transform to martensite, the amount of retained austenite becomes excessive, and good part strength and stretch flangeability cannot be obtained.
  • the cooling stop temperature may be room temperature.
  • Ms is the temperature (Ms point) at which martensitic transformation begins to occur, and a value measured by the test described below is used.
  • the second average cooling rate v2 is 5° C./s or more, and preferably 8° C./s or more.
  • the upper limit of the second average cooling rate v2 is not particularly limited, but from the viewpoint of reducing the capital investment burden, it is preferably 60.0° C./s or less.
  • the cooling rate outside the temperature region T3 is not particularly limited.
  • the Ms point is a value measured by a Formaster test as follows. Using a Formaster testing machine, the steel slab is subjected to a process up to the end of the retention process, and then cooled to room temperature at a second average cooling rate of 5°C/s or more. The temperature at which martensitic transformation occurs and expansion begins during the second cooling is defined as the Ms point.
  • the upper limit of the second average cooling rate is not particularly limited, but can be, for example, 100°C/s or less.
  • the ratio (ratio p) of the number of martensite blocks in which metastable carbides exist to the total number of martensite blocks can be increased, thereby improving low temperature toughness.
  • the temperature X (° C.) is a value higher than room temperature.
  • the temperature X (° C.) of the second heating is higher than the cooling stop temperature of the second cooling step, heating is performed from the cooling stop temperature to the temperature X (° C.) of the second heating.
  • the cooling stop temperature of the second cooling step and the temperature X (°C) of the second heating may be the same, in which case it means that the cooling stop temperature is maintained at the temperature X (°C) of the second heating.
  • the value of the variable part Z in the above formula 2 is too small, that is, if the temperature X is too low and/or the holding time Y is too short, metastable carbides are not sufficiently precipitated, and the proportion p becomes low. Therefore, from the viewpoint of increasing the proportion p, the value of the variable part Z is 7000 or more, and preferably 8000 or more. On the other hand, if the value of the variable part Z is too high, i.e., if the temperature X is too high and/or the holding time Y is too long, the metastable carbides will transition to cementite and the proportion p will decrease. For this reason, from the viewpoint of increasing the proportion p, the value of the variable part Z is 13000 or less, and preferably 12000 or less.
  • the temperature X (unit: ° C.) preferably satisfies the following formula 3. This increases the number density (number density n) of metastable carbides in the martensite block in which the metastable carbides are present.
  • Formula 3 100 ⁇ X ⁇ 400
  • the temperature X is preferably 100° C. or higher, more preferably 120° C. or higher, and even more preferably 150° C. or higher.
  • the temperature X is preferably 400° C. or less, more preferably 380° C. or less, and further preferably 350° C. or less.
  • hot-dip galvanizing it is preferable to immerse the steel sheet in a zinc plating bath at 440°C to 500°C, and then adjust the coating weight by gas wiping or the like.
  • a plating bath with an Al content of 0.10 mass% to 0.23 mass%, with the balance being Zn and unavoidable impurities.
  • the coating weight of the steel sheet having a hot-dip galvanized layer (hot-dip galvanized steel sheet) (GI) and the steel sheet having a galvannealed layer (galvannealed hot-dip galvanized steel sheet) (GA) is preferably 20 to 80 g/ m2 per side.
  • the coating weight can be adjusted by gas wiping or the like.
  • Skin pass rolling (optional)
  • the obtained high strength steel sheet may be subjected to skin pass rolling.
  • the plating treatment may be performed.
  • the reduction ratio in the skin pass rolling is preferably 0.05% or more from the viewpoint of increasing the yield strength.
  • the upper limit of the reduction ratio is not particularly limited, but is preferably 1.50% from the viewpoint of productivity.
  • Skin pass rolling may be performed either online or offline. The skin pass may be performed at a single time to the desired rolling reduction, or may be performed in several steps.
  • Example 1 the sheet was held at the first heating temperature of 800 ° C. for 200 s, then cooled to 480 ° C. under the condition of an average cooling rate of 18 ° C./s, and then held at 480 ° C. for 1000 s.
  • the expansion curve obtained was fitted with Equation 1 to obtain k and n.
  • ⁇ Low temperature toughness test> Six of the obtained steel plates (1.6 mmt steel plates) were stacked and bonded together to prepare a Charpy impact test piece with a thickness of 9.6 mmt. The notch was a 2 mm U-notch. Using this test piece, a Charpy impact test was performed at -40°C, and the Charpy absorbed energy obtained was measured. In addition, the ductile fracture rate was measured by observing the fracture surface after the test.
  • the low-temperature toughness parameter P was defined as the product of Charpy absorbed energy (unit: J) and ductile fracture surface area ratio (%) (Charpy absorbed energy (J) x ductile fracture surface area ratio (%)) and was calculated. When the low temperature toughness parameter P is 3000 or more, it is determined that the low temperature toughness is excellent.
  • TS is 780 MPa or more, and the parts have excellent strength, ductility, stretch flangeability, bendability of the sheared end surface, and low-temperature toughness.
  • the parts are inferior in one or more of strength, ductility, stretch flangeability, bendability of the sheared end surface, and low-temperature toughness.
  • the composition of the metal electroplating layer is Fe: 95-100% by mass for Fe-based electroplating, Ni: 95-100% by mass for Ni-based electroplating, and the remainder in each case is unavoidable impurities.
  • a punch B2 was pressed into the test piece T1 placed on a support roll A2 so that the bending direction was perpendicular to the rolling direction, and an orthogonal bending (secondary bending) was performed.
  • D1 indicates the width (C) direction
  • D2 indicates the rolling (L) direction.

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Abstract

The present invention provides a high-strength steel sheet having excellent component strength, ductility, stretch flange-ability, shear end face section bendability, and low-temperature toughness. The high-strength steel sheet has a prescribed component composition and a steel composition in which the martensite area ratio is 10% to 80%, the bainite area ratio is 2% to 70%, the ferrite area ratio is 80% or less, the residual austenite area ratio is 15% or less, and the ratio of the number of martensite blocks in which metastable carbide exists to the number of martensite blocks is 2% or more. When the nanohardness is measured at 225 or more points at a position 1/4 sheet thickness, the standard deviation σn of the nanohardness to the average nanohardness [Hn]ave is 0.60×[Hn]ave or less.

Description

高強度鋼板、高強度鋼板を用いてなる部材、部材からなる自動車の骨格構造部品用又は自動車の補強部品、ならびに高強度鋼板及び部材の製造方法High-strength steel plate, member made of high-strength steel plate, structural component for automobile or reinforcing component for automobile made of member, and manufacturing method of high-strength steel plate and member
　本発明は、高強度鋼板、高強度鋼板を用いてなる部材、部材からなる自動車の骨格構造部品用又は自動車の補強部品、ならびに高強度鋼板及び部材の製造方法に関する。 The present invention relates to high-strength steel plates, components made of high-strength steel plates, automobile structural components or automobile reinforcing parts made of components, and methods for manufacturing high-strength steel plates and components.
　車輌の軽量化によるＣＯ２排出量削減を図りつつ、耐衝突性能を向上させることを目的として、自動車用鋼板の高強度化が進められている。また、新たな法規制の導入が相次いでいることを背景に、車体強度の増加を目的として、自動車キャビンの骨格を形成する主要な構造部品や補強部品（以下、「自動車の骨格構造部品」ともいう）に対して、引張強さ（ＴＳ）が７８０ＭＰａ以上の高強度鋼板を適用する事例の増加が顕著である。 Strengthening of steel sheets for automobiles is being promoted for the purpose of improving crashworthiness while reducing CO2 emissions by reducing the weight of vehicles. In addition, against the background of the successive introduction of new laws and regulations, there has been a remarkable increase in the number of cases in which high-strength steel sheets with a tensile strength (TS) of 780 MPa or more are used for the main structural parts and reinforcing parts that form the framework of an automobile cabin (hereinafter also referred to as "automobile framework structural parts") in order to increase the strength of the vehicle body.
　例えば、特許文献１では、所定の成分組成及びミクロ組織を有し、低温変態相に含まれる鉄炭化物の粒径が５００ｎｍ以下である、高強度鋼板が提案されている。 For example, Patent Document 1 proposes a high-strength steel plate having a specified component composition and microstructure, and in which the grain size of the iron carbide contained in the low-temperature transformation phase is 500 nm or less.
特開２００８－３０８７１７号公報JP 2008-308717 A
　自動車の骨格構造部品に用いられる高強度鋼板は、部品強度に優れる（衝突時における衝撃吸収エネルギーが大きい）ことが要求され、鋼板の降伏強さ（ＹＳ）が高いこと、降伏比（ＹＲ＝降伏強さ（ＹＳ）／引張強さ（ＴＳ）が高いことが求められる。また、クラッシュボックス等の部品は、打ち抜き端面や曲げ加工部を有するため、これらの部品に用いる鋼板には、良好な延性、伸びフランジ性及び剪断端面部の曲げ性が要求される。さらに、引張強さが７８０ＭＰａ以上の高強度鋼板を用いた部品を低温環境下で使用する場合には、靭性が悪化し、衝突時に割れが生じるおそれがあるところ、自動車用鋼板については、低温環境下で使用される場合の衝突時の割れを防ぐために、優れた低温靭性が要求される。 High-strength steel plates used in automotive structural components are required to have excellent component strength (high impact energy absorption during a collision), and are required to have a high yield strength (YS) and a high yield ratio (YR = yield strength (YS) / tensile strength (TS). In addition, parts such as crash boxes have punched end faces and bent parts, so the steel plates used in these parts are required to have good ductility, stretch flangeability, and bendability of the shear end faces. Furthermore, when parts using high-strength steel plates with a tensile strength of 780 MPa or more are used in low-temperature environments, there is a risk that their toughness will deteriorate and cracks will occur during a collision. Therefore, automotive steel plates are required to have excellent low-temperature toughness to prevent cracks during a collision when used in low-temperature environments.
　高強度鋼板を自動車部品に適用する比率を上げるため、上述した特性を総合的に満足する鋼板が依然として求められている。 In order to increase the proportion of high-strength steel sheets used in automotive parts, there is still a demand for steel sheets that comprehensively satisfy the above-mentioned properties.
　本発明は、以上の点を鑑みてなされたものであり、かつ、部品強度、延性、伸びフランジ性、剪断端面部の曲げ性及び低温靭性に優れる、高強度鋼板をその製造方法とともに提供することを目的とする。
　また、本発明は、上記の高強度鋼板を用いてなる部材をその製造方法とともに提供することを目的とする。 The present invention has been made in consideration of the above points, and an object of the present invention is to provide a high-strength steel plate having excellent part strength, ductility, stretch flangeability, bendability of a sheared end surface, and low-temperature toughness, together with a manufacturing method thereof.
Another object of the present invention is to provide a member made using the above-mentioned high-strength steel plate, together with a method for producing the member.
　ここで、「高強度鋼板」とは、後述する引張試験により求める引張強さ（ＴＳ）が７８０ＭＰａ以上の鋼板を意味する。
　「部品強度に優れる」とは、後述する引張試験により求める降伏比（ＹＲ）が５５％以上であることを意味する。
　「延性に優れる」とは、後述する引張試験により求める全伸び（Ｅｌ）が１０％以上であることを意味する。
　「伸びフランジ性に優れる」とは、後述する穴広げ試験により求める穴広げ率（λ）が２０％以上であることを意味する。
　「剪断端面部の曲げ性に優れる」とは、後述する剪断端面部を有するサンプルの曲げ試験で求められる限界曲げ半径（Ｒｓ／ｔ）と研削端面部を有するサンプルの曲げ試験で求められる限界曲げ半径（Ｒｇ／ｔ）の比（Ｒｓ／Ｒｇ）が１．５０以下を意味する。
　「低温靭性に優れる」とは、後述するシャルピー衝撃試験において、低温靭性パラメータ（Ｐ）が３０００以上であることを意味する。 Here, the term "high strength steel plate" refers to a steel plate having a tensile strength (TS) of 780 MPa or more as determined by a tensile test described later.
"Excellent component strength" means that the yield ratio (YR) determined by the tensile test described below is 55% or more.
The term "excellent ductility" means that the total elongation (El) determined by the tensile test described below is 10% or more.
The term "excellent stretch flangeability" means that the hole expansion ratio (λ) determined by the hole expansion test described below is 20% or more.
"Excellent bendability at the sheared end surface" means that the ratio (Rs/Rg) of the limit bending radius (Rs/t) determined in a bending test of a sample having a sheared end surface (described later) to the limit bending radius (Rg/t) determined in a bending test of a sample having a ground end surface is 1.50 or less.
The term "excellent low-temperature toughness" means that the low-temperature toughness parameter (P) is 3000 or more in the Charpy impact test described below.
　本発明者らは、鋭意検討した結果、下記構成を採用することにより、上記目的が達成されることを見出し、本発明を完成させた。
　すなわち、本発明の要旨は以下のとおりである。
（１）質量％で、
Ｃ：０．０３０％以上０．５００％以下、
Ｓｉ：０．０１％以上２．５０％以下、
Ｍｎ：０．１０％以上５．００％以下、
Ｐ：０．１００％以下、
Ｓ：０．０２００％以下、
Ａｌ：１．０００％以下、
Ｎ：０．０１００％以下及び
Ｏ：０．０１００％以下
を含有し、残部がＦｅ及び不可避的不純物からなる成分組成と、
　板厚１／４位置において、
マルテンサイトの面積率が１０％以上８０％以下、
ベイナイトの面積率が２％以上７０％以下、
フェライトの面積率が８０％以下、
残留オーステナイトの面積率が１５％以下、かつ
マルテンサイトブロック数に対する準安定炭化物が存在するマルテンサイトブロック数の割合が２％以上である鋼組織と、
を有し、
　板厚１／４位置において、２２５点以上のナノ硬度を測定したとき、ナノ硬度の平均値［Ｈｎ］aveに対して、ナノ硬度の標準偏差σnが０．６０×［Ｈｎ］ave以下である、
高強度鋼板。
（２）前記準安定炭化物が存在するマルテンサイトブロックにおける準安定炭化物の個数密度の平均値が１×１０6個/ｍｍ２以上である、上記（１）の高強度鋼板。
（３）前記成分組成は、さらに、質量％で、
Ｔｉ：０．２００％以下、
Ｎｂ：０．２００％以下、
Ｖ：０．２００％以下、
Ｔａ：０．１０％以下、
Ｗ：０．１０％以下、
Ｂ：０．０１００％以下、
Ｃｒ：１．００％以下、
Ｍｏ：１．００％以下、
Ｎｉ：１．００％以下、
Ｃｏ：０．０１０％以下、
Ｃｕ：１．００％以下、
Ｓｎ：０．２００％以下、
Ｓｂ：０．２００％以下、
Ｃａ：０．０１００％以下、
Ｍｇ：０．０１００％以下、
ＲＥＭ：０．０１００％以下、
Ｚｒ：０．１００％以下、
Ｔｅ：０．１００％以下、
Ｈｆ：０．１０％以下及び
Ｂｉ：０．２００％以下
からなる群より選ばれる少なくとも１種の元素を含有する、上記（１）又は（２）の高強度鋼板。
（４）前記高強度鋼板の板厚１／４位置のビッカース硬さに対して、ビッカース硬さが８５％以下の領域であって、前記高強度鋼板表面から板厚方向に２００μｍ以内の領域である表層軟質層を有し、
　前記高強度鋼板表面から前記表層軟質層の板厚方向深さの１／４位置及び板厚方向深さの１／２位置のそれぞれにおける板面の５０μｍ×５０μｍの領域において、３００点以上のナノ硬度を測定したとき、
　前記高強度鋼板表面から前記表層軟質層の板厚方向深さの１／４位置の板面のナノ硬度が７．０ＧＰａ以上の測定数割合が、全測定数に対して０．１０以下であり、
　前記高強度鋼板表面から前記表層軟質層の板厚方向深さの１／４位置の板面のナノ硬度の標準偏差σが１．８ＧＰａ以下であり、
　さらに、前記高強度鋼板表面から前記表層軟質層の板厚方向深さの１／２位置の板面のナノ硬度の標準偏差σが２．２ＧＰａ以下である、上記（１）～（３）のいずれかの高強度鋼板。
（５）前記高強度鋼板の片面又は両面の表面上において、Ｃｒ、Ｍｎ、Ｆｅ、Ｃｏ、Ｎｉ、Ｃｕ、Ｇａ、Ｇｅ、Ａｓ、Ｒｕ、Ｒｈ、Ｐｄ、Ａｇ、Ｃｄ、Ｉｎ、Ｓｎ、Ｓｂ、Ｏｓ、Ｉｒ、Ｒｔ、Ａｕ、Ｈｇ、Ｔｉ、Ｐｂ及びＢｉから選択される１種又は２種以上を合計で５０質量％超含む金属めっき層を有する、上記（１）～（４）のいずれかの高強度鋼板。
（６）前記高強度鋼板の片面又は両面の最外層に、亜鉛及びアルミニウムの少なくとも一方を合計で５０質量以上含む金属めっき層を有する、上記（１）～（５）のいずれかの高強度鋼板。
（７）上記（１）～（６）のいずれかの高強度鋼板を用いてなる、部材。
（８）上記（７）の部材からなる、自動車の骨格構造部品用又は自動車の補強部品。
（９）上記（１）又は（３）に記載の成分組成を有する鋼スラブに、
　平均のひずみ速度が１×１０－４／s以上１×１０－１／s以下、総圧下率５０％以上の条件で粗圧延を施した後、仕上げ圧延を施し、次いで巻取り処理を施して、熱延板を得る熱間圧延工程、
　次いで、酸洗及び冷間圧延を施して、冷延板を得る酸洗及び冷間圧延工程と、
　次いで、加熱温度が７５０℃以上の条件で第１加熱する第１加熱工程と、
　次いで、Ｔ２以上７５０℃以下の温度域における第１冷却速度が２．０℃／ｓ以上の条件で冷却する第１冷却工程と、
　次いで、３５０℃以上５５０℃以下の滞留温度Ｔ２で下記式１で定義されるＦが０．２０以上０．９０以下を満たす滞留時間ｔ（ｓ）の条件で滞炉させる滞炉工程と、
　次いで、Ｍｓ－２０℃以下まで冷却する工程であって、Ｍｓ－２０℃以上Ｍs以下の温度域における第２平均冷却速度を５℃／ｓ以上の条件とする第２冷却工程と、
　次いで、下記式２を満たす温度Ｘ（℃）と保持時間Ｙ（ｓ）の条件で処理する第２加熱工程と
を含む、高強度鋼板の製造方法。
　　　　　　　　　　　　　　　　　記
式１：Ｆ＝１－ｅｘｐ（－ｋｔｎ）
ｔ：滞留時間（ｓ）
ｋ、ｎ：前記スラブを第１冷却工程終了までの工程に付して得られる試験片に対して、３５０℃以上５５０℃以下の滞留温度Ｔ２で保持させて行われるフォーマスター試験の膨張曲線から求められる定数。
式２：７０００≦（２７３＋Ｘ）（２０＋ｌｏｇ(Ｙ/３６００)）≦１３０００
（１０）前記第２加熱工程において、温度Ｘ（℃）が下記式３を満たす、上記（９）の高強度鋼板の製造方法。
　　　　　　　　　　　　　　　　　記
式３：１００≦Ｘ≦４００
（１１）前記第１加熱工程を露点－３０℃以上の雰囲気下で行う，上記（９）又は（１０）の高強度鋼板の製造方法。
（１２）前記冷間圧延工程後、かつ前記焼鈍工程の前の鋼板の片面もしくは両面において、Ｃｒ、Ｍｎ、Ｆｅ、Ｃｏ、Ｎｉ、Ｃｕ、Ｇａ、Ｇｅ、Ａｓ、Ｒｕ、Ｒｈ、Ｐｄ、Ａｇ、Ｃｄ、Ｉｎ、Ｓｎ、Ｓｂ、Ｏｓ、Ｉｒ、Ｒｔ、Ａｕ、Ｈｇ、Ｔｉ、Ｐｂ及びＢｉから選択される１種又は２種以上を５０質量％超含む金属めっきを施す工程を含む、上記（９）～（１１）のいずれかの高強度鋼板の製造方法。
（１３）前記第１加熱から第２加熱工程の鋼板に、亜鉛及びアルミニウムの少なくとも一方を合計で５０質量％以上含む金属めっきを施す工程を含む、上記（９）～（１２）のいずれかの製造方法。
（１４）上記（１）～（６）のいずれかの高強度鋼板に、成形加工又は接合加工の少なくとも一方を施して部材とする工程を有する、部材の製造方法。 As a result of extensive investigation, the present inventors have found that the above object can be achieved by employing the following configuration, and have completed the present invention.
That is, the gist of the present invention is as follows.
(1) In mass%,
C: 0.030% or more and 0.500% or less,
Si: 0.01% or more and 2.50% or less,
Mn: 0.10% or more and 5.00% or less,
P: 0.100% or less,
S: 0.0200% or less,
Al: 1.000% or less,
A composition comprising N: 0.0100% or less and O: 0.0100% or less, with the balance being Fe and unavoidable impurities;
At the 1/4 plate thickness position,
The area ratio of martensite is 10% or more and 80% or less,
The area ratio of bainite is 2% or more and 70% or less,
The area ratio of ferrite is 80% or less,
a steel structure in which the area ratio of retained austenite is 15% or less and the ratio of the number of martensite blocks in which metastable carbides are present to the number of martensite blocks is 2% or more;
having
When the nanohardness is measured at 225 or more points at the 1/4 position of the sheet thickness, the standard deviation σ n of the nanohardness is 0.60 × [H n ] ave or less with respect to the average value [H n ] ave of the nanohardness.
High strength steel plate.
(2) The high-strength steel plate according to (1) above, wherein the average number density of metastable carbides in the martensite block in which the metastable carbides are present is 1 x 10 6 pieces/mm 2 or more.
(3) The composition further includes, in mass%,
Ti: 0.200% or less,
Nb: 0.200% or less,
V: 0.200% or less,
Ta: 0.10% or less,
W: 0.10% or less,
B: 0.0100% or less,
Cr: 1.00% or less,
Mo: 1.00% or less,
Ni: 1.00% or less,
Co: 0.010% or less,
Cu: 1.00% 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: 0.100% or less,
Te: 0.100% or less,
The high-strength steel plate according to (1) or (2) above, containing at least one element selected from the group consisting of Hf: 0.10% or less and Bi: 0.200% or less.
(4) The high-strength steel plate has a Vickers hardness of 85% or less of the Vickers hardness at a 1/4 position in the plate thickness direction of the high-strength steel plate, and has a surface soft layer which is a region within 200 μm from the surface of the high-strength steel plate in the plate thickness direction;
When the nano hardness was measured at 300 points or more in a 50 μm × 50 μm region of the sheet surface at a 1/4 position and a 1/2 position of the sheet thickness direction depth of the soft surface layer from the surface of the high strength steel sheet,
The ratio of the number of measurements in which the nano hardness of the sheet surface at a position of 1/4 of the sheet thickness direction depth of the soft surface layer from the surface of the high strength steel sheet is 7.0 GPa or more to the total number of measurements is 0.10 or less,
The standard deviation σ of the nano-hardness of the sheet surface at a ¼ position of the sheet thickness direction depth of the soft surface layer from the surface of the high-strength steel sheet is 1.8 GPa or less;
Further, the high-strength steel plate according to any one of (1) to (3) above, wherein the standard deviation σ of the nano-hardness of the plate surface at a position half the depth in the plate thickness direction of the soft surface layer from the surface of the high-strength steel plate is 2.2 GPa or less.
(5) The high-strength steel plate according to any one of (1) to (4) above, having a metal plating layer on one or both surfaces of the high-strength steel plate, the metal plating layer containing more than 50 mass% in total of one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi.
(6) The high-strength steel plate according to any one of (1) to (5) above, having a metal plating layer containing at least one of zinc and aluminum in a total amount of 50% by mass or more on one or both outermost layers of the high-strength steel plate.
(7) A member made using the high-strength steel plate according to any one of (1) to (6) above.
(8) An automobile frame structural part or automobile reinforcing part, comprising the member according to (7) above.
(9) A steel slab having the composition described in (1) or (3) above,
a hot rolling process in which rough rolling is performed under conditions of an average strain rate of 1×10 −4 /s or more and 1×10 −1 /s or less and a total rolling reduction of 50% or more, followed by finish rolling and then coiling to obtain a hot-rolled sheet;
Next, pickling and cold rolling are performed to obtain a cold-rolled sheet.
Next, a first heating step of performing a first heating under a condition of a heating temperature of 750° C. or more;
Next, a first cooling step in which the first cooling rate in a temperature range from T2 to 750°C is 2.0°C/s or more;
Next, a retention step of retaining the material in the furnace at a retention temperature T2 of 350° C. or more and 550° C. or less for a retention time t (s) in which F defined by the following formula 1 is 0.20 or more and 0.90 or less;
Next, a second cooling step of cooling to Ms-20°C or less, in which a second average cooling rate in a temperature range of Ms-20°C or more and Ms or less is set to 5°C/s or more;
Next, a second heating step is performed under conditions of a temperature X (° C.) and a holding time Y (s) that satisfy the following formula 2.
Formula 1: F=1-exp(-kt n )
t: residence time (s)
k, n: Constants determined from the expansion curve of a Formaster test performed on a test piece obtained by subjecting the slab to the process up to the end of the first cooling process and holding it at a residence temperature T2 of 350°C or more and 550°C or less.
Formula 2: 7000≦(273+X)(20+log(Y/3600))≦13000
(10) The method for producing a high strength steel plate according to (9) above, wherein in the second heating step, the temperature X (°C) satisfies the following formula 3:
Formula 3: 100≦X≦400
(11) The method for producing a high strength steel plate according to (9) or (10) above, wherein the first heating step is carried out in an atmosphere with a dew point of −30° C. or higher.
(12) A method for producing a high-strength steel sheet according to any one of (9) to (11) above, comprising a step of applying a metal plating containing more than 50 mass% of one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi to one or both sides of the steel sheet after the cold rolling step and before the annealing step.
(13) Any of the manufacturing methods according to (9) to (12) above, comprising a step of applying metal plating containing 50 mass% or more in total of at least one of zinc and aluminum to the steel sheet subjected to the first heating and the second heating steps.
(14) A method for manufacturing a component, comprising the step of subjecting any one of the high-strength steel plates according to (1) to (6) above to at least one of forming and joining to form a component.
　本発明によれば、部品強度、延性、伸びフランジ性、剪断端面部の曲げ性及び低温靭性に優れる高強度鋼板を提供することができる。また、上記の高強度鋼板を用いてなる部材を提供することができる。
　さらに、本発明によれば、上記の高強度鋼板及び高強度鋼板を用いてなる部材の製造方法を提供することができる。
　加えて、本発明によれば、上記の部材からなる部材からなる自動車の骨格構造部品用又は自動車の補強部品を提供することができる。 According to the present invention, it is possible to provide a high-strength steel sheet excellent in part strength, ductility, stretch flangeability, bendability of a sheared end surface, and low-temperature toughness. Also, it is possible to provide a member made using the high-strength steel sheet.
Furthermore, according to the present invention, there can be provided a method for manufacturing the above-mentioned high-strength steel plate and a method for manufacturing a member using the high-strength steel plate.
In addition, according to the present invention, it is possible to provide an automobile frame structural part or an automobile reinforcing part made of the above-mentioned member.
炭化物が存在するマルテンサイトの電子回折図形の一例である。This is an example of an electron diffraction pattern of martensite in which carbides are present. 実施例のＶ曲げ＋直交ＶＤＡ曲げ試験用サンプルの作製に関する模式図である。図２（ａ）はＶ曲げ加工（一次曲げ加工）に関し、図２（ｂ）は直交ＶＤＡ曲げ（二次曲げ加工）に関する。2A and 2B are schematic diagrams showing the preparation of samples for V-bending and orthogonal VDA bending tests in the examples, in which Fig. 2A shows V-bending (primary bending) and Fig. 2B shows orthogonal VDA bending (secondary bending). 実施例の軸圧壊試験用サンプル及び試験に関する模式図である。図３（ａ）は試験用部材の正面図であり、図３（ｂ）は試験用部材の正面図である。図３（ｃ）は軸圧壊試験を示す概略図である。3(a) is a front view of a test member, FIG. 3(b) is a front view of a test member, and FIG. 3(c) is a schematic diagram showing an axial crush test.
　本発明を、以下の実施形態に基づき説明する。本発明は、以下の実施形態に限定されない。 The present invention will be described based on the following embodiments. The present invention is not limited to the following embodiments.
［高強度鋼板］
　本発明の高強度鋼板（以下、便宜的に、「鋼板」ともいう。）は、後述する成分組成及び鋼組織を有する。 [High-strength steel plate]
The high-strength steel plate of the present invention (hereinafter, for convenience, also referred to as "steel plate") has a chemical composition and a steel structure described below.
〈成分組成〉
　本発明の高強度鋼板の成分組成（以下、便宜的に、「本発明の成分組成」ともいう）について、説明する。本発明の成分組成における「％」は、特に明記しない限り「質量％」を意味する。 <Component composition>
The composition of the high strength steel plate of the present invention (hereinafter, for convenience, also referred to as the "composition of the present invention") will be described. "%" in the composition of the present invention means "mass%" unless otherwise specified.
《Ｃ：０．０３０％以上０．５００％以下》
　Ｃは、鋼の重要な基本成分の１つであり、特に本発明においては、マルテンサイトの面積率に影響する。Ｃ含有量が少なすぎると、マルテンサイトの面積率が減少し、７８０ＭＰａ以上のＴＳを実現することが困難になる。このため、Ｃ含有量は、０．０３０％以上であり、０．０４０％以上が好ましく、０．０５０％以上がより好ましい。
　一方、Ｃ含有量が多すぎると、残留オーステナイトが過度に増加し、打ち抜き時に残留オーステナイトから生成するマルテンサイトの硬度が大きく上昇する。その結果、穴広げ時の亀裂進展が促進され、穴広げ率が低下して、伸びフランジ性が低下する。また、残留オーステナイトが応力誘起変態することによってＹＲが低下し、部品強度が低下する。このため、Ｃ含有量は、０．５００％以下であり、０．４００％以下が好ましく、０．３００％以下がより好ましい。 <<C: 0.030% or more and 0.500% or less>>
C is one of the important basic components of steel, and in particular in the present invention, it affects the area ratio of martensite. If the C content is too low, the area ratio of martensite decreases, making it difficult to achieve a TS of 780 MPa or more. For this reason, the C content is 0.030% or more, preferably 0.040% or more, and more preferably 0.050% or more.
On the other hand, if the C content is too high, the amount of retained austenite increases excessively, and the hardness of martensite generated from the retained austenite during punching increases significantly. As a result, crack propagation during hole expansion is promoted, the hole expansion ratio decreases, and the stretch flangeability decreases. In addition, the residual austenite undergoes stress-induced transformation, which reduces the YR and reduces the strength of the part. For this reason, the C content is 0.500% or less, preferably 0.400% or less, and more preferably 0.300% or less.
《Ｓｉ：０．０１％以上２．５０％以下》
　Ｓｉは、マルテンサイト中のセメンタイトの析出抑制や固溶強化によって、鋼板の強度を上昇させる成分である。この効果を得るため、Ｓｉ含有量は、０．０１％以上であり、０．０５％以上が好ましく、０．１０％以上がより好ましい。
　一方、Ｓｉ含有量が多すぎると、ベイナイト変態時の炭化物析出が著しく抑制され、残留オーステナイトが過度に増加し、打ち抜き時に残留オーステナイトから生成するマルテンサイトの硬度が大きく上昇する。その結果、穴広げ時の亀裂進展が促進され、穴広げ率が低下して、伸びフランジ性が低下する。また、残留オーステナイトが応力誘起変態することによってＹＲが低下し、部品強度が低下する。このため、Ｓｉ含有量は、２．５０％以下であり、２．００％以下が好ましく、１．５０％以下がより好ましい。 <Si: 0.01% or more and 2.50% or less>
Silicon is a component that increases the strength of a steel sheet by suppressing the precipitation of cementite in martensite and by solid solution strengthening. To obtain this effect, the silicon content is 0.01% or more, preferably 0.05% or more, and more preferably 0.10% or more.
On the other hand, if the Si content is too high, the carbide precipitation during bainite transformation is significantly suppressed, the residual austenite increases excessively, and the hardness of martensite generated from the residual austenite during punching increases significantly. As a result, crack growth during hole expansion is promoted, the hole expansion ratio decreases, and the stretch flangeability decreases. In addition, the residual austenite undergoes stress-induced transformation, which reduces the YR and reduces the part strength. For this reason, the Si content is 2.50% or less, preferably 2.00% or less, and more preferably 1.50% or less.
《Ｍｎ：０．１０％以上５．００％以下》
　Ｍｎは、鋼の重要な基本成分の１つであり、特に本発明においては、マルテンサイトの面積率に影響する。
　Ｍｎ含有量が少なすぎると、マルテンサイトの面積率が減少し、７８０ＭＰａ以上のＴＳを実現することが困難になる。このため、Ｍｎ含有量は、０．１０％以上であり、０．９０％以上が好ましく、１．８０％以上がより好ましい。
　一方、Ｍｎ含有量が多すぎると、オーステナイトが安定化し、残留オーステナイトが過度に増加し、打ち抜き時に残留オーステナイトから生成するマルテンサイトの硬度が大きく上昇する。その結果、穴広げ時の亀裂進展が促進され、穴広げ率が低下して、伸びフランジ性が低下する。また、残留オーステナイトが応力誘起変態することによってＹＲが低下し、部品強度が低下する。このため、Ｍｎ含有量は、５．００％以下であり、４．２０％以下が好ましく、３．６０％以下がより好ましい。 <<Mn: 0.10% or more and 5.00% or less>>
Mn is one of the important basic components of steel, and particularly in the present invention, it affects the area ratio of martensite.
If the Mn content is too low, the area ratio of martensite decreases, making it difficult to achieve a TS of 780 MPa or more. Therefore, the Mn content is 0.10% or more, preferably 0.90% or more, and more preferably 1.80% or more.
On the other hand, if the Mn content is too high, the austenite is stabilized, the residual austenite is excessively increased, and the hardness of the martensite generated from the residual austenite during punching is greatly increased. As a result, crack propagation during hole expansion is promoted, the hole expansion ratio is reduced, and the stretch flangeability is reduced. In addition, the residual austenite is transformed by stress induced transformation, which reduces the YR and reduces the strength of the part. For this reason, the Mn content is 5.00% or less, preferably 4.20% or less, and more preferably 3.60% or less.
《Ｐ：０．１００％以下》
　Ｐは、旧オーステナイト粒界に偏析して粒界を脆化させるため、鋼板の極限変形能を低下させることから、λを低下させ、曲げ性を低下させ得る成分である。このため、Ｐの含有量は０．１００％以下であり、好ましくは０．０７０％以下である。
　Ｐの含有量の下限は特に限定されないが、Ｐは固溶強化元素であり、鋼板の強度を上昇させることができることから、０．００１％以上とすることが好ましい。 《P: 0.100% or less》
P is a component that segregates at prior austenite grain boundaries to embrittle the grain boundaries, thereby reducing the ultimate deformability of the steel sheet, thereby reducing λ and reducing bendability. For this reason, the P content is 0.100% or less, and preferably 0.070% or less.
The lower limit of the P content is not particularly limited, but since P is a solid solution strengthening element and can increase the strength of the steel sheet, it is preferable to set the lower limit to 0.001% or more.
《Ｓ：０．０２００％以下》
　Ｓは、硫化物として存在し、鋼板の極限変形能を低下させることから、λを低下させ、曲げ性を低下させ得る成分である。このため、Ｓの含有量は０．０２００％以下であり、好ましくは０．００５０％以下である。
　Ｓの含有量の下限は特に限定されないが、生産技術上の制約から、０．０００１％以上とすることが好ましい。 《S: 0.0200% or less》
S exists as a sulfide and is a component that reduces the ultimate deformability of the steel sheet, thereby reducing λ and reducing bendability. Therefore, the S content is 0.0200% or less, preferably 0.0050% or less.
The lower limit of the S content is not particularly limited, but due to constraints on production technology, it is preferably set to 0.0001% or more.
《Ａｌ：１．０００％以下》
　Ａｌは、十分な脱酸を行ない、鋼中介在物を低減させるのに有効な成分であるが、Ａｌ含有量が多すぎると、フェライトが多量に生成し、穴広げ率が低下して、伸びフランジ性を低下させ得る。このため、Ａｌ含有量は、１．０００％以下であり、０．５００％以下が好ましく、０．１００％以下がより好ましい。
　一方、安定して脱酸を行なうためには、Ａｌ含有量は、０．０１０％以上が好ましく、０．０１５％以上がより好ましく、０．０２０％以上がさらに好ましい。 <<Al: 1.000% or less>>
Al is an effective component for sufficient deoxidation and reducing inclusions in steel, but if the Al content is too high, a large amount of ferrite is generated, the hole expansion ratio decreases, and the stretch flangeability may decrease. Therefore, the Al content is 1.000% or less, preferably 0.500% or less, and more preferably 0.100% or less.
On the other hand, in order to perform stable deoxidation, the Al content is preferably 0.010% or more, more preferably 0.015% or more, and even more preferably 0.020% or more.
《Ｎ：０．０１００％以下》
　Ｎは、窒化物として存在し、鋼板の極限変形能を低下させることから、λを低下させ、曲げ性を低下させ得る成分である。このため、Ｎの含有量は０．０１００％以下であり、好ましくは０．００５０％以下である。
　Ｎの含有量の下限は特に限定されないが、生産技術上の制約から、Ｎの含有量は０．０００１％以上とすることが好ましい。 <N: 0.0100% or less>
N exists as a nitride and is a component that reduces the ultimate deformability of the steel sheet, thereby reducing λ and reducing bendability. Therefore, the N content is 0.0100% or less, preferably 0.0050% or less.
Although there is no particular lower limit for the N content, due to constraints on production technology, the N content is preferably 0.0001% or more.
《Ｏ：０．０１００％以下》
　Ｏは、酸化物として存在し、鋼板の極限変形能を低下させることから、λを低下させ、曲げ性を低下させ得る成分である。このため、Ｏの含有量は０．０１００％以下であり、好ましくは０．００５０％以下である。
　Ｏの含有量の下限は特に限定されないが、生産技術上の制約から、Ｏの含有量は０．０００１％以上とすることが好ましい。 <O: 0.0100% or less>
O exists as an oxide and is a component that reduces the ultimate deformability of the steel sheet, thereby reducing λ and reducing bendability. Therefore, the O content is 0.0100% or less, and preferably 0.0050% or less.
Although there is no particular lower limit for the O content, due to constraints on production technology, the O content is preferably 0.0001% or more.
《任意成分》
　本発明の高強度鋼板は、上記の成分組成に加えて、さらに、質量％で、
Ｔｉ：０．２００％以下、Ｎｂ：０．２００％以下、Ｖ：０．２００％以下、
Ｔａ：０．１０％以下、Ｗ：０．１０％以下、
Ｂ：０．０１００％以下、
Ｃｒ：１．００％以下、Ｍｏ：１．００％以下、Ｎｉ：１．００％以下、
Ｃｏ：０．０１０％以下、
Ｃｕ：１．００％以下、
Ｓｎ：０．２００％以下、
Ｓｂ：０．２００％以下、
Ｃａ：０．０１００％以下、Ｍｇ：０．０１００％以下、ＲＥＭ：０．０１００％以下、
Ｚｒ：０．０２０％以下、Ｔｅ：０．０２０％以下、
Ｈｆ：０．１０％以下及び
Ｂｉ：０．２００％以下
からなる群より選ばれる少なくとも１種の元素を含有していてもよい。これらの元素は、単独でも、２種以上の組み合わせてでもよい。 Optional Ingredients
The high strength steel plate of the present invention further comprises, in addition to the above-mentioned composition, in mass%:
Ti: 0.200% or less, Nb: 0.200% or less, V: 0.200% or less,
Ta: 0.10% or less, W: 0.10% or less,
B: 0.0100% or less,
Cr: 1.00% or less, Mo: 1.00% or less, Ni: 1.00% or less,
Co: 0.010% or less,
Cu: 1.00% 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: 0.020% or less, Te: 0.020% or less,
At least one element selected from the group consisting of Hf: 0.10% or less and Bi: 0.200% or less may be contained. These elements may be contained alone or in combination of two or more kinds.
　Ｔｉ、Ｎｂ又はＶを含有する場合、粗大な析出物や介在物が多量に生成し、鋼板の極限変形能を低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｔｉ、Ｎｂ又はＶの含有量はそれぞれ０．２００％以下にすることが好ましく、より好ましくは０．１００％以下である。Ｔｉ、Ｎｂ又はＶの含有量の下限は特に限定されないが、熱間圧延時あるいは連続焼鈍時に、微細な炭化物、窒化物もしくは炭窒化物を形成することによって、鋼板の強度を上昇させることから、Ｔｉ、Ｎｂ又はＶの含有量はそれぞれ０．００１％以上とすることが好ましい。 When Ti, Nb or V is contained, a large amount of coarse precipitates or inclusions are generated, which reduces the ultimate deformability of the steel sheet, and thus reduces λ and bendability. To avoid this, the Ti, Nb or V content is preferably 0.200% or less, and more preferably 0.100% or less. There is no particular lower limit for the Ti, Nb or V content, but since they form fine carbides, nitrides or carbonitrides during hot rolling or continuous annealing, thereby increasing the strength of the steel sheet, it is preferable that the Ti, Nb or V content is 0.001% or more.
　Ｔａ又はＷを含有する場合、粗大な析出物や介在物が多量に生成し、鋼板の極限変形能が低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｔａ又はＷの含有量はそれぞれ０．１０％以下にすることが好ましく、より好ましくは０．０８％以下である。Ｔａ又はＷの含有量の下限は特に限定されないが、熱間圧延時あるいは連続焼鈍時に、微細な炭化物、窒化物もしくは炭窒化物を形成することによって、鋼板の強度を上昇させることから、Ｔａ又はＷの含有量はそれぞれ０．０１％以上とすることが好ましい。 When Ta or W is contained, a large amount of coarse precipitates or inclusions are generated, which reduces the ultimate deformability of the steel sheet, and thus reduces λ and bendability. To avoid this, the Ta or W content is preferably 0.10% or less, and more preferably 0.08% or less. There is no particular lower limit for the Ta or W content, but since they increase the strength of the steel sheet by forming fine carbides, nitrides, or carbonitrides during hot rolling or continuous annealing, the Ta or W content is preferably 0.01% or more.
　Ｂを含有する場合、鋳造時あるいは熱間圧延時において鋼板内部に割れが生成し、鋼板の極限変形能が低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｂの含有量は０．０１００％以下にすることが好ましく、より好ましくは０．０００３％以上である。Ｂの含有量の下限は特に限定されないが、焼鈍中にオーステナイト粒界に偏析し、焼入れ性を向上させる元素であることから、Ｂの含有量は０．０００３％以上とすることが好ましい。 When B is contained, cracks may form inside the steel sheet during casting or hot rolling, reducing the ultimate deformability of the steel sheet, which in turn reduces λ and bendability. To avoid this, the B content is preferably 0.0100% or less, and more preferably 0.0003% or more. There is no particular lower limit for the B content, but since B is an element that segregates to austenite grain boundaries during annealing and improves hardenability, the B content is preferably 0.0003% or more.
　Ｃｒ、Ｍｏ又はＮｉを含有する場合、粗大な析出物や介在物が増加し、鋼板の極限変形能が低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｃｒ、Ｍｏ又はＮｉの含有量はそれぞれ１．００％以下にすることが好ましく、より好ましくは０．８０％以下である。Ｃｒ、Ｍｏ又はＮｉの含有量の下限は特に限定されないが、焼入れ性を向上させる元素であることから、Ｃｒ、Ｍｏ又はＮｉの含有量はそれぞれ０．０１％以上とすることが好ましい。 When Cr, Mo or Ni is contained, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet, and thus reducing λ and bendability. To avoid this, it is preferable that the Cr, Mo or Ni content is 1.00% or less, and more preferably 0.80% or less. There is no particular lower limit for the Cr, Mo or Ni content, but since these elements improve hardenability, it is preferable that the Cr, Mo or Ni content is 0.01% or more.
　Ｃｏを含有する場合、粗大な析出物や介在物が増加し、鋼板の極限変形能が低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｃｏの含有量は０．０１０％以下にすることが好ましく、より好ましくは０．００８％以下である。Ｃｏの含有量の下限は特に限定されないが、焼入れ性を向上させる元素であることから、Ｃｏの含有量は０．００１％以上とすることが好ましい。 When Co is contained, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet, which in turn reduces λ and bendability. To avoid this, the Co content is preferably 0.010% or less, and more preferably 0.008% or less. There is no particular lower limit for the Co content, but since Co is an element that improves hardenability, it is preferable that the Co content be 0.001% or more.
　Ｃｕを含有する場合、粗大な析出物や介在物が増加し、鋼板の極限変形能が低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｃｕの含有量は１．００％以下にすることが好ましく、より好ましくは０．８０％以下である。Ｃｕの含有量の下限は特に限定されないが、焼入れ性を向上させる元素であることから、Ｃｕの含有量は０．０１％以上とすることが好ましい。 When Cu is contained, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet, and thus reducing λ and bendability. To avoid this, the Cu content is preferably 1.00% or less, and more preferably 0.80% or less. There is no particular lower limit for the Cu content, but since Cu is an element that improves hardenability, it is preferable that the Cu content be 0.01% or more.
　Ｓｎを含有する場合、鋳造時あるいは熱間圧延時において鋼板内部に割れが生成し、鋼板の極限変形能が低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｓｎの含有量は０．２００％以下にすることが好ましく、より好ましくは０．１００％以下である。Ｓｎの含有量の下限は特に限定されないが、Ｓｎは焼入れ性を向上させる元素であることから、Ｓｎの含有量は０．００１％以上とすることが好ましい。 When Sn is contained, cracks may form inside the steel sheet during casting or hot rolling, reducing the ultimate deformability of the steel sheet, which in turn reduces λ and bendability. To avoid this, the Sn content is preferably 0.200% or less, and more preferably 0.100% or less. There is no particular lower limit for the Sn content, but since Sn is an element that improves hardenability, the Sn content is preferably 0.001% or more.
　Ｓｂを含有する場合、粗大な析出物や介在物が増加し、鋼板の極限変形能が低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｓｂの含有量は０．２００％以下にすることが好ましく、より好ましくは０．１００％以下である。Ｓｂの含有量の下限は特に限定されないが、表層軟化層の厚さを制御し、強度調整を可能にする元素であることから、Ｓｂの含有量は０．００１％以上とすることが好ましい。 When Sb is contained, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet, which in turn reduces λ and bendability. To avoid this, the Sb content is preferably 0.200% or less, and more preferably 0.100% or less. There is no particular lower limit for the Sb content, but since Sb is an element that controls the thickness of the surface softened layer and enables strength adjustment, the Sb content is preferably 0.001% or more.
　Ｃａ、Ｍｇ又はＲＥＭを含有する場合、粗大な析出物や介在物が増加し、鋼板の極限変形能が低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｃａ、Ｍｇ又はＲＥＭの含有量はそれぞれ０．０１００％以下にすることが好ましく、より好ましくは０．００５０％以下である。Ｃａ、Ｍｇ又はＲＥＭの含有量の下限は特に限定されないが、窒化物や硫化物の形状を球状化し、鋼板の極限変形能を向上する元素であることから、Ｃａ、Ｍｇ又はＲＥＭの含有量はそれぞれ０．０００１％以上とすることが好ましい。 When Ca, Mg or REM is contained, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet, which in turn reduces λ and bendability. To avoid this, it is preferable that the Ca, Mg or REM content is 0.0100% or less, and more preferably 0.0050% or less. There is no particular lower limit for the Ca, Mg or REM content, but since these elements spheroidize the shape of nitrides and sulfides and improve the ultimate deformability of the steel sheet, it is preferable that the Ca, Mg or REM content is 0.0001% or more.
　Ｚｒ又はＴｅを含有する場合、粗大な析出物や介在物が増加し、鋼板の極限変形能が低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｚｒ又はＴｅの含有量はそれぞれ０．１００％以下にすることが好ましく、より好ましくは０．０８０％以下である。Ｚｒ又はＴｅの含有量の下限は特に限定されないが、窒化物や硫化物の形状を球状化し、鋼板の極限変形能を向上する元素であることから、Ｚｒ又はＴｅの含有量はそれぞれ０．００１％以上とすることが好ましい。 When Zr or Te is contained, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet, which in turn reduces λ and bendability. To avoid this, it is preferable that the Zr or Te content is 0.100% or less, and more preferably 0.080% or less. There is no particular lower limit for the Zr or Te content, but since these elements spheroidize the shape of nitrides and sulfides and improve the ultimate deformability of the steel sheet, it is preferable that the Zr or Te content is 0.001% or more.
　Ｈｆを含有する場合、粗大な析出物や介在物が増加し、鋼板の極限変形能が低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｈｆの含有量は０．１０％以下にすることが好ましく、より好ましくは０．０８％以下である。Ｈｆの含有量の下限は特に限定されないが、窒化物や硫化物の形状を球状化し、鋼板の極限変形能を向上する元素であることから、Ｈｆの含有量は０．０１％以上とすることが好ましい。 When Hf is contained, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet, which in turn reduces λ and bendability. To avoid this, the Hf content is preferably 0.10% or less, and more preferably 0.08% or less. There is no particular lower limit for the Hf content, but since Hf is an element that spheroidizes the shape of nitrides and sulfides and improves the ultimate deformability of the steel sheet, the Hf content is preferably 0.01% or more.
　Ｂｉを含有する場合、粗大な析出物や介在物が増加し、鋼板の極限変形能が低下し、ひいてはλが低下し、曲げ性が低下することを回避するため、Ｂｉの含有量は０．２００％以下にすることが好ましく、より好ましくは０．１００％以下である。Ｂｉの含有量の下限は特に限定されないが、偏析を軽減する元素であることから、Ｂｉの含有量は０．００１％以上とすることが好ましい。 When Bi is contained, the amount of coarse precipitates and inclusions increases, reducing the ultimate deformability of the steel sheet, which in turn reduces λ and bendability. To avoid this, the Bi content is preferably 0.200% or less, and more preferably 0.100% or less. There is no particular lower limit for the Bi content, but since Bi is an element that reduces segregation, it is preferable that the Bi content be 0.001% or more.
　本発明の一実施形態に従う高強度鋼板は、上記の必須成分及び場合により任意成分を含有し、残部がＦｅ及び不可避的不純物からなる成分組成を有する。ここで、不可避的不純物としては、Ｚｎ、Ｐｂ、Ａｓ、Ｇｅ、Ｓｒ及びＣｓが挙げられる。これらの不可避的不純物は、合計で０．１００％以下の量で含有されることが許容される。 The high-strength steel plate according to one embodiment of the present invention has a composition containing the above essential components and optional components, with the balance being Fe and unavoidable impurities. Here, the unavoidable impurities include Zn, Pb, As, Ge, Sr, and Cs. These unavoidable impurities are permitted to be contained in an amount of 0.100% or less in total.
〈鋼組織〉
　本発明の高強度鋼板の鋼組織について説明する。 <Steel structure>
The steel structure of the high strength steel plate of the present invention will be described.
《マルテンサイトの面積率：１０％以上８０％以下》
　マルテンサイトを含有することにより、７８０ＭＰａ以上のＴＳを容易に実現できる。このため、マルテンサイトの面積率は、１０％以上であり、１５％以上が好ましく、２０％以上がより好ましい。
　一方、マルテンサイトが多すぎると、Ｅｌが低下し延性が劣化する。このため、マルテンサイトの面積率は、８０％以下であり、７５％以下が好ましく、７０％以下がより好ましい。 <Area ratio of martensite: 10% to 80%>
By containing martensite, a TS of 780 MPa or more can be easily achieved. Therefore, the area ratio of martensite is 10% or more, preferably 15% or more, and more preferably 20% or more.
On the other hand, if the martensite content is too high, El decreases and ductility deteriorates, so the area ratio of martensite is 80% or less, preferably 75% or less, and more preferably 70% or less.
　ここで、マルテンサイトは、下部ベイナイト、後述する焼鈍において実施する冷却中に自己焼戻しを生じたマルテンサイト、後述する第２加熱工程により焼戻されたマルテンサイト等を含む。
　マルテンサイトの観察位置は、後述するように、鋼板の板厚の１／４位置である。 Here, martensite includes lower bainite, martensite that has undergone self-tempering during cooling performed in the annealing step described later, and martensite that has been tempered in the second heating step described later.
As described later, the martensite was observed at a position corresponding to 1/4 of the thickness of the steel plate.
《ベイナイトの面積率：２％以上７０％以下》
　ベイナイトを含有することにより、組織間の硬度差が軽減されλが上昇する。また、界面でのき裂進展を抑制することから低温靭性が向上する。このため、ベイナイトの面積率は、２％以上であり、３％以上が好ましく、４％以上がより好ましい。
　一方、ベイナイトが多すぎると、十分な量のマルテンサイト量を確保できず、ＴＳが低下する。このため、ベイナイトの面積率は、７０％以下であり、６０％以下が好ましく、５０％以下がより好ましい。 <Area ratio of bainite: 2% to 70%>
By including bainite, the hardness difference between the structures is reduced and λ is increased. In addition, the crack propagation at the interface is suppressed, and thus low-temperature toughness is improved. For this reason, the area fraction of bainite is 2% or more, preferably 3% or more, and more preferably 4% or more.
On the other hand, if the amount of bainite is too large, a sufficient amount of martensite cannot be secured, and TS decreases. Therefore, the area ratio of bainite is 70% or less, preferably 60% or less, and more preferably 50% or less.
　ここで、ベイナイトは、Ｍｓ以上７００℃以下の温度域で生成する角状に生成したベイニティックフェライト、鉄系炭化物、残留オーステナイトの混合組織である。
　ベイナイトの観察位置は、後述するように、鋼板の板厚の１／４位置である。 Here, bainite is a mixed structure of angular bainitic ferrite, iron-based carbides, and retained austenite that is formed in a temperature range of Ms or higher and 700° C. or lower.
The observation position of bainite is a quarter position of the sheet thickness of the steel sheet, as described later.
《フェライトの面積率：８０％以下》
　フェライトの面積率を８０％以下とすることで、所望の強度を容易に得ることができる。フェライトの面積率は０％であっても本発明の効果は得られる。一方、フェライトが多すぎると、十分な量のマルテンサイト量を確保できずに、所望のＴＳが得られない。このため、フェライトの面積率は、８０％以下であり、７５％以下が好ましく、７０％以下がより好ましい。一方、Ｅｌを上昇させ、延性をさらに向上するためには、フェライトの面積率は、１０％以上が好ましく、１５％以上がより好ましい。 Ferrite area ratio: 80% or less
By setting the area ratio of ferrite to 80% or less, the desired strength can be easily obtained. The effect of the present invention can be obtained even if the area ratio of ferrite is 0%. On the other hand, if the amount of ferrite is too much, a sufficient amount of martensite cannot be secured, and the desired TS cannot be obtained. Therefore, the area ratio of ferrite is 80% or less, preferably 75% or less, and more preferably 70% or less. On the other hand, in order to increase El and further improve ductility, the area ratio of ferrite is preferably 10% or more, and more preferably 15% or more.
　ここで、フェライトは、比較的高温で生成している軟質なＢＣＣ鉄であり、アロトリオモルフフェライト、イディオモルフフェライトを含む。
　フェライトの観察位置は、後述するように、鋼板の板厚の１／４位置である。 Here, ferrite is soft BCC iron formed at relatively high temperatures, and includes allotriomorph ferrite and idiomorph ferrite.
The observation position of ferrite was a quarter position of the sheet thickness of the steel sheet, as described later.
　マルテンサイト、ベイナイト、フェライトの面積率の測定方法は、以下のとおりである。
　まず、鋼板から、その圧延方向に平行な板厚断面（板厚１／４位置のＬ断面）が観察面となるように、サンプルを切り出す。サンプルの観察面を、ダイヤモンドペーストを用いて鏡面研磨し、その後、コロイダルシリカを用いて仕上げ研磨を施し、さらに１体積％ナイタールを用いてエッチングすることにより、組織を現出させる。
　次いで、サンプルの観察面を、加速電圧１０ｋＶの条件で、走査型電子顕微鏡（ＳＥＭ）を用いて３０００倍の倍率で観察し、３視野（１視野４０μｍ×３０μｍ）分のＳＥＭ画像を得る。
　得られたＳＥＭ画像から、Ａｄｏｂｅ Ｐｈｏｔｏｓｈｏｐ（Ａｄｏｂｅ Ｓｙｓｔｅｍｓ社製）を用いて、各組織の面積率を算出する。具体的には、各組織の面積を測定面積で除して得られる値を、各組織の面積率とする。各組織の面積率を３視野分算出し、それらの平均値を各組織の面積率とする。 The method for measuring the area ratios of martensite, bainite, and ferrite is as follows.
First, a sample is cut out from a steel sheet so that a thickness cross section (L cross section at 1/4 of the sheet thickness) parallel to the rolling direction becomes an observation surface. The observation surface of the sample is mirror-polished with diamond paste, then finish-polished with colloidal silica, and further etched with 1 volume % nital to reveal the structure.
Next, the observation surface of the sample is observed at a magnification of 3000 times using a scanning electron microscope (SEM) under the condition of an acceleration voltage of 10 kV, and SEM images of three visual fields (one visual field is 40 μm×30 μm) are obtained.
From the obtained SEM image, the area ratio of each structure is calculated using Adobe Photoshop (manufactured by Adobe Systems). Specifically, the value obtained by dividing the area of each structure by the measured area is regarded as the area ratio of each structure. The area ratio of each structure is calculated for three fields of view, and the average value thereof is regarded as the area ratio of each structure.
　ＳＥＭ画像において、フェライトは灰色を呈し、白いコントラストを呈する炭化物を内包しない平坦な組織領域である。
　ベイナイトは灰色を呈する角状のベイニティックフェライトと白いコントラストを呈する鉄系炭化物、針状の形状の残留オーステナイトから構成される混合組織領域である。
　マルテンサイトは内部に微細な凹凸を有する階層構造を持つ組織である。
　これらは互いに識別できる。 In SEM images, ferrite appears as a grey, flat region of the structure that does not contain carbides, which provides a white contrast.
Bainite is a mixed structure region consisting of gray angular bainitic ferrite, white contrasting iron carbides, and needle-shaped retained austenite.
Martensite has a hierarchical structure with minute internal irregularities.
These can be distinguished from one another.
《残留オーステナイトの面積率：１５％以下》
　残留オーステナイトを少なくすることにより、良好な部品強度及び伸びフランジ性が得られる。このため、残留オーステナイトの面積率は、１５％以下であり、１０％以下が好ましい。下限は特に限定されず、残留オーステナイトの面積率が０％であっても、この効果は得られる。 <Area ratio of retained austenite: 15% or less>
By reducing the amount of retained austenite, good part strength and stretch flangeability can be obtained. Therefore, the area ratio of the retained austenite is 15% or less, and preferably 10% or less. There is no particular lower limit, and this effect can be obtained even if the area ratio of the retained austenite is 0%.
　残留オーステナイトの面積率の測定方法は、以下のとおりである。
　まず、鋼板を、その板厚１／４位置（鋼板表面から深さ方向で板厚の１／４に相当する位置）が測定面となるように研削し、その後、化学研磨によりさらに０．１ｍｍ研磨し、サンプルを得る。
　サンプルの測定面について、Ｘ線回折装置により、ＣｏのＫα線源を用いて、ｆｃｃ鉄（オーステナイト）の（２００）面、（２２０）面及び（３１１）面、ならびに、ｂｃｃ鉄の（２００）面、（２１１）面及び（２２０）面の積分反射強度を測定する。
　ｂｃｃ鉄の各面の積分反射強度に対するｆｃｃ鉄の各面の積分反射強度の強度比を求める。９つの強度比の平均値を、残留オーステナイトの体積率とする。。この残留オーステナイトの体積率を３次元的に均一であるとみなして、鋼板の板厚１／４位置での残留オーステナイトの面積率とする。 The method for measuring the area ratio of retained austenite is as follows.
First, the steel plate is ground so that the 1/4 position of the plate thickness (the position corresponding to 1/4 of the plate thickness in the depth direction from the surface of the steel plate) becomes the measurement surface, and then the plate is further polished by 0.1 mm by chemical polishing to obtain a sample.
For the measurement surface of the sample, an X-ray diffractometer is used to measure 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, using a Co Kα radiation source.
The intensity ratio of the integrated reflection intensity of each surface of the fcc iron to the integrated reflection intensity of each surface of the bcc iron is calculated. The average value of the nine intensity ratios is taken as the volume fraction of the retained austenite. This volume fraction of the retained austenite is considered to be three-dimensionally uniform, and is taken as the area fraction of the retained austenite at the 1/4 position of the plate thickness of the steel plate.
《残部組織》
　本発明の鋼組織は、上述したマルテンサイト、ベイナイト、フェライト及び残留オーステナイト以外の組織（残部組織）を有していてもよい。
　残部組織としては、マルテンサイト、ベイナイト、フェライト及び残留オーステナイト以外の組織であって、鋼板の組織として公知の組織が挙げられ、例えば、パーライト、フェライト中に析出した合金炭窒化物等が挙げられる。
　なお、ベイナイト中に存在する鉄系炭化物やマルテンサイト中に析出した準安定炭化物及びマルテンサイト中に析出したセメンタイト等の鉄系炭化物は、残部組織に含まれない。
　残部組織の面積率は、本発明の効果が損なわれないという理由から、３％以下が好ましい。 {Remaining structure}
The steel structure of the present invention may have a structure (remaining structure) other than the above-mentioned martensite, bainite, ferrite, and retained austenite.
The remaining structure is a structure other than martensite, bainite, ferrite, and retained austenite, and may be any structure known as a steel sheet structure, such as pearlite and alloy carbonitrides precipitated in ferrite.
It should be noted that iron-based carbides present in bainite, metastable carbides precipitated in martensite, and iron-based carbides such as cementite precipitated in martensite are not included in the remaining structure.
The area ratio of the remaining structure is preferably 3% or less so as not to impair the effects of the present invention.
《マルテンサイトブロック数に対する準安定炭化物が存在するマルテンサイトブロック数の割合：２％以上》
　マルテンサイトブロック中に析出した準安定炭化物によって、優れた部品強度、延性、剪断端面曲げ性、伸びフランジ性を保ちつつ、低温靭性が向上する。これはマルテンサイトブロック中に析出した準安定炭化物が低温でのき裂の発生、進展を抑制するためであると考えられる。
　この効果を得るためには、マルテンサイトブロック数に対する準安定炭化物が存在するマルテンサイトブロック数の割合（以下、「割合ｐ」ともいう）は、２％以上であり、５％以上が好ましく、１０％以上がより好ましく、２０％以上がさらに好ましい。３０％以上が特に好ましい。割合ｐの上限は、特に限定されず、１００％であってもよい。 <The ratio of the number of martensite blocks containing metastable carbides to the number of martensite blocks: 2% or more>
The metastable carbides precipitated in the martensite block improve low-temperature toughness while maintaining excellent part strength, ductility, shear edge bendability, and stretch flangeability. This is believed to be because the metastable carbides precipitated in the martensite block suppress the initiation and propagation of cracks at low temperatures.
In order to obtain this effect, the ratio of the number of martensite blocks in which metastable carbides exist to the number of martensite blocks (hereinafter also referred to as "ratio p") is 2% or more, preferably 5% or more, more preferably 10% or more, even more preferably 20% or more, and particularly preferably 30% or more. The upper limit of ratio p is not particularly limited and may be 100%.
　ここで、準安定炭化物は、マルテンサイトの焼戻し過程で析出する準安定な炭化物である。準安定炭化物は、例えば、セメンタイト以外のＦｅ炭化物（鉄系炭化物）であり、イプシロン（ε）炭化物、イータ（η）炭化物及びカイ（χ）炭化物からなる群より選ばれる少なくとも１種の炭化物が挙げられる。 Here, metastable carbides are metastable carbides that precipitate during the tempering process of martensite. Metastable carbides are, for example, Fe carbides (iron-based carbides) other than cementite, and include at least one type of carbide selected from the group consisting of epsilon (ε) carbides, eta (η) carbides, and chi (χ) carbides.
　マルテンサイトブロック数に対する準安定炭化物が存在するマルテンサイトブロック数の割合（割合ｐ）の測定方法は、以下のとおりである。
　まず、鋼板を、その板厚１／４位置（鋼板表面から深さ方向で板厚の１／４に相当する位置）が観察面となるように研削し、その後、電解研磨して、サンプルを作製する。作製したサンプルの観察面を、透過型電子顕微鏡（ＴＥＭ）を用いて、加速電圧２００ｋＶの条件で、観察する。
　マルテンサイトブロックの［１００］方位から電子線を入射すると、母相マルテンサイトの電子回折図形が得られる。隣接するマルテンサイトブロックどうしは、ブロック境界を介して結晶方位が異なるので、明視野像でコントラストが異なることから、互いに区別される。マルテンサイトとフェライト及びベイナイトどうしは、マルテンサイト中には高密度の転位が観察され、フェライト及びベイナイト中は転位密度が比較的低いことから互いに区別される。 The method for measuring the ratio (ratio p) of the number of martensite blocks containing metastable carbides to the number of martensite blocks is as follows.
First, a steel sheet is ground so that a 1/4 position of the sheet thickness (a position corresponding to 1/4 of the sheet thickness in the depth direction from the surface of the steel sheet) becomes an observation surface, and then electrolytic polishing is performed to prepare a sample. The observation surface of the prepared sample is observed using a transmission electron microscope (TEM) at an acceleration voltage of 200 kV.
When an electron beam is incident on the martensite block from the [100] direction, an electron diffraction pattern of the parent martensite is obtained. Adjacent martensite blocks have different crystal orientations across the block boundaries, and therefore can be distinguished from each other by the different contrast in the bright-field image. Martensite can be distinguished from ferrite and bainite by the high density of dislocations observed in martensite and the relatively low dislocation density in ferrite and bainite.
　図１は、炭化物が存在するマルテンサイトの電子回折図形の一例である。
　観察した単一のマルテンサイトブロックに炭化物が存在する場合、図１に示すように、母相マルテンサイト（α）の電子回折図形に加えて、炭化物の電子回折図形が得られる。
　図１中、黒丸は、電子線が［１００］方位から入射した場合における、母相マルテンサイトの電子回折斑点、白丸は、炭化物の電子回折斑点を示す。
　図１における母相マルテンサイトの電子回折斑点間の距離Ｄ１と、炭化物の電子回折斑点間の距離Ｄ２とから、炭化物の面間隔ｄｃと母相マルテンサイトの面間隔ｄｍとの比ｄｃ／ｄｍを、下記式３を用いて計算する。
式３：ｄｃ／ｄｍ＝Ｄ１／Ｄ２ FIG. 1 is an example of an electron diffraction pattern of martensite in which carbides are present.
When carbides are present in a single martensite block observed, the electron diffraction pattern of the carbides is obtained in addition to the electron diffraction pattern of the parent martensite (α), as shown in FIG.
In FIG. 1, black circles indicate electron diffraction spots of the martensite parent phase when the electron beam is incident from the [100] direction, and white circles indicate electron diffraction spots of carbides.
From the distance D1 between the electron diffraction spots of the martensite parent phase and the distance D2 between the electron diffraction spots of the carbide in FIG. 1, the ratio dc/dm of the lattice spacing dc of the carbide to the lattice spacing dm of the martensite parent phase is calculated using the following formula 3.
Equation 3: dc/dm=D1/D2
　観察したマルテンサイトブロックについて、炭化物の面間隔ｄｃと母相マルテンサイトの面間隔ｄｍとの比ｄｃ／ｄｍが１．０２０以上１．１５０以下である場合、そのマルテンサイトブロックは、準安定炭化物が存在するマルテンサイトブロックであると定義する。すなわち、「マルテンサイトブロック数に対する準安定炭化物が存在するマルテンサイトブロック数の割合」は、「マルテンサイトブロック数に対する比ｄｃ／ｄｍが１．０２０以上１．１５０以下であるマルテンサイトブロック数の割合」といい換えることができる。 If the ratio dc/dm of the interplanar spacing dc of carbides to the interplanar spacing dm of parent martensite is 1.020 or more and 1.150 or less for the observed martensite block, the martensite block is defined as a martensite block containing metastable carbides. In other words, "the ratio of the number of martensite blocks containing metastable carbides to the number of martensite blocks" can be rephrased as "the ratio of the number of martensite blocks with a ratio dc/dm of 1.020 or more and 1.150 or less to the number of martensite blocks."
　母相マルテンサイトの電子回折斑点間の距離Ｄ１は不変の値であるが、炭化物の電子回折斑点間の距離Ｄ２は、その炭化物によって値が変化する。
　例えば、炭化物がセメンタイトである場合、距離Ｄ２は距離Ｄ１と等しい。このため、比ｄｃ／ｄｍの値は、１である（Ｄ１／Ｄ２＝ｄｃ／ｄｍ＝１）。
　一方、準安定炭化物（ε炭化物など）の距離Ｄ２は、セメンタイトよりも短いので、比ｄｃ／ｄｍ（＝Ｄ１／Ｄ２）値は１よりも大きい。このため、準安定炭化物が存在する場合における比ｄｃ／ｄｍの下限値を「１．０２０」とする。
　なお、比ｄｃ／ｄｍの上限については、ε炭化物の距離Ｄ２に基づいて、「１．１５０」と定める。 The distance D1 between electron diffraction spots of the martensite parent phase is a constant value, but the distance D2 between electron diffraction spots of carbides varies depending on the carbide.
For example, if the carbide is cementite, then the distance D2 is equal to the distance D1, so the value of the ratio dc/dm is 1 (D1/D2=dc/dm=1).
On the other hand, since the distance D2 of metastable carbides (such as ε carbide) is shorter than that of cementite, the ratio dc/dm (=D1/D2) is greater than 1. For this reason, the lower limit of the ratio dc/dm in the presence of metastable carbides is set to "1.020".
The upper limit of the ratio dc/dm is set to "1.150" based on the distance D2 of the ε carbide.
　準安定炭化物は、マルテンサイトブロックの内部に存在してもよいし、ブロック境界などの境界部分に存在してもよいが、マルテンサイトブロックの内部に存在することが好ましい。 Metastable carbides may be present inside the martensite blocks or at boundary portions such as block boundaries, but are preferably present inside the martensite blocks.
　５０個のマルテンサイトブロックを観察し、準安定炭化物が存在するマルテンサイトブロック数を、観察したマルテンサイトブロック数で割った値（準安定炭化物が存在するマルテンサイトブロック数／５０）を求める。求めた値に１００を乗じることで、マルテンサイトブロック数に対する準安定炭化物が存在するマルテンサイトブロック数の割合（割合ｐ（％））とする。 50 martensite blocks are observed, and the number of martensite blocks containing metastable carbides is divided by the number of martensite blocks observed (number of martensite blocks containing metastable carbides/50). The value thus obtained is multiplied by 100 to obtain the ratio of the number of martensite blocks containing metastable carbides to the total number of martensite blocks (ratio p (%)).
《準安定炭化物が存在するマルテンサイトブロックにおける準安定炭化物の数密度の平均値：１×１０６個／ｍｍ２以上》
　低温靭性がより優れるという理由から、マルテンサイトブロックにおける準安定炭化物の数密度が高いことが好ましい。これは、準安定炭化物の数密度が高いと、低温でのマルテンサイト中のき裂の進展抵抗がより大きくなるためと考えられる。
　具体的には、準安定炭化物が存在するマルテンサイトブロックにおける準安定炭化物の数密度の平均値（以下、「数密度ｎ」ともいう）は、１×１０６個／ｍｍ２以上が好ましく、１０×１０６個／ｍｍ２以上がより好ましく、１００×１０６個／ｍｍ２以上がさらに好ましい。
　数密度ｎの上限は、特に限定されず、数密度ｎ、例えば１０００００００×１０６個／ｍｍ２以下であることができ、１００００００×１０６個／ｍｍ２以下が好ましく、１０００００×１０６個／ｍｍ２以下がより好ましく、１００００×１０６個／ｍｍ２以下がさらに好ましい。 Average number density of metastable carbides in martensite blocks containing metastable carbides: 1 x 10 6 /mm 2 or more
A high number density of metastable carbides in the martensite blocks is preferred for better low temperature toughness reasons, as it is believed that a high number density of metastable carbides provides greater resistance to crack propagation in the martensite at low temperatures.
Specifically, the average number density of metastable carbides in a martensite block in which metastable carbides are present (hereinafter also referred to as "number density n") is preferably 1 x 10 pcs/mm 2 or more, more preferably 10 x 10 pcs/mm 2 or more, and even more preferably 100 x 10 pcs/mm 2 or more.
The upper limit of the number density n is not particularly limited, and the number density n can be, for example, 10,000,000×10 6 pieces/mm 2 or less, preferably 1,000,000×10 6 pieces/mm 2 or less, more preferably 100,000×10 6 pieces/mm 2 or less, and even more preferably 10,000×10 6 pieces/mm 2 or less.
　準安定炭化物が存在するマルテンサイトブロックにおける準安定炭化物の数密度の平均値（数密度ｎ）の測定方法は、以下のとおりである。
　上述したＴＥＭを用いた割合ｐの測定に際して、準安定炭化物が存在する単一のマルテンサイトブロックにおいて制限視野電子回折図形を取得し、準安定炭化物から得られた電子回折斑点を用いて、暗視野像を得る。暗視野像において、準安定炭化物は、白いコントラストを呈する。 The method for measuring the average number density (number density n) of metastable carbides in a martensite block in which metastable carbides are present is as follows.
In measuring the fraction p using the above-mentioned TEM, a selected area electron diffraction pattern is obtained for a single martensite block in which metastable carbides are present, and a dark-field image is obtained using the electron diffraction spots obtained from the metastable carbides. In the dark-field image, the metastable carbides show white contrast.
　単一のマルテンサイトブロック内部で３００ｎｍ×３００ｎｍの領域を撮影し、準安定炭化物の個数を数える。なお、３００ｎｍ×３００ｎｍの領域に、ブロック境界を介した隣接するマルテンサイトブロックが存在していても構わない。
　準安定炭化物が存在するマルテンサイトブロックの面積を、制限視野電子回折図形を得た単一のマルテンサイトブロックの面積と定義する。隣接するマルテンサイトブロックどうしは、ブロック境界を介して結晶方位が異なるので、明視野像でコントラストが異なることから、互いに区別される。
　上記の測定を３視野で行ない、準安定炭化物の個数を、準安定炭化物が存在するマルテンサイトブロックの面積で除した値（＝準安定炭化物の個数／準安定炭化物が存在するマルテンサイトブロックの面積）を３視野分求める。それらの平均値を、準安定炭化物が存在するマルテンサイトブロックにおける準安定炭化物の数密度の平均値（数密度ｎ）とする。 An area of 300 nm×300 nm is photographed within a single martensite block, and the number of metastable carbides is counted. Note that adjacent martensite blocks may exist across a block boundary within the 300 nm×300 nm area.
The area of a martensite block with metastable carbides is defined as the area of a single martensite block for which a selected-area electron diffraction pattern was obtained. Adjacent martensite blocks are distinguished from each other by their contrast in bright-field images due to different crystal orientations across the block boundaries.
The above measurement is carried out in three visual fields, and the number of metastable carbides is divided by the area of the martensite block in which the metastable carbides are present (= number of metastable carbides / area of the martensite block in which the metastable carbides are present) for each of the three visual fields. The average of these values is regarded as the average number density (number density n) of the metastable carbides in the martensite block in which the metastable carbides are present.
《準安定炭化物の円相当径の平均値：２０ｎｍ以下》
　マルテンサイトブロックにおける準安定炭化物の円相当径の平均値が小さいほど、低温でマルテンサイト中でき裂が発生しにくくなるために、低温靭性がより優れる。このため、マルテンサイトブロックにおける準安定炭化物の円相当径の平均値は、２０ｎｍ以下が好ましく、５ｎｍ以下がより好ましい。 Average circle equivalent diameter of metastable carbides: 20 nm or less
The smaller the average circle-equivalent diameter of the metastable carbides in the martensite blocks, the less likely cracks will occur in the martensite at low temperatures, and the better the low-temperature toughness will be. For this reason, the average circle-equivalent diameter of the metastable carbides in the martensite blocks is preferably 20 nm or less, and more preferably 5 nm or less.
　マルテンサイトブロックにおける準安定炭化物の円相当径の平均値の測定方法は、以下のとおりである。
　上述したＴＥＭを用いた割合ｐの測定に際して、準安定炭化物が存在する単一のマルテンサイトブロックにおいて制限視野電子回折図形を取得し、準安定炭化物から得られた電子回折斑点を用いて、暗視野像を得る。暗視野像において、準安定炭化物は、白いコントラストを呈する。
　単一のマルテンサイトブロック内部で３００ｎｍ×３００ｎｍの領域の暗視野像を撮影し、画像処理を実施して、準安定炭化物が区別できるように二値化画像を得る。二値化画像を粒子解析することにより、全ての準安定炭化物粒子それぞれについて、円相当径を求める。準安定炭化物どうしが暗視野像において重なっている場合は、二値化画像に対して、Ｗａｔｅｒｓｈｅｄ法を用いた分割を実施する。
　３００ｎｍ×３００ｎｍの領域に存在する全ての準安定炭化物それぞれについて、円相当径を求める（３視野分）。３視野分の円相当径の平均値を求め、これを、マルテンサイトブロックにおける準安定炭化物の円相当径の平均値とする。 The method for measuring the average value of the circle equivalent diameter of metastable carbides in a martensite block is as follows.
In measuring the fraction p using the above-mentioned TEM, a selected area electron diffraction pattern is obtained for a single martensite block in which metastable carbides are present, and a dark-field image is obtained using the electron diffraction spots obtained from the metastable carbides. In the dark-field image, the metastable carbides show white contrast.
A dark field image of a 300 nm x 300 nm area within a single martensite block is taken and image processing is performed to obtain a binary image in which metastable carbides can be distinguished. The binary image is subjected to particle analysis to determine the circle equivalent diameter for each metastable carbide particle. If metastable carbides overlap in the dark field image, the binary image is segmented using the Watershed method.
The circle equivalent diameter is determined for each of all metastable carbides present in the 300 nm × 300 nm region (three visual fields). The average of the circle equivalent diameters for the three visual fields is determined and this is set as the average circle equivalent diameter of the metastable carbides in the martensite block.
〈ナノ硬度〉
　鋼板のナノ硬度の標準偏差σnを低下させることで、優れた部品強度、延性、伸びフランジ性、低温靭性を保ちつつ、剪断端面曲げ性が向上する。これは局所領域での塑性変形抵抗が組織内で均一化することで、剪断加工部での塑性変形不均一性が抑制され、剪断加工部での曲げ変形能が向上するためと考えられる。
　ここで、例えばビッカース硬さ試験といったナノインデンテーション法以外の硬さ試験法では、組織のサブミクロンレベルの局所領域での塑性変形抵抗は得られない。よって、ナノインデンテーション法を用いることで初めて本発明の課題を解決することができる。 <Nano hardness>
By reducing the standard deviation σ n of the nanohardness of steel sheets, the shear edge bendability is improved while maintaining excellent part strength, ductility, stretch flangeability, and low-temperature toughness. This is thought to be because the plastic deformation resistance in localized regions is made uniform within the structure, suppressing the non-uniformity of plastic deformation in the sheared area and improving the bending deformability in the sheared area.
Here, the plastic deformation resistance in a local region of the structure at the submicron level cannot be obtained by a hardness test method other than the nanoindentation method, such as a Vickers hardness test, and therefore the problem of the present invention can be solved only by using the nanoindentation method.
　上記の効果を得るため、ナノ硬度の標準偏差σnは、ナノ硬度の平均値［Ｈｎ］aveに対して、０．６０×［Ｈｎ］ave以下であり、好ましくは０．５０×［Ｈｎ］ave以下である。ナノ硬度の標準偏差σnの下限は、特に限定されず、０であってもよい。 In order to obtain the above-mentioned effect, the standard deviation σ n of the nanohardness is 0.60×[H n ] ave or less, preferably 0.50×[H n ] ave or less, where [H n ] ave is the average value of the nanohardness. The lower limit of the standard deviation σ n of the nanohardness is not particularly limited and may be 0.
　ナノ硬度の平均値［Ｈｎ］aveは、好ましくは３．０ＧＰａ以上９．０ＧＰａ以下であり、より好ましくは３．５ＧＰａ以上８．５ＧＰａ以下である。 The average nano-hardness [Hn] ave is preferably 3.0 GPa or more and 9.0 GPa or less, and more preferably 3.5 GPa or more and 8.5 GPa or less.
　ここで、ナノ硬度の標準偏差σnの測定方法を説明する。ナノ硬度を求めるためにはバーコビッチ圧子を備えたナノインデンテーション装置を用いる。
　鋼板の圧延方向に平行な板厚断面（Ｌ断面）が観察面となるよう試料を切り出した後、観察面をダイヤモンドペーストを用いて鏡面研磨し、その後、コロイダルシリカを用い仕上げ研磨を施す。ナノインデンテーション装置を用いて、荷重制御で、荷重速度及び除荷速度を５０μＮ／ｓ、最大荷重を５００μＮ、データ採取ピッチを５ｍｓｅｃとする条件で、試料について２２５点以上のナノ硬度を測定する。測定の際、高強度鋼板表面から板厚１／４位置を測定位置とし、圧痕間の距離は２μｍ以上空けることとする。
　得られた２２５点以上のナノ硬度測定結果からヒストグラムを作成し、標準偏差を求め、その値をナノ硬度の標準偏差σnとする。得られた２２５点以上のナノ硬度測定結果の平均値を、［Ｈｎ］aveとする。 Here, a method for measuring the standard deviation σ n of nanohardness will be described. In order to measure the nanohardness, a nanoindentation device equipped with a Berkovich indenter is used.
After cutting out a sample so that the plate thickness cross section (L cross section) parallel to the rolling direction of the steel plate is the observation surface, the observation surface is mirror-polished using diamond paste, and then finish-polished using colloidal silica. Using a nanoindentation device, the nanohardness of 225 or more points is measured for the sample under load control conditions of a loading rate and unloading rate of 50 μN/s, a maximum load of 500 μN, and a data collection pitch of 5 msec. During the measurement, the measurement position is set to 1/4 the plate thickness from the surface of the high-strength steel plate, and the distance between the indentations is set to 2 μm or more.
A histogram is created from the nanohardness measurement results obtained at 225 or more points, and the standard deviation is calculated and the result is defined as the standard deviation of the nanohardness σ n . The average value of the nanohardness measurement results obtained at 225 or more points is defined as [H n ] ave .
〈表層軟質層〉
　高強度鋼板には、素地鋼板の表層において表層軟質層が形成されていることが好ましい。プレス成形時及び車体衝突時に前記表層軟質層が曲げ割れ進展の抑制に寄与するため、耐曲げ破断特性を向上する。
　ここで、素地鋼板は、溶融亜鉛めっき鋼板、合金化溶融亜鉛めっき鋼板、電気亜鉛めっき鋼板及びその他金属めっき鋼板といっためっき処理が施されている鋼板の場合は、前記の各種めっきの素地（下地）である高強度鋼板であり、めっき処理が施されていない場合は、高強度鋼板であることとする。
　表層とは、素地鋼板表面から板厚方向深さ２００μｍまでの厚さ２００μｍに対応する領域をいう。
　軟質層は、素地鋼板の板厚１／４位置の断面（鋼板表面に平行な面）のビッカース硬さに対して、８５％以下のビッカース硬さの領域をいう。軟質層は、素地鋼板の表層における脱炭層を包含する。
　表層軟質層は、表層に含まれる軟質層をいい、表層全体が軟質層であっても、表層の一部が軟質層であってもよい。表層軟質層は、素地鋼板表面から板厚方向に２００μｍ以内の厚さに対応する領域であることができる。
　例えば、素地鋼板の板厚１／４位置の断面（鋼板表面に平行な面）のビッカース硬さに対して８５％以下の領域が、素地鋼板表面から板厚方向に所定の深さで形成されているとして、所定の深さが板厚方向に２００μｍ以内の場合、表面から板厚方向の所定の深さまでの厚さに対応する領域が表層軟質層であり、所定の深さが板厚方向に２００μｍ超の場合、素地鋼板表面から板厚方向深さ２００μｍまでの厚さ２００μｍに対応する領域が表層軟質層である。
　表層軟質層を有する場合、表層軟質層の厚さの下限は特に限定されず、８μｍ以上が好ましく、１７μｍ超がより好ましい。
　ビッカース硬さは、ＪＩＳ　Ｚ　２２４４－１（２０２０）に基づいて、荷重を１０ｇｆとして測定する。 <Soft surface layer>
It is preferable that the high-strength steel sheet has a soft surface layer formed on the surface of the base steel sheet. The soft surface layer contributes to suppressing the propagation of bending cracks during press forming and vehicle body collision, thereby improving bending fracture resistance.
Here, the base steel sheet refers to a high-strength steel sheet that is the base (undercoat) for the various platings in the case of a steel sheet that has been subjected to a plating treatment, such as a hot-dip galvanized steel sheet, a galvannealed steel sheet, an electrolytic galvanized steel sheet, or a steel sheet that has been plated with other metals, and refers to a high-strength steel sheet in the case of a steel sheet that has not been plated.
The surface layer refers to a region corresponding to a thickness of 200 μm from the surface of the base steel sheet to a depth of 200 μm in the sheet thickness direction.
The soft layer refers to a region having a Vickers hardness of 85% or less of the Vickers hardness of a cross section (a plane parallel to the steel sheet surface) at 1/4 of the sheet thickness of the base steel sheet. The soft layer includes a decarburized layer in the surface layer of the base steel sheet.
The surface soft layer refers to a soft layer included in the surface layer, and may be a soft layer in its entirety or a part of it. The surface soft layer may be a region corresponding to a thickness of 200 μm or less from the surface of the base steel sheet in the sheet thickness direction.
For example, if a region having a Vickers hardness of 85% or less on a cross section (plane parallel to the steel plate surface) at 1/4 of the plate thickness of the base steel plate is formed at a predetermined depth from the surface of the base steel plate in the plate thickness direction, when the predetermined depth is within 200 μm in the plate thickness direction, the region corresponding to the thickness from the surface to the predetermined depth in the plate thickness direction is the surface soft layer, and when the predetermined depth is more than 200 μm in the plate thickness direction, the region corresponding to a thickness of 200 μm from the surface of the base steel plate to a depth of 200 μm in the plate thickness direction is the surface soft layer.
When the soft surface layer is provided, the lower limit of the thickness of the soft surface layer is not particularly limited, but is preferably 8 μm or more, and more preferably more than 17 μm.
The Vickers hardness is measured based on JIS Z 2244-1 (2020) at a load of 10 gf.
　プレス成形時の優れた曲げ性と衝突時の優れた曲げ破断特性を得るためには、素地鋼板表面から表層軟質層の板厚方向深さの１／４位置（素地鋼板表面から深さ方向に表層軟質層の厚さの１／４位置）の板面の５０μｍ×５０μｍの領域において、３００点以上のナノ硬度を測定したとき、ナノ硬度が７．０ＧＰａ以上の割合が０．１０以下である必要があることが好ましい。ナノ硬度が７．０ＧＰａ以上の割合が０．１０以下の場合、硬質な組織（マルテンサイトなど）、介在物などの割合が小さいことを意味し、硬質な組織（マルテンサイトなど）、介在物などのプレス成形時及び衝突時のボイドの生成や連結、さらには亀裂の進展をより抑制することが可能となり、優れたＲ／ｔ及びＳＦｍａｘが得られる。 In order to obtain excellent bendability during press forming and excellent bending fracture properties during collision, when the nano hardness is measured at 300 or more points in a 50 μm×50 μm region of the sheet surface at a 1/4 position from the surface of the base steel sheet to the depth in the sheet thickness direction of the soft surface layer (at a 1/4 position from the surface of the base steel sheet to the thickness of the soft surface layer in the depth direction), it is preferable that the proportion of nano hardness of 7.0 GPa or more should be 0.10 or less. When the proportion of nano hardness of 7.0 GPa or more is 0.10 or less, it means that the proportion of hard structures (martensite, etc.), inclusions, etc. is small, and it is possible to further suppress the generation and connection of voids and crack growth in hard structures (martensite, etc.) and inclusions during press forming and collision, and to obtain excellent R/t and SF max .
　本発明において、プレス成形時の優れた曲げ性と衝突時の優れた曲げ破断特性を得るためには、素地鋼板表面から表層軟質層の板厚方向深さの１／４位置の板面のナノ硬度の標準偏差σが１．８ＧＰａ以下であり、さらに、素地鋼板表面から表層軟質層の板厚方向深さの１／２位置の板面のナノ硬度の標準偏差σが２．２ＧＰａ以下であることが好ましい。素地鋼板表面から表層軟質層の板厚方向深さの１／４位置の板面のナノ硬度の標準偏差σが１．８ＧＰａ以下であり、さらに、素地鋼板表面から表層軟質層の板厚方向深さの１／２位置の板面のナノ硬度の標準偏差σが２．２ＧＰａ以下の場合、ミクロ領域における組織硬度差が小さいことを意味し、プレス成形時及び衝突時のボイドの生成や連結、さらには亀裂の進展をより抑制することが可能となり、優れたＲ／ｔ及びＳＦｍａｘが得られる。 In the present invention, in order to obtain excellent bendability during press forming and excellent bending fracture properties during collision, it is preferable that the standard deviation σ of the nanohardness of the sheet surface at a position 1/4 of the sheet thickness direction depth of the soft surface layer from the surface of the base steel sheet is 1.8 GPa or less, and further, the standard deviation σ of the nanohardness of the sheet surface at a position 1/2 of the sheet thickness direction depth of the soft surface layer from the surface of the base steel sheet is 2.2 GPa or less. When the standard deviation σ of the nanohardness of the sheet surface at a position 1/4 of the sheet thickness direction depth of the soft surface layer from the surface of the base steel sheet is 1.8 GPa or less, and further, the standard deviation σ of the nanohardness of the sheet surface at a position 1/2 of the sheet thickness direction depth of the soft surface layer from the surface of the base steel sheet is 2.2 GPa or less, it means that the difference in microscopic hardness is small, and it is possible to further suppress the generation and connection of voids during press forming and collision, and further the growth of cracks, and excellent R/t and SF max are obtained.
　また、素地鋼板表面から表層軟質層の板厚方向深さの１／４位置の板面のナノ硬度の標準偏差σのより好ましい範囲は、１．７ＧＰａ以下である。素地鋼板表面から表層軟質層の板厚方向深さの１／２位置の板面のナノ硬度の標準偏差σのより好ましい範囲は、２．１ＧＰａ以下である。 In addition, a more preferred range for the standard deviation σ of the nanohardness of the sheet surface at a position 1/4 of the way from the base steel sheet surface to the soft surface layer in the sheet thickness direction is 1.7 GPa or less. A more preferred range for the standard deviation σ of the nanohardness of the sheet surface at a position 1/2 of the way from the base steel sheet surface to the soft surface layer in the sheet thickness direction is 2.1 GPa or less.
　ここで、板厚方向深さの１／４位置、１／２位置の板面のナノ硬度とは、以下の方法により測定される硬度である。
　まず、めっき層が形成されている場合は、めっき層剥離後、素地鋼板表面から表層軟質層の板厚方向深さの１／４位置まで機械研磨を実施し、ダイヤモンド及びアルミナでのバフ研磨を実施し、さらにコロイダルシリカ研磨を実施する。バーコビッチ形状のダイヤモンド圧子により、荷重：５００μＮ、測定領域：５０μｍ×５０μｍ、打点間隔：２μｍの条件でナノ硬度を測定する。
　また、表層軟質層の板厚方向深さの１／２位置まで機械研磨を実施し、ダイヤモンド及びアルミナでのバフ研磨を実施、さらにコロイダルシリカ研磨を実施する。そして、バーコビッチ形状のダイヤモンド圧子により、荷重：５００μＮ、測定領域：５０μｍ×５０μｍ、打点間隔：２μｍの条件でナノ硬度を測定する。 Here, the nano-hardness of the plate surface at the 1/4 and 1/2 positions in the plate thickness direction depth is a hardness measured by the following method.
First, if a plating layer is formed, after peeling off the plating layer, mechanical polishing is performed from the surface of the base steel sheet to a position 1/4 of the depth in the sheet thickness direction of the soft surface layer, buff polishing with diamond and alumina is performed, and further colloidal silica polishing is performed. The nano hardness is measured using a Berkovich-shaped diamond indenter under the following conditions: load: 500 μN, measurement area: 50 μm × 50 μm, and impact spacing: 2 μm.
The soft surface layer is mechanically polished to a position halfway down the thickness direction, buffed with diamond and alumina, and then polished with colloidal silica.Then, the nano-hardness is measured with a Berkovich diamond indenter under the following conditions: load: 500 μN, measurement area: 50 μm × 50 μm, and impact spacing: 2 μm.
　ここで、表層軟質層の厚さは、以下の方法により測定することができる。素地鋼板の圧延方向に平行な板厚断面（Ｌ断面）を湿式研磨により平滑化した後、ビッカース硬度計を用いて、荷重１０ｇｆで、素地鋼板表面から板厚方向に１μｍの位置より、板厚方向１００μｍの位置まで、１μｍ間隔で測定を行った。その後は板厚中心まで２０μｍ間隔で測定を行った。硬度が板厚１／４位置の硬度に比して８５％以下に減少した領域を軟質層（表層軟質層）と定義し、当該領域の板厚方向の厚さを軟質層の厚さとた。 Here, the thickness of the surface soft layer can be measured by the following method. After smoothing the thickness cross section (L cross section) of the base steel sheet parallel to the rolling direction by wet polishing, measurements were taken at 1 μm intervals using a Vickers hardness tester with a load of 10 gf from a position 1 μm from the surface of the base steel sheet in the thickness direction to a position 100 μm in the thickness direction. After that, measurements were taken at 20 μm intervals up to the center of the thickness. The area where the hardness has decreased to 85% or less compared to the hardness at the 1/4 position in the thickness direction was defined as the soft layer (surface soft layer), and the thickness in the thickness direction of this area was taken as the thickness of the soft layer.
〈第一めっき層〉
　本発明の一実施形態に伴う高強度鋼板は、素地鋼板の片面又は両面の表面上において、金属めっき層である第一めっき層を有することが好ましい。第一めっき層は、素地鋼板表面に直接形成されており、Ｃｒ、Ｍｎ、Ｆｅ、Ｃｏ、Ｎｉ、Ｃｕ、Ｇａ、Ｇｅ、Ａｓ、Ｒｕ、Ｒｈ、Ｐｄ、Ａｇ、Ｃｄ、Ｉｎ、Ｓｎ、Ｓｂ、Ｏｓ、Ｉｒ、Ｒｔ、Ａｕ、Ｈｇ、Ｔｉ、Ｐｂ及びＢｉから選択される１種又は２種以上を合計で５０質量％超含む金属めっき層であり、溶融亜鉛めっき層、合金化溶融亜鉛めっき層、電気亜鉛めっき層の亜鉛めっき層、溶融アルミニウムめっき層は除かれる。第一めっき層は、金属電気めっき層が好ましく、以下では、金属電気めっき層を例に説明する。 <First plating layer>
The high-strength steel sheet according to an embodiment of the present invention preferably has a first plating layer, which is a metal plating layer, on one or both surfaces of the base steel sheet. The first plating layer is formed directly on the surface of the base steel sheet and is a metal plating layer containing more than 50 mass% in total of one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb and Bi, and is not a hot-dip galvanized layer, an alloyed hot-dip galvanized layer, a zinc plating layer of an electrogalvanized layer, or a hot-dip aluminum plating layer. The first plating layer is preferably a metal electroplated layer, and the following description will be given taking a metal electroplated layer as an example.
　金属電気めっき層が鋼板表面に形成されることで、プレス成形時及び車体衝突時に最表層の前記金属電気めっき層が曲げ割れ発生の抑制に寄与するため、耐曲げ破断特性がさらに向上する。 By forming a metal electroplating layer on the surface of the steel sheet, the outermost metal electroplating layer helps to prevent bending cracks during press forming and vehicle collisions, further improving bending fracture resistance.
　金属電気めっき層の金属種としては、Ｃｒ、Ｍｎ、Ｆｅ、Ｃｏ、Ｎｉ、Ｃｕ、Ｇａ、Ｇｅ、Ａｓ、Ｒｕ、Ｒｈ、Ｐｄ、Ａｇ、Ｃｄ、Ｉｎ、Ｓｎ、Ｓｂ、Ｏｓ、Ｉｒ、Ｒｔ、Ａｕ、Ｈｇ、Ｔｉ、Ｐｂ、Ｂｉのいずれでもかまわないが、Ｆｅであることがより好ましい。以下では、Ｆｅ系電気めっき層を例に説明する。 The metal type of the metal electroplating layer may be any of Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi, but Fe is more preferable. The following describes an example of an Fe-based electroplating layer.
　Ｆｅ系電気めっき層の付着量は、０ｇ／ｍ２超とし、好ましくは２．０ｇ／ｍ２以上とする。Ｆｅ系電気めっき層の片面あたりの付着量の上限は特に限定されないが、コストの観点から、Ｆｅ系電気めっき層の片面あたりの付着量を６０ｇ／ｍ２以下とすることが好ましい。Ｆｅ系電気めっき層の付着量は、好ましくは５０ｇ／ｍ２以下であり、より好ましくは４０ｇ／ｍ２以下であり、さらに好ましくは３０ｇ／ｍ２以下とする。 The coating weight of the Fe-based electroplating layer is more than 0 g/ m2 , and preferably 2.0 g/ m2 or more. There is no particular upper limit to the coating weight of the Fe-based electroplating layer per side, but from the viewpoint of cost, the coating weight of the Fe-based electroplating layer per side is preferably 60 g/ m2 or less. The coating weight of the Fe-based electroplating layer is preferably 50 g/ m2 or less, more preferably 40 g/ m2 or less, and even more preferably 30 g/ m2 or less.
　Ｆｅ系電気めっき層の付着量は、以下のとおり測定する。Ｆｅ系電気めっき鋼板から１０×１５ｍｍサイズのサンプルを採取して樹脂に埋め込み、断面埋め込みサンプルとする。同断面の任意の３か所を走査型電子顕微鏡（Ｓｃａｎｎｉｎｇ Ｅｌｅｃｔｒｏｎ Ｍｉｃｒｏｓｃｏｐｅ；ＳＥＭ）を用いて加速電圧１５ｋＶで、Ｆｅ系めっき層の厚みに応じて倍率２０００～１００００倍で観察し、３視野の厚みの平均値に鉄の比重を乗じることによって、Ｆｅ系めっき層の片面あたりの付着量に換算する。 The adhesion weight of the Fe-based electroplating layer is measured as follows. A 10 x 15 mm sample is taken from the Fe-based electroplated steel sheet and embedded in resin to create a cross-section embedded sample. Three random locations on the cross section are observed using a scanning electron microscope (SEM) at an accelerating voltage of 15 kV and a magnification of 2,000 to 10,000 times depending on the thickness of the Fe-based plating layer, and the average thickness of the three fields of view is multiplied by the specific gravity of iron to convert it into the adhesion weight of the Fe-based plating layer per side.
　Ｆｅ系電気めっき層としては、純Ｆｅの他、Ｆｅ－Ｂ合金、Ｆｅ－Ｃ合金、Ｆｅ－Ｐ合金、Ｆｅ－Ｎ合金、Ｆｅ－Ｏ合金、Ｆｅ－Ｎｉ合金、Ｆｅ－Ｍｎ合金、Ｆｅ－Ｍｏ合金、Ｆｅ－Ｗ合金等の合金めっき層が使用できる。Ｆｅ系電気めっき層の成分組成は特に限定されないが、Ｂ、Ｃ、Ｐ、Ｎ、Ｏ、Ｎｉ、Ｍｎ、Ｍｏ、Ｚｎ、Ｗ、Ｐｂ、Ｓｎ、Ｃｒ、Ｖ及びＣｏからなる群から選ばれる１又は２以上の元素を合計で１０質量％以下含み、残部はＦｅ及び不可避的不純物からなる成分組成とすることが好ましい。Ｆｅ以外の元素の量を合計で１０質量％以下とすることで、電解効率の低下を防ぎ、低コストでＦｅ系電気めっき層を形成することができる。Ｆｅ－Ｃ合金の場合、Ｃの含有量は０．０８質量％以下とすることが好ましい。 As the Fe-based electroplating layer, in addition to pure Fe, alloy plating layers such as Fe-B alloy, Fe-C alloy, Fe-P alloy, Fe-N alloy, Fe-O alloy, Fe-Ni alloy, Fe-Mn alloy, Fe-Mo alloy, and Fe-W alloy can be used. The composition of the Fe-based electroplating layer is not particularly limited, but it is preferable that the composition contains one or more elements selected from the group consisting of B, C, P, N, O, Ni, Mn, Mo, Zn, W, Pb, Sn, Cr, V, and Co in a total amount of 10 mass% or less, with the remainder consisting of Fe and unavoidable impurities. By keeping the amount of elements other than Fe to a total of 10 mass% or less, it is possible to prevent a decrease in electrolysis efficiency and form the Fe-based electroplating layer at low cost. In the case of an Fe-C alloy, the C content is preferably 0.08 mass% or less.
〈第二めっき層〉
　本発明の一実施形態に従う高強度鋼板は、高強度鋼板の片面又は両面の最外層として、金属めっき層である第二めっき層を有していてもよい。第二めっき層は、亜鉛及びアルミニウムの少なくとも一方を合計で５０質量％以上含み、溶融亜鉛めっき層、合金化溶融亜鉛めっき層、電気亜鉛めっき層、溶融アルミニウムめっき層等であることができる。 <Second plating layer>
The high-strength steel sheet according to one embodiment of the present invention may have a second plating layer, which is a metal plating layer, as an outermost layer on one or both sides of the high-strength steel sheet. The second plating layer contains at least one of zinc and aluminum in a total amount of 50 mass % or more, and may be a hot-dip galvanized layer, a galvannealed layer, an electrolytic galvanized layer, a hot-dip aluminum plating layer, or the like.
　第二めっき層は、素地鋼板表面の片面又は両面に直接形成されていてもよく、第一めっき層上に形成されていてもよい。 The second plating layer may be formed directly on one or both surfaces of the base steel sheet surface, or it may be formed on the first plating layer.
　ここで、溶融亜鉛めっき層、合金化溶融亜鉛めっき層、電気亜鉛めっき層は、Ｚｎ（亜鉛）を主成分（Ｚｎ含有量が５０．０質量％以上）とするめっき層をいう。
　また、アルミニウムめっき鋼板のめっき層は、Ａｌ（アルミニウム）を主成分（Ａｌ含有量が５０．０質量％以上）とするめっき層をいう。 Here, the hot-dip galvanized layer, alloyed hot-dip galvanized layer, and electrolytic galvanized layer refer to a plating layer containing Zn (zinc) as a main component (Zn content of 50.0 mass % or more).
The plating layer of an aluminum-plated steel sheet refers to a plating layer containing Al (aluminum) as a main component (Al content is 50.0 mass % or more).
　ここで、溶融亜鉛めっき層は、例えば、Ｚｎと、２０．０質量％以下のＦｅ、０．００１質量％以上１．０質量％以下のＡｌにより構成することが好適である。また、溶融亜鉛めっき層には、任意に、Ｐｂ、Ｓｂ、Ｓｉ、Ｓｎ、Ｍｇ、Ｍｎ、Ｎｉ、Ｃｒ、Ｃｏ、Ｃａ、Ｃｕ、Ｌｉ、Ｔｉ、Ｂｅ、Ｂｉ及びＲＥＭからなる群から選ばれる１種又は２種以上の元素を合計で０．０質量％超３．５質量％以下含有させてもよい。また、溶融亜鉛めっき層のＦｅ含有量は、より好ましくは７．０質量％未満である。なお、前記の元素以外の残部は、不可避的不純物である。
　合金化溶融亜鉛めっき層は、例えば、２０質量％以下のＦｅ、０．００１質量％以上１．０質量％以下のＡｌにより構成することが好適である。また、合金化溶融亜鉛めっき層には、任意に、Ｐｂ、Ｓｂ、Ｓｉ、Ｓｎ、Ｍｇ、Ｍｎ、Ｎｉ、Ｃｒ、Ｃｏ、Ｃａ、Ｃｕ、Ｌｉ、Ｔｉ、Ｂｅ、Ｂｉ及びＲＥＭからなる群から選ばれる１種又は２種以上の元素を合計で０質量％超３．５質量％以下含有させてもよい。合金化溶融亜鉛めっき層のＦｅ含有量は、より好ましくは７．０質量％以上、さらに好ましくは８．０質量％以上である。また、合金化溶融亜鉛めっき層のＦｅ含有量は、より好ましくは１５．０質量％以下、さらに好ましくは１２．０質量％以下である。なお、前記の元素以外の残部は、不可避的不純物である。 Here, the hot-dip galvanized layer is preferably composed of, for example, Zn, 20.0 mass% or less Fe, and 0.001 mass% or more and 1.0 mass% or less Al. The hot-dip galvanized layer may optionally contain one or more elements 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 more than 0.0 mass% and 3.5 mass% or less. The Fe content of the hot-dip galvanized layer is more preferably less than 7.0 mass%. The remainder other than the above elements is unavoidable impurities.
The galvannealed layer is preferably composed of, for example, 20% by mass or less Fe and 0.001% by mass or more and 1.0% by mass or less Al. The galvannealed layer may optionally contain one or more elements 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 more than 0% by mass and 3.5% by mass or less. The Fe content of the galvannealed layer is more preferably 7.0% by mass or more, and even more preferably 8.0% by mass or more. The Fe content of the galvannealed layer is more preferably 15.0% by mass or less, and even more preferably 12.0% by mass or less. The remainder other than the above elements is unavoidable impurities.
　加えて、上記亜鉛めっき層の片面あたりのめっき付着量は、特に限定されるものではないが、２０ｇ/ｍ２以上８０ｇ/ｍ２以下とすることが好ましい。 In addition, the plating weight of the zinc plating layer per side is not particularly limited, but is preferably 20 g/m 2 or more and 80 g/m 2 or less.
　上記亜鉛めっき層のめっき付着量は、以下のようにして測定する。１０質量％塩酸水溶液１Ｌに対し、Ｆｅに対する腐食抑制剤（朝日化学工業（株）製「イビット７００ＢＫ」（登録商標））を０．６ｇ添加した処理液を調整する。次いで、該処理液に、亜鉛めっき層を備えた鋼板のサンプルを浸漬し、亜鉛めっき層を溶解させる。そして、溶解前後でのサンプルの質量減少量を測定し、その値を、素地鋼板の表面積（めっきで被覆されていた部分の表面積）で除することにより、めっき付着量（ｇ/ｍ２）を算出する。 The coating weight of the zinc plating layer is measured as follows. A treatment solution is prepared by adding 0.6 g of a corrosion inhibitor for Fe (Ivit 700BK (registered trademark) manufactured by Asahi Chemical Industry Co., Ltd.) to 1 L of a 10 mass % aqueous hydrochloric acid solution. A sample of a steel sheet having a zinc plating layer is then immersed in the treatment solution to dissolve the zinc plating layer. The mass loss of the sample before and after dissolution is then measured, and the value is divided by the surface area of the base steel sheet (the surface area of the part that was covered with plating) to calculate the coating weight (g/ m2 ).
〈その他〉
　高強度鋼板の板厚は特に限定されず、０．３ｍｍ以上３．０ｍｍ以下とすることができる。 <others>
The thickness of the high-strength steel plate is not particularly limited, and can be 0.3 mm or more and 3.0 mm or less.
［高強度鋼板の製造方法］
　次に、本発明の高強度鋼板の製造方法（以下、便宜的に「本発明の製造方法」ともいう）を説明する。本発明の製造方法は、上述した本発明の高強度鋼板を製造する方法でもある。ここで、製造方法に関する温度は、特に断らない限り、いずれも鋼スラブ又は鋼板の表面温度を基準とする。 [Method of manufacturing high-strength steel plate]
Next, a method for producing a high-strength steel plate according to the present invention (hereinafter, for convenience, also referred to as the "production method of the present invention") will be described. The production method of the present invention is also a method for producing the high-strength steel plate according to the present invention described above. Here, the temperature in the production method is based on the surface temperature of the steel slab or steel plate, unless otherwise specified.
〈熱間圧延、酸洗及び冷間圧延〉
　本発明の製造方法においては、まず、上述した本発明の成分組成を有する鋼スラブに、熱間圧延、酸洗及び冷間圧延を施すことにより、冷延板を得る。 <Hot rolling, pickling and cold rolling>
In the manufacturing method of the present invention, first, a steel slab having the above-mentioned component composition of the present invention is subjected to hot rolling, pickling and cold rolling to obtain a cold-rolled sheet.
《鋼スラブの製造》
　鋼スラブとしては、例えば、鋼素材を溶製して本発明の成分組成を有する溶鋼とし、得られた溶鋼を固めたものを用いることができる。溶鋼方法は、特に限定されず、転炉溶鋼、電気炉溶鋼等の公知の溶製方法を用いることができる。溶鋼から鋼スラブを製造する方法は特に限定されず、連続鋳造法、造塊法、薄スラブ鋳造法等の公知の方法を用いることができる。マクロ偏析を防止する点から、連続鋳造法で製造することが好ましい。 <Production of Steel Slabs>
The steel slab may be, for example, a molten steel having the composition of the present invention obtained by melting a steel material and solidifying the molten steel. The method for melting steel is not particularly limited, and known melting methods such as converter molten steel and electric furnace molten steel can be used. The method for producing a steel slab from molten steel is not particularly limited, and known methods such as continuous casting, ingot casting, and thin slab casting can be used. From the viewpoint of preventing macrosegregation, it is preferable to produce the steel slab by the continuous casting method.
《熱間圧延工程》
　製造された鋼スラブを、例えば、一旦室温まで冷却した後、再び加熱して熱間圧延（粗圧延及び仕上げ圧延）を施し、その後、巻き取りする。こうして、熱延板が得られる。ただし、製造した鋼スラブを、室温まで冷却しないで温片のまま加熱炉に装入してもよいし、わずかに保熱した後に直ちに粗圧延してもよい。 <Hot rolling process>
The produced steel slab is, for example, cooled to room temperature once, then heated again and hot rolled (rough rolling and finish rolling), and then coiled. In this way, a hot-rolled sheet is obtained. However, the produced steel slab may be charged into a heating furnace as a hot piece without being cooled to room temperature, or may be roughly rolled immediately after being slightly kept at room temperature.
（粗圧延）　
　鋼スラブを以下の条件で粗圧延することにより、粗圧延板が得られる。
　鋼スラブを加熱する温度（スラブ加熱温度）は、炭化物の溶解や圧延荷重の低減の観点から、１１００℃以上が好ましい。一方、スケールロスの増大を防止するため、スラブ加熱温度は、１３００℃以下が好ましい。スラブ加熱温度まで加熱した鋼スラブに対して、粗圧延を行う。 (Rough rolling)
A rough rolled sheet is obtained by rough rolling the steel slab under the following conditions.
The temperature at which the steel slab is heated (slab heating temperature) is preferably 1100° C. or higher from the viewpoint of dissolving carbides and reducing the rolling load. On the other hand, in order to prevent an increase in scale loss, the slab heating temperature is preferably 1300° C. or lower. The steel slab heated to the slab heating temperature is subjected to rough rolling.
　粗圧延中の平均のひずみ速度と総圧下率を適正化することで、ナノ硬度の標準偏差が低下し、０．６０×［Ｈｎ］ave以下とすることができる。これは粗圧延中のオーステナイト粒の塑性変形及び動的再結晶過程においてＳｉやＭｎといった溶質原子が転位や再結晶粒の粒界を介して高速拡散することで、適正分布し、局所領域における塑性変形抵抗が均一化するためと考えられる。 By optimizing the average strain rate and total reduction rate during rough rolling, the standard deviation of nanohardness can be reduced to 0.60 × [H n ] ave or less. This is thought to be because solute atoms such as Si and Mn rapidly diffuse through dislocations and grain boundaries of recrystallized grains during the plastic deformation and dynamic recrystallization of austenite grains during rough rolling, resulting in an appropriate distribution and uniform plastic deformation resistance in local regions.
　ここで、粗圧延中の平均のひずみ速度とは、粗圧延の最初のミルから最終のミルまででの総圧延率ε（－）を、粗圧延の最初のミルでの圧延開始から最終のミルでの圧延完了までに要した時間ｔＲ（ｓ）で割った値（ε／ｔＲ）と定義する。 Here, the average strain rate during rough rolling is defined as the total rolling reduction ε(-) from the first mill to the final mill in rough rolling divided by the time t R (s) required from the start of rolling at the first mill to the completion of rolling at the final mill in rough rolling (ε/t R ).
　粗圧延中の平均のひずみ速度が１×１０－１／s超又は総圧下率が５０％未満の場合、オーステナイト粒の塑性変形及び動的再結晶中のＳｉやＭｎといった溶質原子の拡散が不十分となり、ナノ硬度の標準偏差が０．６０×［Ｈｎ］ave超となる。
　一方、粗圧延中の平均のひずみ速度が１×１０－４／s未満のときは、オーステナイト粒の転位の回復が促進され、再結晶の駆動力が低下し動的再結晶が抑制されることでＳｉやＭｎといった溶質原子の拡散が不十分となり、ナノ硬度の標準偏差が０．６０×［Ｈｎ］ave超となる。
　よって、粗圧延は、平均のひずみ速度は１×１０－４／s以上１×１０－１／s以下、総圧下率５０％以上の条件で行うこととする。粗圧延中の平均のひずみ速度は好ましくは１×１０－３／s以上１×１０－２／ｓ以下である。粗圧延中の総圧下率は好ましくは６０％以上である。 If the average strain rate during rough rolling exceeds 1×10 −1 /s or the total rolling reduction is less than 50%, the plastic deformation of the austenite grains and the diffusion of solute atoms such as Si and Mn during dynamic recrystallization become insufficient, and the standard deviation of the nanohardness exceeds 0.60×[H n ] ave .
On the other hand, when the average strain rate during rough rolling is less than 1 × 10 −4 /s, the recovery of dislocations in austenite grains is promoted, the driving force for recrystallization is reduced, and dynamic recrystallization is suppressed, resulting in insufficient diffusion of solute atoms such as Si and Mn, and the standard deviation of nanohardness exceeds 0.60 × [H n ] ave .
Therefore, rough rolling is performed under the conditions of an average strain rate of 1×10 −4 /s to 1×10 −1 /s and a total reduction of 50% or more. The average strain rate during rough rolling is preferably 1×10 −3 /s to 1×10 −2 /s. The total reduction during rough rolling is preferably 60% or more.
　粗圧延終了温度はオーステナイト粒の再結晶の完了させる観点から、９５０℃以上とすることが好ましい。粗圧延終了温度は、例えば、１２５０℃以下とすることができる。
　スラブ加熱温度を低めにした場合は、熱間圧延におけるトラブルを防止する観点から、仕上げ圧延の前に、バーヒーター等を用いて粗圧延板を加熱することが好ましい。 From the viewpoint of completing the recrystallization of austenite grains, the rough rolling end temperature is preferably set to 950° C. or more. The rough rolling end temperature can be set to, for example, 1250° C. or less.
When the slab heating temperature is set to a low temperature, it is preferable to heat the roughly rolled sheet using a bar heater or the like before finish rolling in order to prevent problems during hot rolling.
（仕上げ圧延）
　仕上げ圧延を実施する際の温度（仕上げ圧延温度）は、Ａｒ３変態点以上が好ましい。これにより、圧延負荷が低減する。さらに、オーステナイトの未再結晶状態での圧下率が低下し、圧延方向に伸長した異常な組織の発達が抑制され、加工性が優れる。 (Finish rolling)
The temperature when performing the finish rolling (finish rolling temperature) is preferably equal to or higher than the Ar3 transformation point. This reduces the rolling load. Furthermore, the rolling reduction in the non-recrystallized state of austenite is reduced, and the development of abnormal structures elongated in the rolling direction is suppressed, resulting in excellent workability.
　仕上げ圧延は、粗圧延板どうしを接合して連続的に実施してもよい。仕上げ圧延を実施する前に、粗圧延板を一旦巻き取ってもよい。 Finish rolling may be performed continuously by joining the rough rolled sheets together. The rough rolled sheets may be wound up once before finishing rolling is performed.
　圧延荷重を低減するために、仕上げ圧延の一部又は全部を、潤滑圧延としてもよい。潤滑圧延は、鋼板形状及び材質を均一化する観点からも好ましい。潤滑圧延する際の摩擦係数は、０．１０以上０．２５以下の範囲が好ましい。 To reduce the rolling load, some or all of the finishing rolling may be performed as lubricated rolling. Lubricated rolling is also preferred from the viewpoint of making the steel sheet shape and material uniform. The friction coefficient during lubricated rolling is preferably in the range of 0.10 to 0.25.
（巻取り）
　仕上げ圧延後に巻き取りを行い、熱延板が得られる。熱間圧延後の巻取温度は、後述する冷間圧延及び焼鈍に際しての通板性を良好にする観点から、３００℃以上７００℃以下が好ましい。 (Winding)
The coiling temperature after the hot rolling is preferably 300° C. or more and 700° C. or less from the viewpoint of improving the sheet passing property during the cold rolling and annealing described later.
《酸洗及び冷間圧延工程》
　熱間圧延により得られた熱延板を酸洗する。酸洗により、熱延板の表面の酸化物が除去されて、最終製品である高強度鋼板において、優れた化成処理性及びめっき層の品質等を得ることができる。酸洗は、一回でもよいし、複数回に分けても実施してもよい。 Pickling and cold rolling process
The hot-rolled sheet obtained by hot rolling is pickled. By pickling, oxides on the surface of the hot-rolled sheet are removed, and excellent chemical conversion treatability and quality of the plating layer can be obtained in the final product, a high-strength steel sheet. Pickling may be performed once or multiple times.
　酸洗後の熱延板に、任意で軟質化熱処理を施してから、冷間圧延を施す。こうして、冷延板が得られる。冷間圧延の条件は、特に限定されないが、冷間圧延の総圧下率は、２０％以上７５％以下が好ましい。圧延パスの回数及び各パスの圧下率は、特に限定されない。 After pickling, the hot-rolled sheet is optionally subjected to a softening heat treatment and then cold-rolled. In this way, a cold-rolled sheet is obtained. There are no particular limitations on the cold-rolling conditions, but the total reduction ratio of the cold rolling is preferably 20% or more and 75% or less. There are no particular limitations on the number of rolling passes and the reduction ratio of each pass.
《第一めっき工程（任意）》
　本発明の一実施形態においては、熱間圧延工程後（冷間圧延を施す場合は、冷間圧延工程後）、かつ焼鈍工程の前の鋼板の片面又は両面において、金属めっき層である第一めっき層を形成する第一めっき工程を含んでいてもよい。第一めっき工程は、金属電気めっき工程であることが好ましい。 First plating process (optional)
In one embodiment of the present invention, a first plating step may be included in which a first plating layer, which is a metal plating layer, is formed on one or both sides of the steel sheet after the hot rolling step (after the cold rolling step if cold rolling is performed) and before the annealing step. The first plating step is preferably a metal electroplating step.
　例えば、上記のようにして得られた冷延板の表面に金属電気めっき処理等の金属めっき処理（第一めっき処理）を施して、焼鈍前金属めっき層が少なくとも片面に形成された焼鈍前金属めっき鋼板としてもよい。なお、ここでいう金属めっき層は、上記第一めっき層であることができる。焼鈍前金属めっき鋼板は、焼鈍前金属電気めっき層を備えた焼鈍前金属電気めっき鋼板であることが好ましい。 For example, a metal plating process (first plating process) such as metal electroplating may be applied to the surface of the cold-rolled sheet obtained as described above to produce a pre-annealed metal-plated steel sheet having a pre-annealed metal plating layer formed on at least one side. The metal plating layer referred to here may be the above-mentioned first plating layer. The pre-annealed metal-plated steel sheet is preferably a pre-annealed metal electroplated steel sheet provided with a pre-annealed metal electroplating layer.
　金属電気めっき処理方法は特に限定されないが、前述したように素地鋼板上に形成させる金属めっき層としては、金属電気めっき層とすることが好ましいため、金属電気めっき処理を施すことが好ましい。例えば、Ｆｅ系電気めっき浴では硫酸浴、塩酸浴あるいは両者の混合などが適用できる。また、焼鈍前金属電気めっき層の付着量は、通電時間等によって調整することができる。なお、焼鈍前金属電気めっき鋼板とは、金属電気めっき層が焼鈍工程を経ていないことを意味し、金属電気めっき処理前の熱延板、熱延後酸洗処理板又は冷延板について予め焼鈍された態様を除外するものではない。 The metal electroplating method is not particularly limited, but as described above, it is preferable to form a metal electroplating layer on the base steel sheet, and therefore it is preferable to perform a metal electroplating process. For example, a sulfuric acid bath, a hydrochloric acid bath, or a mixture of both can be used for an Fe-based electroplating bath. The amount of adhesion of the metal electroplating layer before annealing can be adjusted by the current application time, etc. Note that the metal electroplated steel sheet before annealing means that the metal electroplating layer has not been subjected to an annealing process, and does not exclude the case where the hot-rolled sheet before the metal electroplating process, the pickled sheet after hot rolling, or the cold-rolled sheet has been annealed in advance.
　ここで、電気めっき層の金属種としては、Ｃｒ、Ｍｎ、Ｆｅ、Ｃｏ、Ｎｉ、Ｃｕ、Ｇａ、Ｇｅ、Ａｓ、Ｒｕ、Ｒｈ、Ｐｄ、Ａｇ、Ｃｄ、Ｉｎ、Ｓｎ、Ｓｂ、Ｏｓ、Ｉｒ、Ｒｔ、Ａｕ、Ｈｇ、Ｔｉ、Ｐｂ、Ｂｉのいずれでもかまわないが、Ｆｅであることがより好ましいため、Ｆｅ系電気めっきの製造方法を以下に述べる。 The metal species of the electroplating layer can be any of Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi, but since Fe is more preferable, the manufacturing method of Fe-based electroplating is described below.
　通電開始前のＦｅ系電気めっき浴中のＦｅイオン含有量は、Ｆｅ２＋として０．５ｍｏｌ／Ｌ以上とすることが好ましい。Ｆｅ系電気めっき浴中のＦｅイオン含有量が、Ｆｅ２＋として０．５ｍｏｌ／Ｌ以上であれば、十分なＦｅ付着量を得ることができる。また、十分なＦｅ付着量を得るために、通電開始前のＦｅ系電気めっき浴中のＦｅイオン含有量は、２．０ｍｏｌ／Ｌ以下とすることが好ましい。
　また、Ｆｅ系電気めっき浴中にはＦｅイオンに加えて、Ｂ、Ｃ、Ｐ、Ｎ、Ｏ、Ｎｉ、Ｍｎ、Ｍｏ、Ｚｎ、Ｗ、Ｐｂ、Ｓｎ、Ｃｒ、Ｖ及びＣｏからなる群から選ばれる少なくとも一種の元素を含有することができる。Ｆｅ系電気めっき浴中でのこれらの元素の合計含有量は、焼鈍前Ｆｅ系電気めっき層中でこれらの元素の合計含有量が１０質量％以下となるようにすることが好ましい。なお、金属元素は金属イオンとして含有すればよく、非金属元素はホウ酸、リン酸、硝酸、有機酸等の一部として含有することができる。また、硫酸鉄めっき液中には、硫酸ナトリウム、硫酸カリウム等の伝導度補助剤や、キレート剤、ｐＨ緩衝剤が含まれていてもよい。 The Fe ion content in the Fe-based electroplating bath before the start of energization is preferably 0.5 mol/L or more in terms of Fe 2+ . If the Fe ion content in the Fe-based electroplating bath is 0.5 mol/L or more in terms of Fe 2+ , a sufficient Fe deposition amount can be obtained. In addition, in order to obtain a sufficient Fe deposition amount, the Fe ion content in the Fe-based electroplating bath before the start of energization is preferably 2.0 mol/L or less.
The Fe-based electroplating bath may contain at least one element selected from the group consisting of B, C, P, N, O, Ni, Mn, Mo, Zn, W, Pb, Sn, Cr, V, and Co in addition to Fe ions. The total content of these elements in the Fe-based electroplating bath is preferably such that the total content of these elements in the Fe-based electroplating layer before annealing is 10 mass% or less. The metal elements may be contained as metal ions, and the nonmetal elements may be contained as part of boric acid, phosphoric acid, nitric acid, organic acid, etc. The iron sulfate plating solution may also contain a conductivity aid such as sodium sulfate or potassium sulfate, a chelating agent, or a pH buffer.
　Ｆｅ系電気めっき浴のその他の条件は、特に限定されない。Ｆｅ系電気めっき液の温度は、定温保持性の観点から、３０℃以上とすることが好ましく、また、８５℃以下が好ましい。Ｆｅ系電気めっき浴のｐＨも、特に限定されないが、水素発生による電流効率の低下を防ぐ観点から１．０以上とすることが好ましく、また、Ｆｅ系電気めっき浴の電気伝導度の観点から３．０以下が好ましい。電流密度は、生産性の観点から１０Ａ／ｄｍ２以上とすることが好ましく、また、Ｆｅ系電気めっき層の付着量制御を容易にする観点から１５０Ａ／ｄｍ２以下であることが好ましい。通板速度は、生産性の観点から５ｍｐｍ以上とすることが好ましく、また、付着量を安定的に制御する観点から１５０ｍｐｍ以下とすることが好ましい。 Other conditions of the Fe-based electroplating bath are not particularly limited. The temperature of the Fe-based electroplating solution is preferably 30° C. or higher from the viewpoint of constant temperature retention, and is preferably 85° C. or lower. The pH of the Fe-based electroplating bath is also not particularly limited, but is preferably 1.0 or higher from the viewpoint of preventing a decrease in current efficiency due to hydrogen generation, and is preferably 3.0 or lower from the viewpoint of the electrical conductivity of the Fe-based electroplating bath. The current density is preferably 10 A/dm 2 or higher from the viewpoint of productivity, and is preferably 150 A/dm 2 or lower from the viewpoint of facilitating control of the deposition amount of the Fe-based electroplating layer. The sheet passing speed is preferably 5 mpm or higher from the viewpoint of productivity, and is preferably 150 mpm or lower from the viewpoint of stably controlling the deposition amount.
　Ｆｅ系電気めっき処理を施す前の処理として、冷延板表面を清浄化するための脱脂処理及び水洗、さらには、冷延板表面を活性化するための酸洗処理及び水洗を施すことができる。これらの前処理に引き続いてＦｅ系電気めっき処理を実施する。脱脂処理及び水洗の方法は特に限定されず、通常の方法を用いることができる。
　酸洗処理においては、硫酸、塩酸、硝酸及びこれらの混合物等各種の酸が使用できる。中でも、硫酸、塩酸及びこれらの混合物が好ましい。酸の濃度は特に限定されないが、酸化皮膜の除去能力及び過酸洗による肌荒れ（表面欠陥）防止等の観点から、１質量％以上２０質量％以下が好ましい。
　また、酸洗処理液には、消泡剤、酸洗促進剤、酸洗抑制剤等を含有してもよい。 As a treatment before the Fe-based electroplating treatment, a degreasing treatment and water washing for cleaning the surface of the cold-rolled sheet, and further, a pickling treatment and water washing for activating the surface of the cold-rolled sheet can be performed. The Fe-based electroplating treatment is performed following these pretreatments. The method of the degreasing treatment and water washing is not particularly limited, and a conventional method can be used.
In the pickling treatment, various acids such as sulfuric acid, hydrochloric acid, nitric acid, and mixtures thereof can be used. Among them, sulfuric acid, hydrochloric acid, and mixtures thereof are preferred. The concentration of the acid is not particularly limited, but is preferably 1% by mass or more and 20% by mass or less from the viewpoints of the ability to remove the oxide film and the prevention of roughness (surface defects) due to over-pickling.
The pickling solution may also contain an antifoaming agent, a pickling promoter, a pickling inhibitor, and the like.
《第１加熱工程》
　次いで、得られた冷延板を７５０℃以上で第１加熱を施す。冷延板は、電気めっき処理されたものであってもよいが、当該処理を施されていなくてもよい。 <<First heating step>>
Next, the obtained cold-rolled sheet is subjected to a first heating at a temperature of 750° C. or more. The cold-rolled sheet may be one that has been subjected to an electroplating treatment, but may not be one that has been subjected to said treatment.
　第１加熱温度が低すぎると、オーステナイトへの逆変態が十分進行せず、マルテンサイト分の面積率が低くなり、所望のＴＳが得られない。このため、第１加熱温度は、７５０℃以上であり、７７０℃以上が好ましい。
　加熱温度の上限は、特に限定されないが、操業性等の観点から、９５０℃以下が好ましい。 If the first heating temperature is too low, the reverse transformation to austenite does not proceed sufficiently, the area ratio of martensite is reduced, and the desired TS cannot be obtained. Therefore, the first heating temperature is 750° C. or higher, and preferably 770° C. or higher.
The upper limit of the heating temperature is not particularly limited, but is preferably 950° C. or less from the viewpoint of operability and the like.
　冷延板を第１加熱温度で加熱する時間（加熱時間）は、特に限定されないが、短すぎると、オーステナイトへの逆変態が十分進行しないおそれがあることから、３０ｓ以上が好ましく、６０ｓ以上がより好ましい。
　加熱時間の上限は、特に限定されず、例えば、６０００ｓ以下とすることができ、３０００ｓ以下が好ましい。ここで、「ｓ」は、秒を意味する。 The time for which the cold-rolled sheet is heated at the first heating temperature (heating time) is not particularly limited, but if the time is too short, the reverse transformation to austenite may not proceed sufficiently, so the time is preferably 30 seconds or more, and more preferably 60 seconds or more.
The upper limit of the heating time is not particularly limited, and can be, for example, 6000 s or less, and preferably 3000 s or less, where "s" means seconds.
　第１加熱の焼鈍雰囲気の露点は－３０℃以上とすることが好ましい。焼鈍工程における焼鈍雰囲気の露点を－３０℃以上で行うことで、脱炭反応が促進され、表層軟質層がより深く形成される。これにより、素地鋼板表面から表層軟質層の板厚方向深さの１／４位置の板面の５０μｍ×５０μｍの領域において、３００点以上のナノ硬度を測定したとき、ナノ硬度が７．０ＧＰａ以上の割合が０．１０以下となる。焼鈍工程の焼鈍雰囲気は、より好ましくは－１５℃以上、さらにより好ましくは－５℃以上である。焼鈍工程の焼鈍雰囲気の露点の上限は特に限定されないが、Ｆｅ系電気めっき層表面の酸化を好適に防ぎ、亜鉛めっき層を設ける際のめっき密着性を良好にする観点から、焼鈍工程の焼鈍雰囲気の露点は３０℃以下とすることが好ましい。 The dew point of the annealing atmosphere in the first heating is preferably -30°C or higher. By performing the annealing process at a dew point of -30°C or higher, the decarburization reaction is promoted and the soft surface layer is formed deeper. As a result, when the nano hardness is measured at 300 or more points in a 50 μm x 50 μm area on the sheet surface at a position 1/4 of the depth in the sheet thickness direction of the soft surface layer from the surface of the base steel sheet, the proportion of nano hardness of 7.0 GPa or higher is 0.10 or less. The annealing atmosphere in the annealing process is more preferably -15°C or higher, and even more preferably -5°C or higher. There is no particular upper limit to the dew point of the annealing atmosphere in the annealing process, but from the viewpoint of suitably preventing oxidation of the surface of the Fe-based electroplating layer and improving plating adhesion when a zinc plating layer is provided, the dew point of the annealing atmosphere in the annealing process is preferably 30°C or lower.
《第１冷却工程》
　次いで加熱した冷延板を冷却するが、その際、Ｔ２以上７５０℃以下の温度域Ｔ１を経る。冷延板を、温度域Ｔ１において、以下に説明する第１平均冷却速度ｖ１で冷却することにより、パーライト変態が抑制され、続く滞留工程においてベイナイト変態を活用することができる。 <<First cooling step>>
The heated cold-rolled sheet is then cooled, passing through a temperature region T1 from T2 to 750° C. In the temperature region T1 , the cold-rolled sheet is cooled at a first average cooling rate v1 described below, whereby pearlite transformation is suppressed and bainite transformation can be utilized in the subsequent retention step.
　第１平均冷却速度ｖ１が低すぎると、第１加熱で生成したオーステナイトにおいてパーライト変態が生じ、未変態オーステナイトが減少し、続く滞留工程においてベイナイト変態が活用できずベイナイトの面積率が減少し、低温靭性及び伸びフランジ性が劣化する。このため、第１平均冷却速度ｖ１は、２．０℃／ｓ以上であり、３．０℃／ｓ以上が好ましく、５．０℃／ｓ以上がより好ましい。
　第１平均冷却速度ｖ１の上限は特に限定されないが、設備投資負担の軽減の観点から、６０．０℃／ｓ以下が好ましい。
　温度域Ｔ１での冷却は、連続冷却が好ましい。
　第１加熱温度から７５０℃までの冷却速度は特に限定されない。 If the first average cooling rate v1 is too low, pearlite transformation occurs in the austenite generated in the first heating, untransformed austenite decreases, and bainite transformation cannot be utilized in the subsequent retention step, resulting in a decrease in the area ratio of bainite and deterioration of low-temperature toughness and stretch flangeability. Therefore, the first average cooling rate v1 is 2.0° C./s or more, preferably 3.0° C./s or more, and more preferably 5.0° C./s or more.
The upper limit of the first average cooling rate v1 is not particularly limited, but from the viewpoint of reducing the capital investment burden, it is preferably 60.0° C./s or less.
The cooling in the temperature region T1 is preferably continuous cooling.
The cooling rate from the first heating temperature to 750° C. is not particularly limited.
《滞炉工程》
　温度域Ｔ１を経た冷延板は、次いで、３５０℃以上５５０℃以下の滞留温度Ｔ２で保持する滞炉工程を行う。第１冷却工程後の冷延板を、３５０℃以上５５０℃以下の滞留温度Ｔ２で、式１で定義されるＦが０．２０以上０．９０以下を満たす滞留時間ｔ２（ｓ）の条件で滞炉することにより、ベイナイト変態が生じる。この滞留時間ｔ２（ｓ）をフォーマスター試験から得られる膨張曲線から求め、適用することで、ベイナイト変態とそれに伴う鉄炭化物の析出や未変態オーステナイト中へのＣの分配が適正化する。膨張曲線は鋼成分及び第１冷却までの熱履歴に依存するため、鋼成分及び第１加熱温度からＴ２℃になるまでの熱履歴ごとに膨張曲線を求め、適切な滞留時間ｔ２を選択する必要がある。 "Kiln-Stay Process"
The cold-rolled sheet that has passed through the temperature range T1 is then subjected to a residence time step of holding at a residence temperature T2 of 350°C or more and 550°C or less. The cold-rolled sheet that has passed through the first cooling step is then held at a residence temperature T2 of 350°C or more and 550°C or less for a residence time t2 (s) that satisfies F defined by formula 1 of 0.20 or more and 0.90 or less, thereby causing bainite transformation. By determining and applying this residence time t2 (s) from an expansion curve obtained from a Formaster test, the bainite transformation and the associated precipitation of iron carbides and the distribution of C into untransformed austenite are optimized. Since the expansion curve depends on the steel composition and the thermal history up to the first cooling, it is necessary to determine an expansion curve for each steel composition and thermal history from the first heating temperature to T2 °C, and select an appropriate residence time t2 .
　ここで、式１は以下のとおりである。
式１：Ｆ＝１－ｅｘｐ（－ｋｔｎ）
ｔ：滞留時間（ｓ）
ｋ、ｎ：フォーマスター試験の膨張曲線から求められる定数 Here, Equation 1 is as follows:
Formula 1: F=1-exp(-kt n )
t: residence time (s)
k, n: constants obtained from the expansion curve of the Formaster test
　冷延板が滞留に滞留する時間（滞留時間ｔ２）が短すぎると、ベイナイト変態が不十分となり、ベイナイトの面積率が低くなり、またナノ硬度の標準偏差が大きくなり、低温靭性、伸びフランジ性、剪断端面部の曲げ性が劣化する。このため、滞留時間ｔ２はＦが０．２０となる時間ｔ以上であり、好ましくはＦが０．３０となる時間ｔ以上である。
　一方、冷延板が滞留温度Ｔ２に滞留する時間（滞留時間ｔ２）が長すぎると、ベイナイト変態が過度に進行し、マルテンサイト量が減少し、ＴＳが低下する。また、未変態オーステナイト中にセメンタイトが析出し、続く第２加熱工程で準安定炭化物が析出しにくくなり、マルテンサイトブロック数に対する準安定炭化物が存在するマルテンサイトブロック数の割合が低下し、低温靭性が劣化する。このため、滞留時間ｔ２はＦが０．９０となる時間ｔ以下であり、好ましくはＦが０．８０となる時間ｔ以下である。 If the time that the cold-rolled sheet is held in the holding tank (holding time t2 ) is too short, the bainite transformation becomes insufficient, the area ratio of bainite becomes low, the standard deviation of nano-hardness becomes large, and low-temperature toughness, stretch flangeability, and bendability of the sheared edge portion deteriorate. Therefore, the holding time t2 is the time t or more at which F becomes 0.20, and preferably the time t or more at which F becomes 0.30.
On the other hand, if the time during which the cold-rolled sheet is held at the holding temperature T2 (holding time t2 ) is too long, the bainite transformation proceeds excessively, the amount of martensite decreases, and TS decreases. In addition, cementite precipitates in the untransformed austenite, making it difficult for metastable carbides to precipitate in the subsequent second heating step, decreasing the ratio of the number of martensite blocks containing metastable carbides to the number of martensite blocks, and deteriorating low-temperature toughness. For this reason, the holding time t2 is equal to or shorter than the time t at which F becomes 0.90, and preferably equal to or shorter than the time t at which F becomes 0.80.
　式１におけるＦとｔの関係は次のように求める。
　上記鋼スラブを第１冷却までの工程に付し、次いで３５０℃以上５５０℃以下の滞留温度Ｔ２で滞留させる。第１冷却までの工程は、熱間圧延工程、酸洗及び冷間圧延工程、第１加熱工程及び第１冷却までの工程である。
　フォーマスター試験機を用いて、滞留温度Ｔ２での滞留中の膨張曲線を取得する。Ｔ２での滞留は膨張が停留するまで実施する。滞留温度Ｔ２での滞炉開始時の膨張量を０、停留時の膨張量を１とし、膨張曲線を式１でフィッティングを行い、定数ｋとｎを算出する。これにより、滞留温度Ｔ２におけるＦとｔの関係が求まる。
　式１：Ｆ＝１－ｅｘｐ（－ｋｔｎ）
ｔ：滞留時間（ｓ）
ｋ、ｎ：フォーマスター試験の膨張曲線から求められる定数 The relationship between F and t in formula 1 is calculated as follows.
The steel slab is subjected to a process up to the first cooling, and then is retained at a retention temperature T2 of 350° C. to 550° C. The process up to the first cooling includes a hot rolling process, a pickling and cold rolling process, a first heating process, and a process up to the first cooling.
Using a Formaster tester, an expansion curve during retention at retention temperature T2 is obtained. Retention at T2 is continued until expansion stops. The expansion amount at the start of retention at retention temperature T2 is set to 0, and the expansion amount at retention is set to 1. The expansion curve is fitted with Equation 1 to calculate constants k and n. This determines the relationship between F and t at retention temperature T2 .
Formula 1: F=1-exp(-kt n )
t: residence time (s)
k, n: constants obtained from the expansion curve of the Formaster test
《第２冷却工程》
　次いで、滞留工程を経た冷延板を冷却する。冷却停止温度は、Ｍｓ－２０℃以下である。これにより、マルテンサイト変態が十分に進行する。冷却停止温度がＭｓ－２０℃超である場合、未変態オーステナイトがマルテンサイト変態せずに、残留オーステナイト量が過大となり、良好な部品強度及び伸びフランジ性が得られない。冷却停止温度は室温でもよい。ここで、Ｍｓは、マルテンサイト変態が起こり始める温度（Ｍｓ点）であり、後述する試験により、測定した値を用いる。 <<Second cooling step>>
Next, the cold-rolled sheet that has been through the retention step is cooled. The cooling stop temperature is Ms-20°C or less. This allows the martensitic transformation to proceed sufficiently. If the cooling stop temperature exceeds Ms-20°C, the untransformed austenite does not transform to martensite, the amount of retained austenite becomes excessive, and good part strength and stretch flangeability cannot be obtained. The cooling stop temperature may be room temperature. Here, Ms is the temperature (Ms point) at which martensitic transformation begins to occur, and a value measured by the test described below is used.
　冷却において、滞留工程を経た冷延板は、Ｍｓ－２０℃以上Ｍｓ℃以下の温度域Ｔ３を経るが、冷延板を、温度域Ｔ３において、以下に説明する第２平均冷却速度ｖ２で冷却することにより、マルテンサイト中でのセメンタイト析出を抑制できる。 In the cooling, the cold-rolled sheet that has undergone the retention step passes through a temperature region T3 of Ms-20 ° C. or more and Ms ° C. or less. By cooling the cold-rolled sheet at a second average cooling rate v2 described below in the temperature region T3 , cementite precipitation in martensite can be suppressed.
　第２平均冷却速度ｖ２が低すぎると、マルテンサイト中にセメンタイトが析出し、続く第２加熱で準安定炭化物の析出が抑制され、マルテンサイトブロック数に対する準安定炭化物が存在するマルテンサイトブロック数の割合が低下し、低温靭性が劣化する。
　このため、第２平均冷却速度ｖ２は、５℃／ｓ以上であり、８℃／ｓ以上が好ましい。第２平均冷却速度ｖ２の上限は特に限定されないが、設備投資負担の軽減の観点から、６０．０℃／ｓ以下が好ましい。
　温度域Ｔ３以外の冷却速度は特に限定されない。 If the second average cooling rate v2 is too low, cementite precipitates in martensite, and the precipitation of metastable carbides is suppressed in the subsequent second heating, so that the ratio of the number of martensite blocks containing metastable carbides to the number of martensite blocks decreases, and low-temperature toughness deteriorates.
Therefore, the second average cooling rate v2 is 5° C./s or more, and preferably 8° C./s or more. The upper limit of the second average cooling rate v2 is not particularly limited, but from the viewpoint of reducing the capital investment burden, it is preferably 60.0° C./s or less.
The cooling rate outside the temperature region T3 is not particularly limited.
　Ｍｓ点は、次のようにフォーマスター試験により測定した値を用いる。
フォーマスター試験機を用いて、上記鋼スラブを滞留工程終了までの工程に付し、次いで第２平均冷却速度５℃／ｓ以上で室温まで冷却する。第２冷却中にマルテンサイト変態が起こり膨張が始める温度をＭｓ点とする。第２平均冷却速度の上限は、特に限定されないが、例えば１００℃／ｓ以下とすることができる。 The Ms point is a value measured by a Formaster test as follows.
Using a Formaster testing machine, the steel slab is subjected to a process up to the end of the retention process, and then cooled to room temperature at a second average cooling rate of 5°C/s or more. The temperature at which martensitic transformation occurs and expansion begins during the second cooling is defined as the Ms point. The upper limit of the second average cooling rate is not particularly limited, but can be, for example, 100°C/s or less.
《第２加熱工程》
　冷却停止温度まで冷却した冷延板に、次いで、第２加熱を施す。
　第２加熱を実施することにより、上述した冷却中に生成したマルテンサイトブロックにおいて、低温靭性を向上させる準安定炭化物が析出する。これにより、良好な部品強度、伸びフランジ性及び剪断端面部曲げ性を有しつつ、低温靭性に優れる、ＴＳが７８０ＭＰａ以上の高強度鋼板が得られる。 <<Second heating step>>
The cold-rolled sheet cooled to the cooling stop temperature is then subjected to a second heating.
By carrying out the second heating, metastable carbides that improve low-temperature toughness are precipitated in the martensite blocks formed during the above-mentioned cooling, and as a result, a high-strength steel plate having a TS of 780 MPa or more and excellent low-temperature toughness while having good part strength, stretch flangeability, and shear end face bendability can be obtained.
　冷却を施した冷延板に、第２加熱を第２加熱温度Ｘ（単位：℃）まで加熱し、保持時間Ｙ（単位：ｓ）で保持する。保持時間Ｙにおける温度は、Ｘ±２０℃の範囲内とする。
　ここで、ＸとＹは、下記式２を満たすこととする。
　式２：７０００≦（２７３＋Ｘ）×（２０＋ｌｏｇ（Ｙ／３６００））≦１３０００ The cooled cold-rolled sheet is then subjected to a second heating to a second heating temperature X (unit: ° C.) and held for a holding time Y (unit: s). The temperature during the holding time Y is within the range of X±20° C.
Here, X and Y satisfy the following formula 2.
Formula 2: 7000≦(273+X)×(20+log(Y/3600))≦13000
　上記式２中の「（２７３＋Ｘ）×（２０＋ｌｏｇ（Ｙ／３６００））」を、以下、便宜的に、「変数部Ｚ」と呼ぶ。 In the above formula 2, "(273+X) x (20+log(Y/3600))" will be referred to as the "variable part Z" for the sake of convenience.
　上記の条件を満たす第２加熱を行うことにより、マルテンサイトブロック数に対する準安定炭化物が存在するマルテンサイトブロック数の割合（割合ｐ）を高くし、低温靭性の向上を図ることができる。
　ここで、温度Ｘ（℃）は室温より高い値とする。
　第２冷却工程の冷却停止温度より第２加熱の温度Ｘ（℃）が高い場合、冷却停止温度から第２加熱の温度Ｘ（℃）まで加熱を行う。
　第２冷却工程の冷却停止温度と第２加熱の温度Ｘ（℃）が同じであってもよく、この場合、第２加熱の温度Ｘ（℃）である冷却停止温度で保持することを意味する。 By carrying out the second heating that satisfies the above conditions, the ratio (ratio p) of the number of martensite blocks in which metastable carbides exist to the total number of martensite blocks can be increased, thereby improving low temperature toughness.
Here, the temperature X (° C.) is a value higher than room temperature.
When the temperature X (° C.) of the second heating is higher than the cooling stop temperature of the second cooling step, heating is performed from the cooling stop temperature to the temperature X (° C.) of the second heating.
The cooling stop temperature of the second cooling step and the temperature X (°C) of the second heating may be the same, in which case it means that the cooling stop temperature is maintained at the temperature X (°C) of the second heating.
　上記式２の変数部Ｚの値が小さすぎる場合、すなわち、温度Ｘが低すぎるか、かつ／又は、保持時間Ｙが短すぎる場合、準安定炭化物が十分に析出しないため、割合ｐが低くなる。このため、割合ｐが高くなるという観点から、変数部Ｚの値は、７０００以上であり、８０００以上が好ましい。
　一方、変数部Ｚの値が高すぎる場合、すなわち、温度Ｘが高すぎるか、かつ／又は、保持時間Ｙが長すぎる場合、準安定炭化物のセメンタイトへの遷移が起こり、割合ｐが低くなる。このため、割合ｐが高くなるという観点から、変数部Ｚの値は、１３０００以下であり、１２０００以下が好ましい。 If the value of the variable part Z in the above formula 2 is too small, that is, if the temperature X is too low and/or the holding time Y is too short, metastable carbides are not sufficiently precipitated, and the proportion p becomes low. Therefore, from the viewpoint of increasing the proportion p, the value of the variable part Z is 7000 or more, and preferably 8000 or more.
On the other hand, if the value of the variable part Z is too high, i.e., if the temperature X is too high and/or the holding time Y is too long, the metastable carbides will transition to cementite and the proportion p will decrease. For this reason, from the viewpoint of increasing the proportion p, the value of the variable part Z is 13000 or less, and preferably 12000 or less.
　第２加熱において、温度Ｘ（単位：℃）は、下記式３を満たすことが好ましい。これにより、準安定炭化物が存在するマルテンサイトブロックにおける準安定炭化物の数密度（数密度ｎ）が高くなる。
　式３：１００≦Ｘ≦４００ In the second heating, the temperature X (unit: ° C.) preferably satisfies the following formula 3. This increases the number density (number density n) of metastable carbides in the martensite block in which the metastable carbides are present.
Formula 3: 100≦X≦400
　数密度ｎが高くなるという観点から、温度Ｘは、１００℃以上が好ましく、１２０℃以上がより好ましく、１５０℃以上がさらに好ましい。
　一方、同様の観点から、温度Ｘは、４００℃以下が好ましく、３８０℃以下がより好ましく、３５０℃以下がさらに好ましい。 From the viewpoint of increasing the number density n, the temperature X is preferably 100° C. or higher, more preferably 120° C. or higher, and even more preferably 150° C. or higher.
On the other hand, from the same viewpoint, the temperature X is preferably 400° C. or less, more preferably 380° C. or less, and further preferably 350° C. or less.
　第２加熱を施した冷延板を、その後、例えば、室温まで冷却する。その際の冷却速度は特に限定されない。
　こうして、本発明の製造方法により、本発明の高強度鋼板が得られる。
　本発明の製造方法において、後述するめっき処理を施す場合、得られる本発明の高強度鋼板は、めっき層を有するめっき鋼板である。 The cold-rolled sheet that has been subjected to the second heating is then cooled, for example, to room temperature. The cooling rate at which the sheet is cooled is not particularly limited.
Thus, the high strength steel plate of the present invention can be obtained by the manufacturing method of the present invention.
In the production method of the present invention, when a plating treatment described later is carried out, the obtained high-strength steel sheet of the present invention is a plated steel sheet having a plating layer.
　本発明の製造方法における一連の熱処理は、上述した熱履歴を満足すれば、その他の条件は特に限定されず、熱処理を実施する設備等も特に限定されない。 As long as the series of heat treatments in the manufacturing method of the present invention satisfies the thermal history described above, there are no other particular limitations on the conditions, and there are no particular limitations on the equipment used to carry out the heat treatments.
（第二めっき工程（任意））
　本発明の一実施形態においては、前記第１加熱から第２加熱工程の鋼板にめっき処理を施し、金属めっき層である第二めっき層を形成する第二めっき工程を含んでいてもよい。
　第二めっき工程は亜鉛めっき工程又は溶融アルミニウムめっき処理とすることが好ましく、亜鉛めっき工程における亜鉛めっき処理としては、例えば、溶融亜鉛めっき処理、合金化亜鉛めっき処理、電気亜鉛めっき処理、アルミニウムめっき処理が挙げられる。 (Second plating process (optional))
In one embodiment of the present invention, the method may include a second plating step of plating the steel sheet from the first heating step to the second heating step to form a second plating layer which is a metal plating layer.
The second plating step is preferably a zinc plating step or a hot-dip aluminum plating process. Examples of the zinc plating process in the zinc plating step include a hot-dip zinc plating process, a zinc alloy plating process, an electrolytic zinc plating process, and an aluminum plating process.
　溶融亜鉛めっき処理の場合、鋼板を４４０℃以上５００℃以下の亜鉛めっき浴中に浸漬させた後、ガスワイピング等によって、めっき付着量を調整することが好ましい。溶融亜鉛めっき浴としては、前記した亜鉛めっき層の組成となれば特に限定されるものではないが、例えば、Ａｌ含有量が０．１０質量％以上０．２３質量％以下であり、残部がＺｎ及び不可避的不純物からなる組成のめっき浴を用いることが好ましい。 In the case of hot-dip galvanizing, it is preferable to immerse the steel sheet in a zinc plating bath at 440°C to 500°C, and then adjust the coating weight by gas wiping or the like. There are no particular limitations on the hot-dip galvanizing bath as long as it has the composition of the zinc plating layer described above, but it is preferable to use a plating bath with an Al content of 0.10 mass% to 0.23 mass%, with the balance being Zn and unavoidable impurities.
　合金化亜鉛めっき処理の場合、前記の要領で溶融亜鉛めっき処理を施した後、溶融亜鉛めっき層を有する鋼板（溶融亜鉛めっき鋼板）を４５０℃以上６００℃以下の合金化温度に加熱して合金化処理を施すことが好ましい。合金化温度が４５０℃未満では、Ｚｎ－Ｆｅ合金化速度が遅くなり、合金化が困難となる場合がある。一方、合金化温度が６００℃を超えると、未変態オーステナイトがパーライトへ変態し、ＴＳを５９０ＭＰａ以上とすることが困難となる。なお、合金化温度は、より好ましくは５１０℃以上である。また、合金化温度は、より好ましくは５７０℃以下である。 In the case of alloying galvanizing treatment, after performing hot-dip galvanizing treatment as described above, it is preferable to perform alloying treatment by heating the steel sheet having a hot-dip galvanized layer (hot-dip galvanized steel sheet) to an alloying temperature of 450°C or more and 600°C or less. If the alloying temperature is less than 450°C, the Zn-Fe alloying rate slows down and alloying may become difficult. On the other hand, if the alloying temperature exceeds 600°C, untransformed austenite transforms to pearlite, making it difficult to achieve a TS of 590 MPa or more. The alloying temperature is more preferably 510°C or more. Also, the alloying temperature is more preferably 570°C or less.
　また、溶融亜鉛めっき層を有する鋼板（溶融亜鉛めっき鋼板）（ＧＩ）及び合金化溶融亜鉛めっき層を有する鋼板（合金化溶融亜鉛めっき鋼板）（ＧＡ）のめっき付着量はいずれも、片面あたり２０～８０ｇ／ｍ２とすることが好ましい。なお、めっき付着量は、ガスワイピング等により調節することが可能である。 The coating weight of the steel sheet having a hot-dip galvanized layer (hot-dip galvanized steel sheet) (GI) and the steel sheet having a galvannealed layer (galvannealed hot-dip galvanized steel sheet) (GA) is preferably 20 to 80 g/ m2 per side. The coating weight can be adjusted by gas wiping or the like.
　また、その他金属めっき処理の具体例として挙げられる溶融アルミニウムめっき処理を施すときは、前記冷延板焼鈍を施して得た冷延板を６６０～７３０℃のアルミニウムめっき浴中に浸漬し、溶融アルミニウムめっき処理を施し、その後、ガスワイピング等によって、めっき付着量を調整する。 When applying molten aluminum plating, which is a specific example of other metal plating processes, the cold-rolled sheet obtained by annealing the cold-rolled sheet is immersed in an aluminum plating bath at 660 to 730°C to apply the molten aluminum plating process, and then the coating weight is adjusted by gas wiping or the like.
　さらに、電気亜鉛めっき処理を施すときは、特に限定しないが、皮膜厚が２μｍから１５μｍの範囲になるようにすることが好ましい。 Furthermore, when performing electrolytic zinc plating, although there are no particular limitations, it is preferable that the coating thickness be in the range of 2 μm to 15 μm.
《スキンパス圧延（任意）》
　得られた高強度鋼板にスキンパス圧延を実施してもよい。スキンパス圧延後は、めっき処理後に行ってもよい。
　スキンパス圧延での圧下率は、降伏強さを上昇させる観点から、０．０５％以上が好ましい。圧下率の上限は、特に限定されないが、生産性の観点から１．５０％が好ましい。
　スキンパス圧延は、オンラインで実施してもよいし、オフラインで実施してもよい。
　一度に目的の圧下率のスキンパスを実施してもよいし、数回に分けて実施してもよい。 Skin pass rolling (optional)
The obtained high strength steel sheet may be subjected to skin pass rolling. After the skin pass rolling, the plating treatment may be performed.
The reduction ratio in the skin pass rolling is preferably 0.05% or more from the viewpoint of increasing the yield strength. The upper limit of the reduction ratio is not particularly limited, but is preferably 1.50% from the viewpoint of productivity.
Skin pass rolling may be performed either online or offline.
The skin pass may be performed at a single time to the desired rolling reduction, or may be performed in several steps.
　生産性の観点から、上述した焼鈍及びめっき処理等の一連の処理は、ＣＧＬ（Continuous Galvanizing Line）で実施するのが好ましい。 From the viewpoint of productivity, it is preferable to carry out the above-mentioned series of processes such as annealing and plating in a CGL (Continuous Galvanizing Line).
［部材及びその製造方法］
　本発明の部材及びその製造方法について説明する。
　本発明の部材は、上記した本発明の高強度鋼板を用いてなる部材である。部材は、例えば、本発明の高強度鋼板を、プレス加工等により、目的の形状に成形することにより製造することができる。 [Members and their manufacturing methods]
The member of the present invention and its manufacturing method will now be described.
The member of the present invention is a member made using the high-strength steel plate of the present invention described above. The member can be produced, for example, by forming the high-strength steel plate of the present invention into a desired shape by press working or the like.
　本発明の高強度鋼板は、部品の強度、延性、伸びフランジ性、剪断端面部の曲げ性及び剪断端面部の曲げ性に優れる高強度鋼板である。そのため、本発明の高強度鋼板又は高強度鋼板を用いてなる部材を、例えば、自動車の骨格構造部品又は自動車の補強部品に適用することによって、車体軽量化による燃費向上を図ることができ、産業上の利用価値は極めて大きい。 The high-strength steel plate of the present invention is a high-strength steel plate that is excellent in strength, ductility, stretch flangeability, bendability of the sheared end surface, and bendability of the sheared end surface. Therefore, by applying the high-strength steel plate of the present invention or a member made of the high-strength steel plate to, for example, an automobile frame structural part or an automobile reinforcing part, it is possible to improve fuel efficiency by reducing the weight of the vehicle body, and the industrial value of the steel plate is extremely great.
　以下に、実施例を挙げて本発明を具体的に説明する。ただし、本発明は、以下に説明する実施例に限定されない。 The present invention will be specifically explained below with reference to examples. However, the present invention is not limited to the examples described below.
［試験Ｎｏ．１～５１］
〈鋼板の製造〉
　下記表１に示す成分組成（残部はＦｅ及び不可避的不純物からなる）を有する溶鋼を転炉において溶製し、連続鋳造法にて鋼スラブを得た。
　得られた鋼スラブに熱間圧延を施して、熱延板を得た。具体的には、鋼スラブを、１２５０℃に加熱して表２に示す条件で粗圧延し、次いで、仕上げ圧延温度９００℃で仕上げ圧延を施してから、５００℃の条件で巻き取り、その後、室温まで冷却して、熱延板を得た。
　得られた熱延板に、酸洗を施した後、５００℃の条件で軟質化熱処理を施し、次いで、圧延率５０％の条件で冷間圧延を施した。こうして、板厚１．６ｍｍの冷延板を得た。 [Test No. 1 to 51]
<Steel plate manufacturing>
Molten steel having the composition shown in Table 1 below (the balance being Fe and unavoidable impurities) was produced in a converter, and a steel slab was obtained by continuous casting.
The obtained steel slab was subjected to hot rolling to obtain a hot-rolled sheet. Specifically, the steel slab was heated to 1250°C and roughly rolled under the conditions shown in Table 2, then finished rolled at a finish rolling temperature of 900°C, coiled at 500°C, and then cooled to room temperature to obtain a hot-rolled sheet.
The obtained hot-rolled sheet was pickled, then softened by heat treatment at 500° C., and then cold-rolled at a rolling ratio of 50% to obtain a cold-rolled sheet having a thickness of 1.6 mm.
　得られた冷延板を、下記表２に示す第１加熱温度まで加熱し、２００ｓ保持した。
　次いで、温度域Ｔ１（７５０℃以下Ｔ２以上）の平均冷却速度が表２に示す第１平均冷却速度ｖ１となるように冷却し、引き続き、その冷却速度で滞留温度Ｔ２まで冷却しＴ２にて表２に示す滞留時間ｔ２で滞炉させた。 The obtained cold-rolled sheet was heated to the first heating temperature shown in Table 2 below and held for 200 s.
Next, the sintered material was cooled so that the average cooling rate in the temperature range T1 ( T2 or higher but not exceeding 750°C) became the first average cooling rate v1 shown in Table 2, and then cooled to a residence temperature T2 at that cooling rate and retained at T2 for a residence time t2 shown in Table 2.
　パラメータＦ、滞留時間ｔ２は、下記式１を満たす。
　式１：Ｆ＝１－ｅｘｐ（－ｋｔｎ）
　ｔ：保持時間（ｓ）
　ｋ、ｎ：フォーマスター試験の膨張曲線から求められる定数 The parameter F and the residence time t2 satisfy the following formula 1.
Formula 1: F=1-exp(-kt n )
t: retention time (s)
k, n: constants obtained from the expansion curve of the Formaster test
　上記式１におけるｋ、ｎは、以下のようにして求めた。
　得られた冷延板を、幅１０ｍｍ×長さ３ｍｍに切断し、フォーマスター試験に供した。フォーマスター試験には、富士電波工機株式会社製のＦＴＭ―１００を用いた。各実施例において、表２に示す第１加熱温度まで加熱し、２００ｓ保持した。次いで、温度域Ｔ１（７５０℃以下Ｔ２以上）の平均冷却速度が表２に示す第１平均冷却速度ｖ１となるように冷却し、引き続き、その冷却速度で滞留温度Ｔ２まで冷却し、Ｔ２にて１０００ｓ滞留して膨張曲線を取得した。例えば、実施例１では第１加熱温度８００℃で２００ｓ保持し、次いで平均冷却速度が１８℃／ｓの条件で４８０℃まで冷却し、次いで、４８０℃で１０００ｓ滞留して得られた膨張曲線を式１でフィッティングし、ｋとｎを求めた。 The values of k and n in the above formula 1 were determined as follows.
The obtained cold-rolled sheet was cut into a width of 10 mm x length of 3 mm and subjected to a Formaster test. For the Formaster test, FTM-100 manufactured by Fuji Electric Industrial Co., Ltd. was used. In each example, the sheet was heated to the first heating temperature shown in Table 2 and held for 200 s. Next, the sheet was cooled so that the average cooling rate in the temperature range T 1 (750 ° C. or less and T 2 or more) became the first average cooling rate v 1 shown in Table 2, and then cooled to the residence temperature T 2 at that cooling rate, and held at T 2 for 1000 s to obtain an expansion curve. For example, in Example 1, the sheet was held at the first heating temperature of 800 ° C. for 200 s, then cooled to 480 ° C. under the condition of an average cooling rate of 18 ° C./s, and then held at 480 ° C. for 1000 s. The expansion curve obtained was fitted with Equation 1 to obtain k and n.
　滞留温度Ｔ２（３５０℃以上５５０℃以下）での滞炉に次いで、温度域Ｔ３（Ｍｓ－２０℃以上Ｍｓ℃以下）の平均冷却速度が表２に示す第２平均冷却速度ｖ２となるように、表２の冷却停止温度まで冷却した。 Following residence time at residence temperature T 2 (350°C or higher and 550°C or lower), the material was cooled to the cooling stop temperature in Table 2 so that the average cooling rate in temperature range T 3 (Ms-20°C or higher and Ms°C or lower) became the second average cooling rate v 2 shown in Table 2.
　Ｍｓ点は、以下のようにして求めた。
　得られた冷延板を、幅１０ｍｍ×長さ３ｍｍに切断し、フォーマスター試験に供した。フォーマスター試験には、富士電波工機株式会社製のＦＴＭ―１００を用いた。各実施例において、表２に示す第１加熱温度まで加熱し、２００ｓ保持した。次いで、温度域Ｔ１（７５０℃以下Ｔ２以上）の平均冷却速度が表２に示す第１平均冷却速度ｖ１となるように冷却し、引き続き、その冷却速度で滞留温度Ｔ２まで冷却し、Ｔ２にてｔ秒滞留し、次いで、３０℃／ｓで室温まで最終冷却し、最終冷却中の膨張曲線を作成し、膨張が認められた温度をＭｓ点とした。例えば、実施例１では第１加熱温度８００℃で２００ｓ保持し、次いで、第１平均冷却速度が１８℃／ｓの条件で４８０℃まで冷却し、次いで、４８０℃で３０ｓ滞留し、次いで、３０℃／ｓで室温まで最終冷却して得られた膨張曲線からＭｓ点を求めた。 The Ms point was determined as follows.
The obtained cold-rolled sheet was cut into a width of 10 mm x length of 3 mm and subjected to a Formaster test. For the Formaster test, FTM-100 manufactured by Fuji Electric Industrial Co., Ltd. was used. In each example, the sheet was heated to the first heating temperature shown in Table 2 and held for 200 s. Next, the sheet was cooled so that the average cooling rate in the temperature range T 1 (750 ° C. or less and T 2 or more) became the first average cooling rate v 1 shown in Table 2, and then cooled to the residence temperature T 2 at that cooling rate, held at T 2 for t seconds, and then finally cooled to room temperature at 30 ° C./s, and an expansion curve during final cooling was created, and the temperature at which expansion was observed was taken as the Ms point. For example, in Example 1, the sheet was held at the first heating temperature of 800 ° C. for 200 s, then cooled to 480 ° C. under the condition of a first average cooling rate of 18 ° C./s, then held at 480 ° C. for 30 s, and then finally cooled to room temperature at 30 ° C./s. The Ms point was obtained from the expansion curve obtained.
《めっき処理》
　一部の冷延板に対しては、滞留温度Ｔ２の滞炉後に、溶融亜鉛めっき処理を実施して、両面にめっき層（溶融亜鉛めっき層）を形成した。すなわち、溶融亜鉛めっき鋼板（ＧＩ）を得た。その後、第２平均冷却速度ｖ２にて、冷却停止温度まで冷却した。
　溶融亜鉛めっき処理には、Ａｌ：０．２０質量％を含有し、残部がＺｎ及び不可避的不純物からなる溶融亜鉛めっき浴（浴温：４７０℃）を使用した。
　溶融亜鉛めっき層の片面あたりの付着量は、４５～７２ｇ／ｍ２程度とした。
　形成した溶融亜鉛めっき層の組成は、Ｆｅ：０．１～１．０質量％及びＡｌ：０．２～１．０質量％を含有し、残部がＦｅ及び不可避的不純物からなる組成であった。 Plating
For some of the cold-rolled sheets, after the retention time at the retention temperature T2 , a hot-dip galvanizing process was performed to form a coating layer (hot-dip galvanized layer) on both sides. That is, a hot-dip galvanized steel sheet (GI) was obtained. Thereafter, the sheet was cooled to a cooling stop temperature at a second average cooling rate v2 .
For the hot-dip galvanizing treatment, a hot-dip galvanizing bath (bath temperature: 470° C.) containing 0.20 mass % Al, with the balance being Zn and unavoidable impurities, was used.
The coating weight of the hot-dip galvanized layer per side was set to about 45 to 72 g/ m2 .
The composition of the formed hot-dip galvanized layer contained 0.1 to 1.0 mass % of Fe, 0.2 to 1.0 mass % of Al, and the remainder being Fe and unavoidable impurities.
　別の一部の冷延板に対しては、滞留温度Ｔ２の滞炉後に、合金化溶融亜鉛めっき処理を実施して、両面にめっき層（合金化溶融亜鉛めっき層）を形成した。すなわち、合金化溶融亜鉛めっき鋼板（ＧＡ）を得た。その後、第２平均冷却速度ｖ２にて、冷却停止温度まで冷却した。
　溶融亜鉛めっき処理には、Ａｌ：０．１４質量％を含有し、残部がＺｎ及び不可避的不純物からなる溶融亜鉛めっき浴（浴温：４７０℃）を使用した。
　合金化処理温度は、５５０℃とした。
　合金化溶融亜鉛めっき層の片面あたりの付着量は、４５ｇ／ｍ２程度とした。
　形成した合金化溶融亜鉛めっき層の組成は、Ｆｅ：７～１５質量％及びＡｌ：０．１～１．０質量％を含有し、残部がＦｅ及び不可避的不純物からなる組成であった。 Another part of the cold-rolled sheet was subjected to a galvannealing process after being held in the furnace at a holding temperature T2 to form a plated layer (galvannealed layer) on both sides. That is, a galvannealed steel sheet (GA) was obtained. Thereafter, the steel sheet was cooled to a cooling stop temperature at a second average cooling rate v2 .
For the hot-dip galvanizing treatment, a hot-dip galvanizing bath (bath temperature: 470° C.) containing 0.14 mass % Al, with the balance being Zn and unavoidable impurities, was used.
The alloying treatment temperature was set to 550°C.
The coating weight of the galvannealed layer on one side was about 45 g/ m2 .
The composition of the formed alloyed hot-dip galvanized layer contained 7 to 15 mass % Fe, 0.1 to 1.0 mass % Al, and the balance being Fe and unavoidable impurities.
　溶融亜鉛めっき層を形成した場合は「ＧＩ」、合金化溶融亜鉛めっき層を形成した場合は「ＧＡ」、めっき層を形成しなかった場合は「ＣＲ」を、下記表３の「めっき種類」の欄に記載した。 If a hot-dip galvanized layer was formed, it is indicated as "GI", if an alloyed hot-dip galvanized layer was formed, it is indicated as "GA", and if no plating layer was formed, it is indicated as "CR" in the "Plating type" column in Table 3 below.
　冷却停止点まで冷却した鋼板を、表２に示す温度Ｘ［℃］まで再加熱し、保持時間Ｙ［ｓ］保持した。変数部Ｚは、（２７３＋Ｘ）×（２０＋ｌｏｇ（Ｙ／３６００））である。 The steel plate cooled to the cooling stop point was reheated to the temperature X [℃] shown in Table 2 and held for a holding time of Y [s]. The variable Z is (273 + X) x (20 + log (Y/3600)).
〈鋼組織の観察〉
　得られた鋼板について、上述した方法にしたがって、マルテンサイト、ベイナイト、フェライト及び残留オーステナイトの面積率、マルテンサイトブロック数に対する準安定炭化物が存在するマルテンサイトブロック数の割合（割合ｐ）、準安定炭化物が存在するマルテンサイトブロックにおける準安定炭化物の数密度の平均値（数密度ｎ）、ナノ硬度の標準偏差（測定点２２５点）を測定した。結果を下記表３に示す。
　残部組織についても、一般的な公知の方法により面積率を測定した。残部組織に関して、下記表３中の「θ」は、フェライト中に析出したセメンタイトを意味する。 Observation of steel structure
For the obtained steel sheets, the area ratios of martensite, bainite, ferrite and retained austenite, the ratio of the number of martensite blocks containing metastable carbides to the number of martensite blocks (ratio p), the average number density of metastable carbides in the martensite blocks containing metastable carbides (number density n), and the standard deviation of nanohardness (225 measurement points) were measured according to the above-mentioned methods. The results are shown in Table 3 below.
The area ratio of the remaining structure was also measured by a commonly known method. Regarding the remaining structure, "θ" in Table 3 below means cementite precipitated in ferrite.
〈評価〉
　得られた鋼板について、以下に説明する試験を実施して、各種特性を評価した。結果を下記表３に示す。 <evaluation>
The obtained steel sheets were subjected to the tests described below to evaluate various properties. The results are shown in Table 3 below.
《引張試験》
　引張試験は、ＪＩＳ Ｚ ２２４１に準拠して実施した。
　具体的には、得られた鋼板から、鋼板の圧延方向に対して直角方向が長手方向となるように、ＪＩＳ５号試験片を採取した。採取した試験片を用いて、クロスヘッド速度が１．６７×１０－１ｍｍ／ｓの条件で、引張試験を実施して、降伏強さ（ＹＳ）［ＭＰａ］、引張強さ（ＴＳ）［ＭＰａ］及び全伸び（Ｅｌ）［％］を測定した。さらに、降伏比（ＹＲ）（＝１００×ＹＳ／ＴＳ）［％］を算出した。
　引張強さ（ＴＳ）が７８０ＭＰａ以上である場合、高強度であると判断した。
　降伏比（ＹＲ）が５５％以上である場合、部品強度に優れると判断した。
全伸び（Ｅｌ）が１０％以上である場合、延性に優れると判断した。 Tensile test
The tensile test was carried out in accordance with JIS Z 2241.
Specifically, JIS No. 5 test pieces were taken from the obtained steel plate so that the direction perpendicular to the rolling direction of the steel plate was the longitudinal direction. Using the taken test pieces, a tensile test was carried out under the condition of a crosshead speed of 1.67×10 −1 mm/s to measure the yield strength (YS) [MPa], tensile strength (TS) [MPa] and total elongation (El) [%]. Furthermore, the yield ratio (YR) (= 100×YS/TS) [%] was calculated.
A tensile strength (TS) of 780 MPa or more was determined to be high strength.
A yield ratio (YR) of 55% or more was determined to be excellent in part strength.
When the total elongation (El) was 10% or more, it was determined that the ductility was excellent.
《穴広げ試験》
　穴広げ試験は、ＪＩＳ Ｚ ２２５６に準拠して実施した。
　具体的には、得られた鋼板を剪断して、１００ｍｍ×１００ｍｍのサイズの試験片を採取した。採取した試験片に、クリアランス１２．５％で直径１０ｍｍの穴を打ち抜いた。その後、内径７５ｍｍのダイスを用いて、しわ押さえ力９ｔｏｎ（８８．２６ｋＮ）で抑えた状態で、頂角６０°の円錐ポンチを穴に押し込み、亀裂発生限界における穴直径Ｄｆ［ｍｍ］を測定した。初期の穴直径をＤ０［ｍｍ］として、下記の式から、穴広げ率λ［％］を求めた。
　λ＝｛（Ｄｆ－Ｄ０）／Ｄ０｝×１００
　穴広げ率（λ）が２０％以上である場合、伸びフランジ性に優れると判断した。 <Hole expansion test>
The hole expansion test was carried out in accordance with JIS Z 2256.
Specifically, the obtained steel plate was sheared to obtain a test piece having a size of 100 mm x 100 mm. A hole having a diameter of 10 mm was punched in the obtained test piece with a clearance of 12.5%. Then, a conical punch having an apex angle of 60° was pressed into the hole using a die having an inner diameter of 75 mm while being held down with a wrinkle holding force of 9 ton (88.26 kN), and the hole diameter Df [mm] at the crack initiation limit was measured. The initial hole diameter was D0 [mm], and the hole expansion ratio λ [%] was calculated from the following formula.
λ = {(Df - D0) / D0} x 100
When the hole expansion ratio (λ) was 20% or more, it was determined that the stretch flangeability was excellent.
《剪断端面部曲げ試験》
　曲げ試験は、ＪＩＳ Ｚ ２２４８に準拠して実施した。
　具体的には、得られた鋼板から、鋼板の圧延方向に対して平行方向が曲げ試験の軸方向となるように、幅３０ｍｍ、長さ１００ｍｍの短冊状の試験片を採取した。なお、剪断端面サンプルの曲げ試験では、長手方向の端面を剪断端面とし、研削端面サンプルの曲げ試験では、長手方向の端面を研削端面とした。
　採取した試験片を用いて、押込み荷重１００ｋＮ、押付け保持時間５秒の条件で、９０°Ｖ曲げ試験を実施した。
　適当な曲げ半径Ｒで、５つの試験片について、曲げ試験を実施した。次いで、曲げ頂点の稜線部における亀裂発生の有無を確認した。
　亀裂発生の有無は、曲げ頂点の稜線部を、デジタルマイクロスコープ（ＲＨ－２０００、ハイロックス社製）を用いて、４０倍の倍率で観察することにより、確認した。
　５つの試験片のいずれにも亀裂が発生しなかった最小の曲げ半径Ｒを求め、板厚ｔで割った値（Ｒ／ｔ）を限界曲げ半径とした。剪断端面サンプルの限界曲げ半径（Ｒｓ／ｔ）と研削端面サンプルの限界曲げ半径（Ｒｇ／ｔ）をそれぞれ求め、剪断端面サンプルの限界曲げ半径（Ｒｓ／ｔ）と研削端面サンプルの限界曲げ半径（Ｒｇ／ｔ）の比（Ｒｓ／Ｒｇ）が１．５０以下の場合、剪断端面部の曲げ性に優れると判断した。 <Shear end face bending test>
The bending test was carried out in accordance with JIS Z 2248.
Specifically, a rectangular test piece having a width of 30 mm and a length of 100 mm was taken from the obtained steel sheet so that the axial direction of the bending test was parallel to the rolling direction of the steel sheet. In the bending test of the sheared end face sample, the end face in the longitudinal direction was the sheared end face, and in the bending test of the ground end face sample, the end face in the longitudinal direction was the ground end face.
Using the collected test pieces, a 90° V-bending test was carried out under conditions of a pressing load of 100 kN and a pressing holding time of 5 seconds.
A bending test was carried out on five test pieces at an appropriate bending radius R. Then, the presence or absence of cracks was confirmed at the ridge line of the bending apex.
The occurrence of cracks was confirmed by observing the ridgeline at the apex of bending at 40x magnification using a digital microscope (RH-2000, manufactured by Hirox Corporation).
The minimum bending radius R at which no cracks were generated in any of the five test pieces was determined, and the value (R/t) obtained by dividing the radius by the plate thickness t was taken as the limit bending radius. The limit bending radius (Rs/t) of the sheared end face sample and the limit bending radius (Rg/t) of the ground end face sample were determined, and when the ratio (Rs/Rg) of the limit bending radius (Rs/t) of the sheared end face sample to the limit bending radius (Rg/t) of the ground end face sample was 1.50 or less, the sheared end face was judged to have excellent bendability.
《低温靭性試験》
　得られた鋼板（１．６ｍｍｔ鋼板）を６枚重ね合わせて接着し、９．６ｍｍｔ厚さのシャルピー衝撃試験片を作成した。ノッチは２ｍｍのＵノッチとした。この試験片を用いて、－４０℃でシャルピー衝撃試験を行い、これによって得られるシャルピー吸収エネルギーを測定した。また、試験後の破面を観察することで、延性破面率を測定した。
低温靭性パラメータＰをシャルピー吸収エネルギー（単位：Ｊ）と延性破面率（％）の積（シャルピー吸収エネルギー（Ｊ）×延性破面率（％））と定義し、算出した。
低温靭性パラメータＰが３０００以上であるとき、低温靭性に優れると判断した。 <Low temperature toughness test>
Six of the obtained steel plates (1.6 mmt steel plates) were stacked and bonded together to prepare a Charpy impact test piece with a thickness of 9.6 mmt. The notch was a 2 mm U-notch. Using this test piece, a Charpy impact test was performed at -40°C, and the Charpy absorbed energy obtained was measured. In addition, the ductile fracture rate was measured by observing the fracture surface after the test.
The low-temperature toughness parameter P was defined as the product of Charpy absorbed energy (unit: J) and ductile fracture surface area ratio (%) (Charpy absorbed energy (J) x ductile fracture surface area ratio (%)) and was calculated.
When the low temperature toughness parameter P is 3000 or more, it is determined that the low temperature toughness is excellent.
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000003
Figure JPOXMLDOC01-appb-T000003
　表３に示すように、発明例ではいずれも、ＴＳが７８０ＭＰａ以上であり、部品の強度、延性、伸びフランジ性、剪断端面部の曲げ性及び低温靭性が優れている。一方、比較例では、部品の強度、延性、伸びフランジ性、剪断端面部の曲げ性及び低温靭性のいずれか一つ以上が劣っている。 As shown in Table 3, in all of the invention examples, TS is 780 MPa or more, and the parts have excellent strength, ductility, stretch flangeability, bendability of the sheared end surface, and low-temperature toughness. On the other hand, in the comparative examples, the parts are inferior in one or more of strength, ductility, stretch flangeability, bendability of the sheared end surface, and low-temperature toughness.
［試験Ｎｏ．５２～７５］
表１に示す成分組成を有し、残部がＦｅ及び不可避的不純物からなる鋼スラブ（鋼素材）を転炉にて溶製し、連続鋳造法にて鋼スラブを得た。得られた鋼スラブを１２５０℃に加熱して、粗圧延し、シートバーを得た。次いで、得られたシートバーに、仕上げ圧延温度：９００℃で仕上げ圧延を施し、表４に示す条件で巻取りを行い、熱延板を得た。得られた熱延板に酸洗を施した後、表４に示す条件で冷間圧延を施し、板厚：１．２ｍｍの冷延板を得た。 [Test No. 52-75]
A steel slab (steel material) having the composition shown in Table 1 and the balance consisting of Fe and unavoidable impurities was melted in a converter, and a steel slab was obtained by a continuous casting method. The obtained steel slab was heated to 1250 ° C and roughly rolled to obtain a sheet bar. The obtained sheet bar was then subjected to finish rolling at a finish rolling temperature of 900 ° C and coiled under the conditions shown in Table 4 to obtain a hot-rolled sheet. The obtained hot-rolled sheet was pickled, and then cold-rolled under the conditions shown in Table 4 to obtain a cold-rolled sheet having a sheet thickness of 1.2 mm.
　得られた冷延鋼に、表４に示す条件で、第一めっき工程（金属電気めっき工程）、第１加熱工程、第１冷却工程、滞炉工程、第二めっき工程、第２冷却工程、第２加熱工程及びにおける処理を行い、鋼板を得た。 The obtained cold-rolled steel was subjected to the first plating process (metal electroplating process), the first heating process, the first cooling process, the furnace dwell process, the second plating process, the second cooling process, and the second heating process under the conditions shown in Table 4 to obtain a steel sheet.
　ここで、表４の金属電気めっき処理の有無(めっき種)の欄が、有(Fe)とある場合は、Ｆｅ系電気めっき処理を行った例であり、（Ni）とある場合は、Ｎｉ系電気めっき処理を行った例である。金属電気層の組成は、Ｆｅ系電気めっきでは、Ｆｅ：９５～１００質量％、Ｎｉ系電気めっきでは、Ｎｉ：９５～１００質量％を含有し、それぞれ残部は不可避的不純物であった。 In Table 4, in the column for metal electroplating (type of plating), if it says Yes (Fe), it is an example where Fe-based electroplating was performed, and if it says (Ni), it is an example where Ni-based electroplating was performed. The composition of the metal electroplating layer is Fe: 95-100% by mass for Fe-based electroplating, Ni: 95-100% by mass for Ni-based electroplating, and the remainder in each case is unavoidable impurities.
　ここで、第二めっき工程では、一部の熱延板（白皮）又は冷延板に対して、溶融亜鉛めっき処理又は合金化亜鉛めっき処理を行い、溶融亜鉛めっき鋼板（以下、ＧＩともいう）又は合金化溶融亜鉛めっき鋼板（以下、ＧＡともいう）を得た。
　亜鉛めっき浴温は、ＧＩ及びＧＡいずれを製造する場合も、４７０℃とした。
　亜鉛めっき付着量は、ＧＩを製造する場合は、片面あたり４５～７５ｇ/ｍ２（両面めっき）し、ＧＡを製造する場合は、片面あたり４０～６５ｇ/ｍ２（両面めっき）した。
　なお、最終的に得られた鋼板の亜鉛めっき層の組成は、ＧＩでは、Ｆｅ：０．１～１．０質量％、Ａｌ：０．２０～０．３３質量％を含有し、残部がＺｎ及び不可避的不純物であった。また、ＧＡでは、Ｆｅ：７．０～１２．０質量％、Ａｌ：０．１０～０．２３質量％を含有し、残部がＺｎ及び不可避的不純物であった。
　また、溶融亜鉛めっき層、合金化溶融亜鉛めっき層はいずれも、素地鋼板の両面に形成した。 Here, in the second plating process, a part of the hot-rolled sheet (white skin) or the cold-rolled sheet was subjected to a hot-dip galvanizing treatment or a galvannealed plating treatment to obtain a hot-dip galvanized steel sheet (hereinafter also referred to as GI) or a galvannealed hot-dip galvanized steel sheet (hereinafter also referred to as GA).
The zinc plating bath temperature was 470° C. for both GI and GA production.
The amount of zinc plating applied was 45 to 75 g/m 2 per side (double-sided plating) when producing GI, and 40 to 65 g/m 2 per side (double-sided plating) when producing GA.
The composition of the zinc plating layer of the finally obtained steel sheet was as follows: GI: 0.1-1.0 mass% Fe, 0.20-0.33 mass% Al, and the balance being Zn and unavoidable impurities, while GA: 7.0-12.0 mass% Fe, 0.10-0.23 mass% Al, and the balance being Zn and unavoidable impurities.
In addition, both the hot-dip galvanized layer and the galvannealed layer were formed on both sides of the base steel sheet.
［表層軟質層の厚さ］
　表層軟質層の測定方法は、以下のとおりである。めっき層（溶融亜鉛めっき層又は合金化亜鉛めっき層、場合により金属電気めっき層）剥離後、素地鋼板の圧延方向に平行な板厚断面（Ｌ断面）を湿式研磨により平滑化した後、ビッカース硬度計を用いて、荷重１０ｇｆで、鋼板表面から板厚方向に１μｍの位置より、板厚方向１００μｍの位置まで、１μｍ間隔で測定を行った。その後は板厚中心まで２０μｍ間隔で測定を行った。硬度が板厚１／４位置の硬度に比して８５％以下に減少した領域を軟質層（表層軟質層）と定義し、当該領域の板厚方向の厚さを軟質層の厚さとした。 [Thickness of surface soft layer]
The method for measuring the surface soft layer is as follows. After peeling off the plating layer (hot-dip galvanized layer or galvannealed layer, and sometimes metal electroplated layer), the thickness cross section (L cross section) parallel to the rolling direction of the base steel sheet was smoothed by wet polishing, and then measurements were made at 1 μm intervals using a Vickers hardness tester from a position 1 μm in the thickness direction from the steel sheet surface to a position 100 μm in the thickness direction under a load of 10 gf. Thereafter, measurements were made at 20 μm intervals to the center of the sheet thickness. The region in which the hardness was reduced to 85% or less compared to the hardness at the 1/4 position of the sheet thickness was defined as the soft layer (surface soft layer), and the thickness in the thickness direction of the region was defined as the thickness of the soft layer.
［表層軟質層のナノ硬度］
　めっき層（溶融亜鉛めっき層又は合金化亜鉛めっき層、場合により金属電気めっき層）剥離後、素地鋼板表面から表層軟質層の板厚方向深さの１／４位置まで機械研磨、ダイヤモンド及びアルミナでのバフ研磨及びコロイダルシリカ研磨を実施した。ナノインデンテーション装置（Ｈｙｓｉｔｒｏｎ社のｔｒｉｂｏ－９５０）を用いて、バーコビッチ形状のダイヤモンド圧子により、
　荷重速度及び除荷速度：５０μＮ／ｓ
　最大荷重：５００μＮ
　測定領域：５０μｍ×５０μｍ
　データ採取ピッチ：５msec
　打点間隔：２μｍ
の条件で計５１２点のナノ硬度を測定した。
　次いで、前記表層軟質層の板厚方向深さの１／２位置まで機械研磨、ダイヤモンド及びアルミナでのバフ研磨及びコロイダルシリカ研磨を実施した。Ｈｙｓｉｔｒｏｎ社のｔｒｉｂｏ－９５０を用い、バーコビッチ形状のダイヤモンド圧子により、上記と同様の条件で計５１２点のナノ硬度を測定した。 [Nano hardness of the soft surface layer]
After peeling off the plating layer (hot-dip galvanized layer or alloyed galvanized layer, or sometimes metal electroplated layer), mechanical polishing was performed from the surface of the base steel sheet to a position 1/4 of the depth in the sheet thickness direction of the soft surface layer, buff polishing with diamond and alumina, and colloidal silica polishing were performed.
Loading rate and unloading rate: 50 μN/s
Maximum load: 500 μN
Measurement area: 50 μm × 50 μm
Data sampling pitch: 5 msec
Dot spacing: 2 μm
The nano-hardness was measured at a total of 512 points under the above conditions.
Next, the surface soft layer was mechanically polished to a position of 1/2 the depth in the plate thickness direction, buffed with diamond and alumina, and polished with colloidal silica. Using a Tribo-950 manufactured by Hysitron, a Berkovich-shaped diamond indenter, the nano hardness was measured at a total of 512 points under the same conditions as above.
　得られた鋼板について、上記の試験方法に従い、部品の強度、延性、伸びフランジ性、剪断端面部の曲げ性及び低温靭性を評価した。結果を表５に併記する。
　さらに、以下の試験方法に従い、Ｖ曲げ＋直交ＶＤＡ曲げ試験、軸圧壊破断試験を行った。結果を表５に併記する。
　なお、板厚１．２ｍｍ超の溶融亜鉛めっき鋼板のＶ曲げ＋直交ＶＤＡ曲げ試験及び軸圧壊試験では、板厚の影響を考慮し、全て板厚１．２ｍｍの鋼板で実施した。板厚１．２ｍｍ超の鋼板は片面研削し、板厚を１．２ｍｍにした。一方、板厚１．２未満の溶融亜鉛めっき鋼板のＶ曲げ＋直交ＶＤＡ曲げ試験及び軸圧壊試験では、板厚の影響が小さいため、研削処理無しで試験を行った。 The obtained steel sheets were evaluated for strength, ductility, stretch flangeability, bendability of the sheared end surface, and low temperature toughness according to the above-mentioned test methods. The results are shown in Table 5.
Furthermore, a V-bend + orthogonal VDA bending test and an axial compression fracture test were performed according to the following test methods. The results are shown in Table 5.
In addition, in the V-bend + orthogonal VDA bending test and axial crush test of hot-dip galvanized steel sheets with a thickness of more than 1.2 mm, the test was performed on steel sheets with a thickness of 1.2 mm, taking into consideration the influence of the sheet thickness. The steel sheets with a thickness of more than 1.2 mm were ground on one side to make the sheet thickness 1.2 mm. On the other hand, in the V-bend + orthogonal VDA bending test and axial crush test of hot-dip galvanized steel sheets with a thickness of less than 1.2 mm, the test was performed without grinding because the influence of the sheet thickness was small.
《Ｖ曲げ＋直交ＶＤＡ曲げ試験》
　Ｖ曲げ＋直交ＶＤＡ曲げ試験は以下のようにして実施した。
　得られた鋼板から、６０ｍｍ×６５ｍｍの試験片を剪断・端面研削加工により採取した。ここで、６０ｍｍの辺は圧延（Ｌ）方向に平行とする。曲率半径/板厚：４．２で幅（Ｃ）方向を軸に圧延（Ｌ）方向に９０°曲げ加工（一次曲げ加工）を施し、試験片を準備した。９０°曲げ加工（一次曲げ加工）では、図２（ａ）に示すように、Ｖ溝を有するダイＡ１の上に載せた鋼板に対して、パンチＢ１を押し込んで試験片Ｔ１を得た。次に、図２（ｂ）に示すように、支持ロールＡ２の上に載せた試験片Ｔ１に対して、曲げ方向が圧延直角方向となるようにして、パンチＢ２を押し込んで直交曲げ（二次曲げ加工）を施した。図２（ａ）及び図２（ｂ）において、Ｄ１は幅（Ｃ）方向、Ｄ２は圧延（Ｌ）方向を示している。 <<V-bend + orthogonal VDA bending test>>
The V-bend + orthogonal VDA bend test was carried out as follows.
From the obtained steel plate, a 60 mm x 65 mm test piece was taken by shearing and end face grinding. Here, the 60 mm side is parallel to the rolling (L) direction. The test piece was prepared by bending the steel plate 90° (primary bending) in the rolling (L) direction with the width (C) direction as the axis at a curvature radius/plate thickness of 4.2. In the 90° bending (primary bending), as shown in FIG. 2(a), a punch B1 was pressed into a steel plate placed on a die A1 having a V groove to obtain a test piece T1 . Next, as shown in FIG. 2(b), a punch B2 was pressed into the test piece T1 placed on a support roll A2 so that the bending direction was perpendicular to the rolling direction, and an orthogonal bending (secondary bending) was performed. In FIG. 2(a) and FIG. 2(b), D1 indicates the width (C) direction, and D2 indicates the rolling (L) direction.
　Ｖ曲げ＋直交ＶＤＡ曲げ試験におけるＶ曲げの条件は、以下のとおりである。
　試験方法：ダイ支持、パンチ押し込み
　成型荷重：１０ｔｏｎ
　試験速度：３０ｍｍ/ｍｉｎ
　保持時間：５ｓ
　曲げ方向：圧延（Ｌ）方向 The V-bending conditions in the V-bending + orthogonal VDA bending test are as follows.
Test method: Die support, punch pressing Molding load: 10 tons
Test speed: 30 mm/min
Holding time: 5 s
Bending direction: Rolling (L) direction
　Ｖ曲げ＋直交ＶＤＡ曲げ試験におけるＶＤＡ曲げの条件は、以下のとおりである。
　試験方法：ロール支持、パンチ押し込み
　ロール径：φ３０ｍｍ
　パンチ先端Ｒ：０．４ｍｍ
　ロール間距離：（板厚×２）＋０．５ｍｍ
　ストローク速度：２０ｍｍ/ｍｉｎ
　試験片サイズ：６０ｍｍ×６０ｍｍ
　曲げ方向：圧延直角（Ｃ）方向 The VDA bending conditions in the V-bending + orthogonal VDA bending test are as follows.
Test method: Roll support, punch pressing Roll diameter: φ30 mm
Punch tip R: 0.4 mm
Distance between rolls: (sheet thickness x 2) + 0.5 mm
Stroke speed: 20 mm/min
Test piece size: 60 mm x 60 mm
Bending direction: perpendicular to rolling (C) direction
　前記ＶＤＡ曲げを施した際に得られるストローク－荷重曲線において、荷重最大時のストロークを求める。前記Ｖ曲げ＋直交ＶＤＡ曲げ試験を３回実施した際の当該荷重最大時のストロークの平均値をＳＦｍａｘ（ｍｍ）とした。求めたＳＦｍａｘが、２６．０ｍｍ以上を満たす場合、衝突時の耐破断性（曲げ破断に対する耐性）に優れると判断した。 The stroke at maximum load was determined from the stroke-load curve obtained when the VDA bending test was performed. The average stroke at maximum load when the V bending + orthogonal VDA bending test was performed three times was taken as SF max (mm). When the determined SF max was 26.0 mm or more, it was determined that the fracture resistance (resistance to bending fracture) during a collision was excellent.
《軸圧壊試験》
　軸圧壊試験は、以下のようにして実施した。
　得られた鋼板から、１５０ｍｍ×１００ｍｍの試験片を剪断加工により採取した。ここで、１５０ｍｍの辺は圧延（Ｌ）方向に平行とする。パンチ肩半径が５．０ｍｍであり、ダイ肩半径が５．０ｍｍである金型を用いて、深さ４０ｍｍとなるように成形加工（曲げ加工）して、図３（ａ）及び図３（ｂ）に示すハット型部材１０を作製した。
　また、ハット型部材の素材として用いた鋼板を、８０ｍｍ×１００ｍｍの大きさに別途切り出した。次に、その切り出した後の鋼板２０と、ハット型部材１０とをスポット溶接し、図３（ａ）及び図３（ｂ）に示すような試験用部材３０を作製した。図３（ａ）は、ハット型部材１０と鋼板２０とをスポット溶接して作製した試験用部材３０の正面図である。図３（ｂ）は、試験用部材３０の斜視図である。スポット溶接部４０の位置は、図３（ｂ）に示すように、鋼板の端部と溶接部が１０ｍｍ、溶接部間が２０ｍｍの間隔となるようにした。次に、図３（ｃ）に示すように、試験用部材３０を地板５０とＴＩＧ溶接により接合して軸圧壊試験用サンプルを作製した。次に、作製した軸圧壊試験用サンプルにインパクター６０を衝突速度１０ｍｍ/ｍｉｎで等速衝突させ、軸圧壊試験用のサンプルを７０ｍｍ圧壊した。図３（ｃ）に示すように、圧壊方向Ｄ３は、試験用部材３０の長手方向と平行な方向とした。
　試験後の試験体３０の外観を観察し、軸圧壊破断（外観割れ）の有無を確認した。
　外観割れが認められなかった場合は「Ａ」を、外観割れが１箇所以下で認められた場合は「Ｂ」を、外観割れが２箇所以上で認められた場合は「Ｃ」を、下記表５に記載した。「Ａ」又は「Ｂ」の場合、衝突時の耐破断性（軸圧壊破断に対する耐性）に優れると判断した。 <Axial crushing test>
The axial crush test was carried out as follows.
A test piece of 150 mm x 100 mm was taken from the obtained steel plate by shearing. Here, the 150 mm side was parallel to the rolling (L) direction. Using a die having a punch shoulder radius of 5.0 mm and a die shoulder radius of 5.0 mm, the test piece was molded (bended) to a depth of 40 mm to produce the hat-shaped member 10 shown in Figures 3(a) and 3(b).
In addition, the steel plate used as the material for the hat-shaped member was separately cut to a size of 80 mm x 100 mm. Next, the cut-out steel plate 20 and the hat-shaped member 10 were spot-welded to prepare a test member 30 as shown in Figures 3(a) and 3(b). Figure 3(a) is a front view of the test member 30 prepared by spot welding the hat-shaped member 10 and the steel plate 20. Figure 3(b) is a perspective view of the test member 30. As shown in Figure 3(b), the position of the spot welded portion 40 was such that the end of the steel plate and the welded portion were spaced 10 mm apart, and the welded portions were spaced 20 mm apart. Next, as shown in Figure 3(c), the test member 30 was joined to the base plate 50 by TIG welding to prepare a sample for an axial crush test. Next, an impactor 60 was made to collide with the prepared sample for an axial crush test at a constant speed of 10 mm/min, and the sample for an axial crush test was crushed by 70 mm. As shown in FIG. 3C, the crushing direction D3 was parallel to the longitudinal direction of the test member 30.
After the test, the appearance of the test specimen 30 was observed to check for the presence or absence of axial compression fracture (external cracks).
When no cracks were observed, the sample was rated as "A", when one or less cracks were observed, the sample was rated as "B", and when two or more cracks were observed, the sample was rated as "C" in Table 5. Samples rated as "A" or "B" were judged to have excellent resistance to fracture during a collision (resistance to fracture due to axial pressure).
Figure JPOXMLDOC01-appb-T000004
Figure JPOXMLDOC01-appb-T000004
Figure JPOXMLDOC01-appb-T000005
Figure JPOXMLDOC01-appb-T000005
　表５に示すように、本発明例では、ＴＳが７８０ＭＰａ以上であり、部品の強度、延性、伸びフランジ性、剪断端面部の曲げ性、低温靭性及び衝突時の耐破断特性（曲げ破断特性及び軸圧壊特性）が優れている。 As shown in Table 5, in the present invention example, TS is 780 MPa or more, and the part has excellent strength, ductility, stretch flangeability, bendability of the sheared end surface, low-temperature toughness, and fracture resistance during collision (bending fracture properties and axial crush properties).
　本発明によれば、部品強度、伸びフランジ性、剪断端面部の曲げ性、低温靭性及び衝突時の耐破断特性（曲げ破断特性及び軸圧壊特性）に優れる高強度鋼板が提供される。特に、本発明の高強度鋼板は、種々の特性に優れるので、種々の大きさ及び形状の自動車の骨格構造部品や補強部品等に適用することが可能である。これにより、車体軽量化による燃費向上を図ることができ、産業上の利用価値は極めて大きい。 The present invention provides a high-strength steel plate that is excellent in part strength, stretch flangeability, bendability of the sheared end surface, low-temperature toughness, and fracture resistance during collision (bending fracture properties and axial crush properties). In particular, the high-strength steel plate of the present invention has a variety of excellent properties, making it applicable to structural components and reinforcing components of automobiles of various sizes and shapes. This allows for improved fuel efficiency by reducing the weight of the vehicle body, and is of great industrial value.
Ａ１　ダイ
Ａ２　支持ロール
Ｂ１　パンチ
Ｂ２　パンチ
Ｔ１　試験片
Ｄ１　幅（Ｃ）方向
Ｄ２　圧延（Ｌ）方向
Ｄ３　圧壊方向
１０　ハット型部材
２０　鋼板
３０　試験用部材
４０　スポット溶接部
５０　地板
６０　インパクター A1 Die A2 Support roll B1 Punch B2 Punch T1 Test piece D1 Width (C) direction D2 Rolling (L) direction D3 Crushing direction 10 Hat-shaped member 20 Steel plate 30 Test member 40 Spot welded portion 50 Base plate 60 Impactor

Claims (14)

  1. 　質量％で、
    Ｃ：０．０３０％以上０．５００％以下、
    Ｓｉ：０．０１％以上２．５０％以下、
    Ｍｎ：０．１０％以上５．００％以下、
    Ｐ：０．１００％以下、
    Ｓ：０．０２００％以下、
    Ａｌ：１．０００％以下、
    Ｎ：０．０１００％以下及び
    Ｏ：０．０１００％以下
    を含有し、残部がＦｅ及び不可避的不純物からなる成分組成と、
    　板厚１／４位置において、
    マルテンサイトの面積率が１０％以上８０％以下、
    ベイナイトの面積率が２％以上７０％以下、
    フェライトの面積率が８０％以下、
    残留オーステナイトの面積率が１５％以下、かつ
    マルテンサイトブロック数に対する準安定炭化物が存在するマルテンサイトブロック数の割合が２％以上である鋼組織と、
    を有し、
    　板厚１／４位置において、２２５点以上のナノ硬度を測定したとき、ナノ硬度の平均値［Ｈｎ］aveに対して、ナノ硬度の標準偏差σnが０．６０×［Ｈｎ］ave以下である、
    高強度鋼板。     In mass percent,
    C: 0.030% or more and 0.500% or less,
    Si: 0.01% or more and 2.50% or less,
    Mn: 0.10% or more and 5.00% or less,
    P: 0.100% or less,
    S: 0.0200% or less,
    Al: 1.000% or less,
    A composition comprising N: 0.0100% or less and O: 0.0100% or less, with the balance being Fe and unavoidable impurities;
    At the 1/4 plate thickness position,
    The area ratio of martensite is 10% or more and 80% or less,
    The area ratio of bainite is 2% or more and 70% or less,
    The area ratio of ferrite is 80% or less,
    a steel structure in which the area ratio of retained austenite is 15% or less and the ratio of the number of martensite blocks in which metastable carbides are present to the number of martensite blocks is 2% or more;
    having
    When the nanohardness is measured at 225 or more points at the 1/4 position of the sheet thickness, the standard deviation σ n of the nanohardness is 0.60 × [H n ] ave or less with respect to the average value [H n ] ave of the nanohardness.
    High strength steel plate.
  2. 　前記準安定炭化物が存在するマルテンサイトブロックにおける準安定炭化物の個数密度の平均値が１×１０6個/ｍｍ２以上である、請求項１に記載の高強度鋼板。 2. The high-strength steel plate according to claim 1, wherein an average number density of the metastable carbides in the martensite block in which the metastable carbides are present is 1×10 6 /mm 2 or more.
  3. 　前記成分組成は、さらに、質量％で、
    Ｔｉ：０．２００％以下、
    Ｎｂ：０．２００％以下、
    Ｖ：０．２００％以下、
    Ｔａ：０．１０％以下、
    Ｗ：０．１０％以下、
    Ｂ：０．０１００％以下、
    Ｃｒ：１．００％以下、
    Ｍｏ：１．００％以下、
    Ｎｉ：１．００％以下、
    Ｃｏ：０．０１０％以下、
    Ｃｕ：１．００％以下、
    Ｓｎ：０．２００％以下、
    Ｓｂ：０．２００％以下、
    Ｃａ：０．０１００％以下、
    Ｍｇ：０．０１００％以下、
    ＲＥＭ：０．０１００％以下、
    Ｚｒ：０．１００％以下、
    Ｔｅ：０．１００％以下、
    Ｈｆ：０．１０％以下及び
    Ｂｉ：０．２００％以下
    からなる群より選ばれる少なくとも１種の元素を含有する、請求項１又は２に記載の高強度鋼板。     The composition further includes, in mass%,
    Ti: 0.200% or less,
    Nb: 0.200% or less,
    V: 0.200% or less,
    Ta: 0.10% or less,
    W: 0.10% or less,
    B: 0.0100% or less,
    Cr: 1.00% or less,
    Mo: 1.00% or less,
    Ni: 1.00% or less,
    Co: 0.010% or less,
    Cu: 1.00% 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: 0.100% or less,
    Te: 0.100% or less,
    The high strength steel plate according to claim 1 or 2, containing at least one element selected from the group consisting of Hf: 0.10% or less and Bi: 0.200% or less.
  4. 　前記高強度鋼板の板厚１／４位置のビッカース硬さに対して、ビッカース硬さが８５％以下の領域であって、前記高強度鋼板表面から板厚方向に２００μｍ以内の領域である表層軟質層を有し、
    　前記高強度鋼板表面から前記表層軟質層の板厚方向深さの１／４位置及び板厚方向深さの１／２位置のそれぞれにおける板面の５０μｍ×５０μｍの領域において、３００点以上のナノ硬度を測定したとき、
    　前記高強度鋼板表面から前記表層軟質層の板厚方向深さの１／４位置の板面のナノ硬度が７．０ＧＰａ以上の測定数割合が、全測定数に対して０．１０以下であり、
    　前記高強度鋼板表面から前記表層軟質層の板厚方向深さの１／４位置の板面のナノ硬度の標準偏差σが１．８ＧＰａ以下であり、
    　さらに、前記高強度鋼板表面から前記表層軟質層の板厚方向深さの１／２位置の板面のナノ硬度の標準偏差σが２．２ＧＰａ以下である、請求項１～３のいずれか一項に記載の高強度鋼板。     The high-strength steel plate has a Vickers hardness of 85% or less of the Vickers hardness at a 1/4 position in the plate thickness direction of the high-strength steel plate, and has a surface soft layer which is a region within 200 μm from the surface of the high-strength steel plate in the plate thickness direction;
    When the nano hardness was measured at 300 points or more in a 50 μm × 50 μm region of the sheet surface at a 1/4 position and a 1/2 position of the sheet thickness direction depth of the soft surface layer from the surface of the high strength steel sheet,
    The ratio of the number of measurements in which the nano hardness of the sheet surface at a position of 1/4 of the sheet thickness direction depth of the soft surface layer from the surface of the high strength steel sheet is 7.0 GPa or more to the total number of measurements is 0.10 or less,
    The standard deviation σ of the nano-hardness of the sheet surface at a ¼ position of the sheet thickness direction depth of the soft surface layer from the surface of the high-strength steel sheet is 1.8 GPa or less;
    Further, the standard deviation σ of the nano hardness of the sheet surface at a position of 1/2 of the sheet thickness direction depth of the soft surface layer from the surface of the high strength steel sheet is 2.2 GPa or less.
  5. 　前記高強度鋼板の片面又は両面の表面上において、Ｃｒ、Ｍｎ、Ｆｅ、Ｃｏ、Ｎｉ、Ｃｕ、Ｇａ、Ｇｅ、Ａｓ、Ｒｕ、Ｒｈ、Ｐｄ、Ａｇ、Ｃｄ、Ｉｎ、Ｓｎ、Ｓｂ、Ｏｓ、Ｉｒ、Ｒｔ、Ａｕ、Ｈｇ、Ｔｉ、Ｐｂ及びＢｉから選択される１種又は２種以上を合計で５０質量％超含む金属めっき層を有する、請求項１～４のいずれか一項に記載の高強度鋼板。 The high-strength steel plate according to any one of claims 1 to 4, having a metal plating layer on one or both surfaces of the high-strength steel plate, the metal plating layer containing more than 50 mass% in total of one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb and Bi.
  6. 　前記高強度鋼板の片面又は両面の最外層に、亜鉛及びアルミニウムの少なくとも一方を合計で５０質量％以上含む金属めっき層を有する、請求項１～５のいずれか一項に記載の高強度鋼板。 The high-strength steel plate according to any one of claims 1 to 5, having a metal plating layer containing at least one of zinc and aluminum in total at 50 mass% or more on one or both outermost layers of the high-strength steel plate.
  7. 　請求項１～６のいずれか１項に記載の高強度鋼板を用いてなる、部材。 A member made using the high-strength steel plate according to any one of claims 1 to 6.
  8. 　請求項７に記載の部材からなる、自動車の骨格構造部品又は自動車の補強部品。 An automobile frame structural part or automobile reinforcing part made of the member described in claim 7.
  9. 　請求項１又は３に記載の成分組成を有する鋼スラブに、
    　平均のひずみ速度が１×１０－４／s以上１×１０－１／s以下、総圧下率５０％以上の条件で粗圧延を施した後、仕上げ圧延を施し、次いで巻取り処理を施して、熱延板を得る熱間圧延工程、
    　次いで、酸洗及び冷間圧延を施して、冷延板を得る酸洗及び冷間圧延工程と、
    　次いで、加熱温度が７５０℃以上の条件で第１加熱する第１加熱工程と、
    　次いで、Ｔ２以上７５０℃以下の温度域における第１冷却速度が２．０℃／ｓ以上の条件で冷却する第１冷却工程と、
    　次いで、３５０℃以上５５０℃以下の滞留温度Ｔ２で下記式１で定義されるＦが０．２０以上０．９０以下を満たす滞留時間ｔ（ｓ）の条件で滞炉させる滞炉工程と、
    　次いで、Ｍｓ－２０℃以下まで冷却する工程であって、Ｍｓ－２０℃以上Ｍs以下の温度域における第２平均冷却速度を５℃／ｓ以上の条件とする第２冷却工程と、
    　次いで、下記式２を満たす温度Ｘ（℃）と保持時間Ｙ（ｓ）の条件で処理する第２加熱工程と
    を含む、高強度鋼板の製造方法。
    　　　　　　　　　　　　　　　　　記
    式１：Ｆ＝１－ｅｘｐ（－ｋｔｎ）
    ｔ：滞留時間（ｓ）
    ｋ、ｎ：前記スラブを前記第１冷却工程終了までの工程に付して得られる試験片に対して、３５０℃以上５５０℃以下の滞留温度Ｔ２で保持させて行われるフォーマスター試験の膨張曲線から求められる定数。
    式２：７０００≦（２７３＋Ｘ）（２０＋ｌｏｇ(Ｙ/３６００)）≦１３０００     A steel slab having the composition according to claim 1 or 3,
    a hot rolling process in which rough rolling is performed under conditions of an average strain rate of 1×10 −4 /s or more and 1×10 −1 /s or less and a total rolling reduction of 50% or more, followed by finish rolling and then coiling to obtain a hot-rolled sheet;
    Next, pickling and cold rolling are performed to obtain a cold-rolled sheet.
    Next, a first heating step of performing a first heating under a condition of a heating temperature of 750° C. or more;
    Next, a first cooling step in which the first cooling rate in a temperature range from T2 to 750°C is 2.0°C/s or more;
    Next, a retention step of retaining the material in the furnace at a retention temperature T2 of 350° C. or more and 550° C. or less for a retention time t (s) in which F defined by the following formula 1 is 0.20 or more and 0.90 or less;
    Next, a second cooling step of cooling to Ms-20°C or less, in which a second average cooling rate in a temperature range of Ms-20°C or more and Ms or less is set to 5°C/s or more;
    Next, a second heating step is performed under conditions of a temperature X (° C.) and a holding time Y (s) that satisfy the following formula 2.
    Formula 1: F=1-exp(-kt n )
    t: residence time (s)
    k, n: Constants obtained from the expansion curve of a Formaster test in which a test piece obtained by subjecting the slab to the process up to the end of the first cooling process is held at a residence temperature T2 of 350°C or more and 550°C or less.
    Formula 2: 7000≦(273+X)(20+log(Y/3600))≦13000
  10. 　前記第２加熱工程において、温度Ｘ（℃）が下記式３を満たす、請求項９に記載の高強度鋼板の製造方法。
    　　　　　　　　　　　　　　　　　記
    式３：１００≦Ｘ≦４００     The method for producing a high strength steel plate according to claim 9 , wherein in the second heating step, the temperature X (° C.) satisfies the following formula 3:
    Formula 3: 100≦X≦400
  11. 　前記第１加熱工程を露点－３０℃以上の雰囲気下で行う、請求項９又は１０に記載の高強度鋼板の製造方法。 The method for manufacturing high-strength steel plate according to claim 9 or 10, wherein the first heating step is carried out in an atmosphere with a dew point of -30°C or higher.
  12. 　前記冷間圧延工程後、かつ前記焼鈍工程の前の鋼板の片面もしくは両面において、Ｃｒ、Ｍｎ、Ｆｅ、Ｃｏ、Ｎｉ、Ｃｕ、Ｇａ、Ｇｅ、Ａｓ、Ｒｕ、Ｒｈ、Ｐｄ、Ａｇ、Ｃｄ、Ｉｎ、Ｓｎ、Ｓｂ、Ｏｓ、Ｉｒ、Ｒｔ、Ａｕ、Ｈｇ、Ｔｉ、Ｐｂ及びＢｉから選択される１種又は２種以上を合計で５０質量％超含む金属めっきを施す工程を含む、請求項９～１１のいずれか一項に記載の高強度鋼板の製造方法。 The method for producing a high-strength steel sheet according to any one of claims 9 to 11, comprising the step of applying a metal plating containing more than 50 mass% in total of one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb and Bi to one or both sides of the steel sheet after the cold rolling step and before the annealing step.
  13. 　前記第１加熱から第２加熱工程の鋼板に、亜鉛及びアルミニウムの少なくとも一方を合計で５０質量％以上含む金属めっきを施す工程を含む、請求項９～１２のいずれか一項に記載の高強度鋼板の製造方法。 The method for manufacturing a high-strength steel sheet according to any one of claims 9 to 12, further comprising a step of applying a metal plating containing at least one of zinc and aluminum in total at 50 mass% to the steel sheet from the first heating step to the second heating step.
  14. 　請求項１～６のいずれか一項に記載の高強度鋼板に、成形加工又は接合加工の少なくとも一方を施して部材とする工程を有する、部材の製造方法。 A method for manufacturing a component, comprising the step of subjecting the high-strength steel plate according to any one of claims 1 to 6 to at least one of forming and joining processes to produce a component.
PCT/JP2022/045373 2022-12-08 2022-12-08 High-strength steel sheet, member formed using high-strength steel sheet, automobile framework structure component or automobile reinforcing component composed of member, and production methods for high-strength steel sheet and member WO2024122037A1 (en)

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