High-strength hot-rolled steel sheet and method for producing same

Abstract

Provided is a high-strength hot-rolled steel sheet having excellent strength, toughness, and delayed fracture resistance after post-heating. A high-strength hot-rolled steel sheet having a component composition containing, in mass%, 0.06-0.23% of C, 0.1-3.0% of Si, 1.5-3.5% of Mn, more than 0% but not more than 0.050% of P, more than 0% but not more than 0.0050% of S, more than 0% but not more than 1.5% of Al, more than 0% but not more than 0.010% of N, and more than 0% but not more than 0.003% of O, and further containing a total of 0.040-0.200% of Ti and Nb, with the remainder comprising Fe and unavoidable impurities; the steel structure has martensite and / or lower bainite as the main phase, the retained austenite is less than 3% by volume fraction, (the amount of solid-solution Ti + the amount of solid-solution Nb) / (the total amount of Ti + the total amount of Nb) is 0.300 or more and less than 0.800, the total amount of Ti and Nb present as precipitates having a particle diameter of 100 nm or more is 0.010-0.030 mass%, and the polar density of {110} < 111 > orientation in the surface layer region is 1.8-5.0.

Description

Technical Field

[0001] The present invention relates to a high-strength hot-rolled steel sheet and a method for manufacturing the same, and particularly to a high-strength hot-rolled steel sheet suitable as a blank for automotive parts and a method for manufacturing the same. Background Art

[0002] From the viewpoints of improving the collision safety of automobiles and increasing fuel efficiency, high strength is required for hot-rolled steel sheets used for automotive parts. On the other hand, in high-strength hot-rolled steel sheets, the generation of cracks due to insufficient workability during pressing becomes significant, so it is necessary to improve the pressing method and the workability of the steel sheet. Regarding the method, research is being conducted to improve workability by heating the steel sheet (original steel sheet). It should be noted that in this specification, the heat treatment applied when processing the steel sheet (original steel sheet) into parts, etc. is also referred to as post-heating. On the other hand, the development of steel sheets considering the characteristics of parts after processing (after post-heating processing) is underway. Moreover, in high-strength steel sheets with a tensile strength (TS) of 1180 MPa or more, delayed fracture also becomes a problem, so it is also necessary to design a blank considering the delayed fracture resistance characteristics after post-heating.

[0003] Patent Document 1 discloses a technique related to a method for improving stretch flangeability by setting the processing temperature (heating temperature of post-heating) to 400 to 1000°C. Patent Document 2 discloses a technique related to a high-strength hot-rolled steel sheet with a TS of 730 MPa or more. In the hot-rolled steel sheet disclosed in Patent Document 2, a structure mainly composed of bainite and having 80% or more of the total Ti amount as solid-solution Ti is formed. Thereby, the heat treatment hardenability with an increase in YS (yield strength) and TS of 100 MPa or more after heat treatment carried out by heating to the temperature range of 500°C to the Ac1 transformation point and holding for 60 min is obtained. Patent Document 3 discloses a technique related to a hot-rolled steel sheet having a TS of 120 kgf / mm 2 or more and excellent delayed fracture resistance characteristics.

[0004] Prior Art Documents

[0005] Patent Documents

[0006] Patent Document 1: Japanese Patent Application Laid-Open No. 2002-113527

[0007] Patent Document 2: Japanese Patent Application Laid-Open No. 2015-57514

[0008] Patent Document 3: Japanese Patent Application Laid-Open No. 6-145894 Summary of the Invention

[0009] However, in Patent Document 1, the strength, toughness, etc. after post-heating are not considered as the properties of the components, and there is room for improvement. In particular, when post-heating is carried out at a high temperature above 400 °C, the steel structure changes greatly. If the steel plate before post-heating (original steel plate) has a high strength of 1180 MPa or more, the influence on its strength becomes significant. Therefore, a billet design considering the influence of post-heating on strength is required. Although the steel plate disclosed in Patent Document 2 does not cause a decrease in strength after heat treatment, it instead causes an excessive increase in strength, and the fine precipitation of carbides is significant. Therefore, there are problems in terms of toughness after heat treatment, and there is room for improvement. In addition, in Patent Document 2, the strength grade of the studied steel plate only stays around 980 Mpa, and there are no insights and teachings regarding high-strength steel above 1180 MPa. Patent Document 3 aims to improve the workability of the original steel plate and is to improve the delayed fracture resistance characteristics without post-heating. However, in Patent Document 3, the delayed fracture test is evaluated by deep drawing, and the delayed fracture resistance characteristics of the more severe end face and the end face after processing after post-heating are not considered, and there is room for improvement.

[0010] The present invention has been completed in view of the above circumstances, and its object is to provide a high-strength hot-rolled steel plate and a manufacturing method thereof having excellent strength, toughness, and delayed fracture resistance characteristics after post-heating.

[0011] The inventors of the present invention focused on the precipitation behavior of Ti and Nb after post-heating of the hot-rolled steel plate, and came up with the idea of improving the characteristics of the steel plate after heating (after post-heating) by controlling the initial coarse Ti-containing precipitates and Nb-containing precipitates before post-heating, as well as the amounts of solid-solution Ti and solid-solution Nb. Furthermore, focusing on the crystal orientation, the inventors came up with the idea of suppressing the delayed fracture after post-heating of the blanking end face by forming the surface layer region into a structure aggregated along a specific orientation. As a result, it was found that on the basis of adjusting the chemical composition, with martensite and / or lower bainite as the main phase, the amount of retained austenite (retained γ) is less than 3% by volume, the pole density of the {110}<111> orientation is 1.8 to 5.0 in the surface layer region from the surface along the center direction of the plate thickness to 100 μm, the ratio of the total amount of solid-solution Ti and solid-solution Nb to the total amount of Ti content and Nb content (solid-solution Ti amount + solid-solution Nb amount) / (total Ti amount + total Nb amount) is 0.300 or more and less than 0.800, and the total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more is 0.010 to 0.030 mass%. Thus, high strength is exhibited in the hot-rolled steel plate (original steel plate), and after post-heating, a strength close to that of the original steel plate, as well as excellent toughness and delayed fracture resistance characteristics, can be obtained, thereby completing the present invention.

[0012] It should be noted that in the present invention, high strength means that the tensile strength (TS) is 1180 MPa or more and less than 1600 MPa.

[0013] In addition, in the present invention, excellent strength after post-heating means that the reduction in the strength of the hot-rolled steel sheet after post-heating is 50 or less with a Vickers hardness tester relative to the strength of the hot-rolled steel sheet (original steel sheet) before post-heating.

[0014] In the present invention, excellent toughness after post-heating means that in the Charpy impact test using a test piece collected from the hot-rolled steel sheet after post-heating, the ductile fracture rate at -20°C is 50% or more. It should be noted that the thickness of the above test piece is 0.6 to 3.0 mm. When the thickness of the hot-rolled steel sheet is greater than 3.0 mm, the test piece collected from the hot-rolled steel sheet is ground on both the front and back surfaces to a thickness of 3.0 mm for the Charpy impact test.

[0015] In the present invention, excellent delayed fracture resistance after post-heating means that a long test piece with a sheared end face collected from the steel sheet after post-heating is subjected to a V-bending process at 90°, and then the part opened due to springback is fastened with bolts or the like, and no cracks are generated when immersed in hydrochloric acid at pH 3 for 96 hours.

[0016] It should be noted that in the present invention, post-heating refers to heat treatment in which the hot-rolled steel sheet (original steel sheet) is heated to 400°C or higher.

[0017] The present invention has the following configuration.

[0018] [1] A high-strength hot-rolled steel sheet having the following composition:

[0019] Containing, by mass%,

[0020] C: 0.06 to 0.23%,

[0021] Si: 0.1 to 3.0%,

[0022] Mn: 1.5 to 3.5%,

[0023] P: greater than 0% and 0.050% or less,

[0024] S: greater than 0% and 0.0050% or less,

[0025] Al: greater than 0% and 1.5% or less,

[0026] N: greater than 0% and 0.010% or less,

[0027] O: greater than 0% and 0.003% or less,

[0028] And further containing a total of 0.040 to 0.200% of Ti and Nb,

[0029] The balance is composed of Fe and inevitable impurities;

[0030] Moreover, the steel structure has martensite and / or lower bainite as the main phase, and the retained austenite is less than 3% by volume fraction.

[0031] The ratio of the total amount of dissolved Ti and the total amount of dissolved Nb to the total amount of Ti and the total amount of Nb, i.e., (dissolved Ti amount + dissolved Nb amount) / (total Ti amount + total Nb amount), is 0.300 or more and less than 0.800.

[0032] The total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more is 0.010 to 0.030% by mass.

[0033] In the surface layer region from the surface along the center direction of the plate thickness to 100 μm, the pole density of the {110}<111> orientation is 1.8 to 5.0.

[0034] [2] The high-strength hot-rolled steel sheet according to [1], wherein the above composition further contains, by mass%, selected from

[0035] Cr: 0.005 to 2.0%,

[0036] Ni: 0.005 to 2.0%,

[0037] Mo: 0.005 to 1.0%,

[0038] V: 0.005 to 0.5%,

[0039] B: 0.0002 to 0.0050%,

[0040] Ca: 0.0001 to 0.0050%,

[0041] REM: 0.0001 to 0.0050%,

[0042] Cu: 0.005 to 0.5%,

[0043] Sb: 0.0010 to 0.10%, and

[0044] Sn: 0.0010 to 0.10%

[0045] one or more of the above.

[0046] [3] A method for manufacturing a high-strength hot-rolled steel sheet, which is the method for manufacturing the high-strength hot-rolled steel sheet according to the above [1] or [2],

[0047] A slab having the above composition is heated in a temperature range of 1150 to 1300 °C and held in this temperature range for 0.2 to 3.5 hours.

[0048] Then, when performing hot rolling,

[0049] The total reduction ratio in the temperature range above 1080°C is 80 - 90%, the total reduction ratio in the temperature range below 900°C is 20% or more, and the reduction ratio in a single pass below T (°C) calculated by the following formula is 25% or less. After rolling under these conditions, air cooling is performed for 1.0 s or more.

[0050] Next, it is cooled in the temperature range up to 550°C at an average cooling rate of 50°C / s or more, and the time from reaching 550°C to the start of quenching is set to 0.5 - 4.0 s.

[0051] Next, it is quenched to a coiling temperature of 100 - 250°C at a cooling rate of 200°C / s or more and coiled at the above coiling temperature.

[0052] T (°C) = 800 + 1000[Ti] + 2500[Nb]

[0053] Here, [Ti] and [Nb] are the contents (mass%) of Ti and Nb respectively, and are 0 when not contained.

[0054] According to the present invention, it is possible to provide a high-strength hot-rolled steel sheet having excellent strength, toughness, and hydrogen embrittlement resistance after post-heating, and a method for manufacturing the same.

[0055] According to the present invention, a high-strength hot-rolled steel sheet having excellent strength, toughness, and hydrogen embrittlement resistance after post-heating or after post-heat treatment, which is suitable as a blank for automotive parts, can be obtained.

[0056] If the high-strength hot-rolled steel sheet of the present invention is used, even after a heat treatment is performed to improve workability, fatigue characteristics, etc., products such as high-strength automotive parts that exhibit high strength, good toughness, and excellent hydrogen embrittlement resistance can be obtained. Detailed Embodiments

[0057] Hereinafter, embodiments of the high-strength hot-rolled steel sheet and the method for manufacturing the same according to the present invention will be described in detail. It should be noted that the present invention is not limited to the following embodiments.

[0058] <High-Strength Hot-Rolled Steel Sheet>

[0059] The high-strength hot-rolled steel sheet of the present invention can be either a black skin steel sheet in the as-hot-rolled state or a hot-rolled steel sheet called a white skin steel sheet that has been pickled after hot rolling. In addition, the target high-strength hot-rolled steel sheet of the present invention preferably has a thickness of 0.6 mm or more. In addition, the high-strength hot-rolled steel sheet of the present invention preferably has a thickness of 10.0 mm or less. When the high-strength hot-rolled steel sheet of the present invention is used as a blank for automotive parts, the thickness is more preferably 1.0 mm or more. In addition, when the high-strength hot-rolled steel sheet of the present invention is used as a blank for automotive parts, the thickness is more preferably 6.0 mm or less. In addition, the sheet width of the high-strength hot-rolled steel sheet of the present invention is preferably 500 mm or more, more preferably 700 mm or more. The sheet width of the high-strength hot-rolled steel sheet of the present invention is preferably 1800 mm or less, more preferably 1400 mm or less.

[0060] The high-strength hot-rolled steel sheet of the present invention has a specific composition and a specific steel structure. Here, the description will be made in the order of the composition and the steel structure.

[0061] First, the composition of the high-strength hot-rolled steel sheet of the present invention will be described. It should be noted that "%" indicating the content of the composition refers to "% by mass".

[0062] The composition of the high-strength hot-rolled steel sheet of the present invention contains, by mass%, C: 0.06 to 0.23%, Si: 0.1 to 3.0%, Mn: 1.5 to 3.5%, P: 0.050% or less (excluding 0), S: 0.0050% or less (excluding 0), Al: 1.5% or less (excluding 0), N: 0.010% or less (excluding 0), O: 0.003% or less (excluding 0), a total of 0.040 to 0.200% of Ti and Nb, and the balance is composed of Fe and inevitable impurities.

[0063] C: 0.06 to 0.23%

[0064] C is an element that is effective in improving TS by generating and strengthening martensite and lower bainite or by forming precipitates in combination with Ti, Nb, N, etc. to inhibit the reduction of strength after post-heating. If the C content is less than 0.06%, such an effect cannot be obtained sufficiently, and the TS of the steel sheet (original steel sheet) of 1180 MPa or more or the excellent strength after post-heating cannot be obtained. On the other hand, if the C content is greater than 0.23%, the reduction of toughness after post-heating becomes significant, and excellent toughness after post-heating cannot be obtained. Therefore, the C content is 0.06 to 0.23%. The C content is preferably 0.07% or more. In addition, the C content is preferably 0.22% or less, more preferably 0.20% or less.

[0065] Si: 0.1 to 3.0%

[0066] Si is an element effective for solid solution strengthening of steel and suppressing the reduction in strength after post-heating. To obtain such an effect, the Si content needs to be 0.1% or more. On the other hand, if the Si content is greater than 3.0%, excessive formation of polygonal ferrite occurs and the steel structure of the present invention cannot be obtained. Therefore, the Si content is 0.1 - 3.0%. The Si content is preferably 0.2% or more. In addition, the Si content is preferably 2.0% or less, more preferably 1.5% or less.

[0067] Mn: 1.5 - 3.5%

[0068] Mn is an element effective for suppressing ferrite, upper bainite and generating lower bainite, martensite. If the Mn content is less than 1.5%, such an effect cannot be fully obtained, and polygonal ferrite, upper bainite, etc. are generated, and the microstructure of the present invention cannot be obtained. On the other hand, if the Mn content is greater than 3.5%, the reduction in toughness and the reduction in delayed fracture resistance become significant, and excellent toughness and delayed fracture resistance after post-heating cannot be obtained. Therefore, the Mn content is 1.5 - 3.5%. The Mn content is preferably 1.6% or more. In addition, the Mn content is preferably 3.0% or less, more preferably 2.5% or less.

[0069] P: greater than 0% and 0.050% or less

[0070] P reduces the toughness and delayed fracture resistance after post-heating, so it is preferably reduced as much as possible. In the present invention, the P content can be allowed up to 0.050%. Therefore, the P content is 0.050% or less. The P content is preferably 0.030% or less. The lower limit is not particularly limited, and the P content can be greater than 0%, but if the P content is less than 0.001%, the production efficiency decreases, so the P content is preferably 0.001% or more.

[0071] S: greater than 0% and 0.0050% or less

[0072] S reduces the toughness and delayed fracture resistance after post-heating, so it is preferably reduced as much as possible. In the present invention, the S content can be allowed up to 0.0050%. Therefore, the S content is 0.0050% or less. The S content is preferably 0.0030% or less, more preferably 0.0020% or less, further preferably 0.0015%. The lower limit is not particularly limited, and the S content can be greater than 0%, but if the S content is less than 0.0002%, the production efficiency decreases, so the S content is preferably 0.0002% or more.

[0073] Al: greater than 0% and 1.5% or less

[0074] Al acts as a deoxidizer and is preferably added in the deoxidation process. The Al content can be greater than 0%, but from the perspective of using it as a deoxidizer, the Al content is preferably 0.01% or more. On the other hand, if a large amount of Al is contained, a large amount of polygonal ferrite is generated and the steel structure of the present invention cannot be obtained. In the present invention, the Al content can be allowed up to 1.5%. Therefore, the Al content is 1.5% or less. The Al content is preferably 0.50% or less, and more preferably 0.20% or less.

[0075] N: greater than 0% and 0.010% or less

[0076] N forms TiN and NbC, hindering the precipitation of fine TiC, NbC, etc. Therefore, it is preferred that its amount be minimized as much as possible. In the present invention, the N content can be allowed up to 0.010%. Therefore, the N content is 0.010% or less. The N content is preferably 0.007% or less. The lower limit is not particularly limited. The N content can be greater than 0%, but if the N content is less than 0.0005%, the production efficiency decreases. Therefore, the N content is preferably 0.0005% or more.

[0077] O: greater than 0% and 0.003% or less

[0078] O reduces the toughness and stress corrosion cracking resistance after post-heating. Therefore, it is preferred that its amount be minimized as much as possible. In the present invention, the O content can be allowed up to 0.003%. Therefore, the O content is 0.003% or less. The O content is preferably 0.002% or less. The lower limit is not particularly limited. The O content can be greater than 0%, but if the O content is less than 0.0002%, the production efficiency decreases. Therefore, the O content is preferably 0.0002% or more.

[0079] Total of Ti and Nb: 0.040 - 0.200%

[0080] Ti and Nb are the most important elements in the present invention, and are elements necessary to obtain excellent strength, toughness, and hydrogen-induced cracking resistance after post-heating by generating fine precipitates such as TiC and NbC after post-heating. If the total content of Ti and Nb is less than 0.040%, such effects cannot be obtained sufficiently, and excellent strength after post-heating cannot be obtained. On the other hand, if the total content of Ti and Nb is greater than 0.200%, the amount of coarse precipitates containing Ti and Nb increases, resulting in a decrease in hydrogen-induced cracking resistance after post-heating. In addition, the precipitates after post-heating become excessive, and excellent toughness after post-heating cannot be obtained. Therefore, the total content of Ti and Nb is 0.040 to 0.200%. The total content of Ti and Nb is preferably 0.050% or more, and more preferably 0.060% or more. In addition, the total content of Ti and Nb is preferably 0.160% or less, and more preferably 0.120% or less. It should be noted that as long as the total content of Ti and Nb satisfies the above range, the content of either one can be 0%.

[0081] The above components are the basic components of the high-strength hot-rolled steel sheet of the present invention. The high-strength hot-rolled steel sheet of the present invention can have a composition containing the above components and the balance being Fe and inevitable impurities.

[0082] In addition to the above components, the high-strength hot-rolled steel sheet of the present invention may contain one or more selected from Cr: 0.005 to 2.0%, Ni: 0.005 to 2.0%, Mo: 0.005 to 1.0%, V: 0.005 to 0.5%, B: 0.0002 to 0.0050%, Ca: 0.0001 to 0.0050%, REM: 0.0001 to 0.0050%, Cu: 0.005 to 0.5%, Sb: 0.0010 to 0.10%, Sn: 0.0010 to 0.10%.

[0083] Cr: 0.005 to 2.0%

[0084] Cr is an element effective for suppressing ferrite and forming lower bainite and martensite. In order to obtain such effects, when Cr is contained, the Cr content is preferably 0.005% or more. On the other hand, if the content of Cr is greater than 2.0%, the decrease in corrosion resistance sometimes becomes significant. Therefore, when Cr is contained, the Cr content is preferably 2.0% or less. The Cr content is more preferably 0.1% or more. In addition, the Cr content is more preferably 0.8% or less.

[0085] Ni: 0.005 to 2.0%

[0086] Ni is an element effective in suppressing ferrite and forming lower bainite and martensite. To achieve such an effect, when Ni is contained, it is preferable to make the Ni content 0.005% or more. On the other hand, if the Ni content is greater than 2.0%, a large amount of residual γ is formed, sometimes resulting in a reduction in toughness after post-heating. Therefore, when Ni is contained, it is preferable to make the Ni content 2.0% or less. The Ni content is more preferably 0.05% or more. In addition, the Ni content is more preferably 0.8% or less, and further preferably 0.5% or less.

[0087] Mo: 0.005 to 1.0%

[0088] Mo is an element effective in improving the hardenability of the steel plate and forming lower bainite and martensite. To achieve such an effect, when Mo is contained, it is preferable to make the Mo content 0.005% or more. On the other hand, if the Mo content is greater than 1.0%, the formation of Mo-based precipitates becomes significant, sometimes resulting in a reduction in toughness after post-heating. Therefore, when Mo is contained, it is preferable to make the Mo content 1.0% or less. The Mo content is more preferably 0.05% or more. In addition, the Mo content is more preferably 0.50% or less.

[0089] V: 0.005 to 0.5%

[0090] V is an element effective in improving the hardenability of the steel plate and forming lower bainite and martensite. To achieve such an effect, when V is contained, it is preferable to make the V content 0.005% or more. On the other hand, if the V content is greater than 0.5%, the formation of V-based precipitates becomes excessive, sometimes resulting in a reduction in toughness after post-heating. Therefore, when V is contained, it is preferable to make the V content 0.5% or less. The V content is more preferably 0.01% or more. In addition, the V content is more preferably 0.1% or less.

[0091] B: 0.0002 to 0.0050%

[0092] B is an element effective in improving the hardenability of the steel plate and forming lower bainite and martensite. To achieve such an effect, when B is contained, it is preferable to make the B content 0.0002% or more. On the other hand, if the B content is greater than 0.0050%, the B-based compounds increase, sometimes resulting in a reduction in toughness and delayed fracture resistance characteristics after post-heating. Therefore, when B is contained, it is preferable to make the B content 0.0050% or less. The B content is more preferably 0.0005% or more. In addition, the B content is more preferably 0.0040% or less.

[0093] Ca: 0.0001 to 0.0050%, REM: 0.0001 to 0.0050%

[0094] Ca and REM (rare earth elements) are elements effective in improving toughness and stress corrosion cracking resistance after post-heating by controlling the morphology of inclusions. To obtain such effects, when Ca and REM are contained, it is preferable that the respective contents are 0.0001% or more. On the other hand, if the contents of Ca and REM are greater than 0.0050% respectively, the influence of the increase in the amount of inclusions becomes excessive, and sometimes the toughness and stress corrosion cracking resistance after post-heating decrease. Therefore, when Ca and REM are contained, the contents of Ca and REM are preferably 0.0050% or less respectively. The Ca content is more preferably 0.0005% or more. In addition, the Ca content is more preferably 0.0030% or less. The REM content is more preferably 0.0005% or more. In addition, the REM content is more preferably 0.0030% or less. It should be noted that REM is a general term for 15 elements including Sc, Y, and lanthanum (La) with atomic number 57 to lutetium (Lu) with atomic number 71, and the REM content referred to here is the total content of these elements.

[0095] Cu: 0.005 - 0.5%, Sb: 0.0010 - 0.10%, Sn: 0.0010 - 0.10%

[0096] Cu, Sb, and Sn are elements effective in retarding corrosion reaction and improving stress corrosion cracking resistance after post-heating. To obtain such effects, when Cu, Sb, and Sn are contained, it is preferably that the Cu content is 0.005% or more, the Sb content is 0.0010% or more, and the Sn content is 0.0010% or more respectively. On the other hand, if the content of Cu is greater than 0.5%, the formation of Cu precipitates becomes excessive, and sometimes the toughness after post-heating decreases. Therefore, when Cu is contained, it is preferable that the Cu content is 0.5% or less. In addition, if the respective contents of Sb and Sn are greater than 0.10%, the grain boundary embrittlement effect becomes excessive, and sometimes the stress corrosion cracking resistance decreases. Therefore, when Sb and Sn are contained, it is preferable that the contents of Sb and Sn are 0.10% or less respectively. The Cu content is more preferably 0.05% or more. In addition, the Cu content is more preferably 0.3% or less. The Sb content is more preferably 0.0050% or more. In addition, the Sb content is more preferably 0.050% or less. The Sn content is more preferably 0.0050% or more. In addition, the Sn content is more preferably 0.050% or less.

[0097] It should be noted that even if the contents of Cr, Ni, Mo, V, B, Ca, REM, Cu, Sb, and Sn are less than the above lower limit values, the effects of the present invention are not impaired. Therefore, when the contents of these components are less than the above lower limit values, these elements are regarded as being contained as inevitable impurities. In addition, in the present invention, in addition to the above component composition, Mg, As, W, Ta, Pb, Zr, Hf, Te, Bi, and Se may be contained in a total amount of 0.3% or less by mass. It should be noted that the contents of these elements are preferably limited to 0.03% or less respectively.

[0098] Next, the steel structure of the high-strength hot-rolled steel sheet of the present invention will be described.

[0099] The steel structure of the high-strength hot-rolled steel sheet of the present invention has martensite and / or lower bainite as the main phase, and the residual γ is less than 3% by volume fraction.

[0100] Main phase: Martensite and / or lower bainite

[0101] In the present invention, in order to obtain high strength and excellent toughness and anti-delayed fracture properties after post-heating, a structure with martensite and / or lower bainite as the main phase is formed. If ferrite, pearlite, residual γ, etc. become the main phase, it is difficult to balance high strength and excellent toughness and anti-delayed fracture properties after post-heating. Therefore, the steel structure has martensite and / or lower bainite as the main phase. It should be noted that martensite can be auto-tempered martensite or tempered martensite, but does not include fresh martensite without carbides inside. In addition, lower bainite can be tempered lower bainite. It should be noted that in the present invention, the main phase refers to the phase accounting for 50% or more by area fraction. The area fraction of the main phase is preferably 60% or more, more preferably 75% or more. It should be noted that in the present invention, the main phase can be martensite, can be lower bainite, or can be the sum of martensite and lower bainite. There is no particular limitation on the upper limit of the area fraction of the main phase, and it can be 100%. As an example, the area fraction of the main phase can be less than 100%, or can be 98% or less.

[0102] Amount of retained austenite (retained γ): Less than 3%

[0103] Retained austenite (retained γ) is a structure that transforms into pearlite after post-heating and significantly reduces strength and toughness. Therefore, it is preferably reduced as much as possible. In the present invention, the retained γ is allowed to be less than 3% by volume fraction. Therefore, the retained γ is less than 3% by volume fraction. The retained γ is preferably less than 2% by volume fraction, more preferably less than 1% by volume fraction. There is no particular limitation on the lower limit of the volume fraction of the retained γ, and the volume fraction of the retained γ can be 0%.

[0104] It should be noted that, as phases other than martensite, lower bainite, and residual γ (other phases), one or more of ferrite, pearlite, and upper bainite can be cited. The total area ratio of the other phases is preferably 30% or less, more preferably 25% or less. The lower limit of the area ratio of the other phases is not particularly limited, and the total area ratio of the other phases can be 0%.

[0105] (Solid solution Ti amount + solid solution Nb amount) / (total Ti amount + total Nb amount): 0.300 or more and less than 0.800

[0106] If the ratio of the sum of the solid solution Ti amount and the solid solution Nb amount to the sum of the Ti content and the Nb content, that is, [(solid solution Ti amount + solid solution Nb amount) / (total Ti amount + total Nb amount)] is less than 0.300, the solid solution Ti amount and the solid solution Nb amount that become precipitates during post-heating and offset the strength reduction become insufficient. As a result, excellent strength after post-heating cannot be obtained. On the other hand, if it is 0.800 or more, the strength improvement due to the precipitation of precipitates after post-heating becomes excessive, and excellent toughness after post-heating cannot be obtained. Therefore, (solid solution Ti amount + solid solution Nb amount) / (total Ti amount + total Nb amount) is 0.300 or more and less than 0.800. It is preferably 0.350 or more. In addition, it is preferably 0.700 or less. It should be noted that (solid solution Ti amount + solid solution Nb amount) / (total Ti amount + total Nb amount) is obtained by the method described in the examples.

[0107] Total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more: 0.010 to 0.030 mass%

[0108] By containing a certain amount or more of Ti-containing precipitates and Nb-containing precipitates with a particle size of 100 nm or more, the growth of these precipitates competes with the precipitation of new TiC, NbC, etc. during post-heating. Thereby, the precipitation of fine TiC, NbC, etc. can be moderately suppressed, and excessive strength improvement and reduction of toughness can be suppressed. In order to obtain such an effect, it is necessary to make the total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more 0.010 mass% or more. On the other hand, if the total amount of the above Ti and Nb is more than 0.030 mass%, the reduction of toughness caused by coarse precipitates becomes significant. Therefore, it is necessary to make the total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more 0.030 mass% or less. Therefore, the total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more is 0.010 to 0.030 mass%. It is preferably 0.013 mass% or more. In addition, it is preferably 0.027 mass% or less. It should be noted that the total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more is obtained by the method described in the examples.

[0109] The pole density of the {110}<111> orientation in the surface layer region from the surface along the center direction of the plate thickness to 100 μm: 1.8 to 5.0

[0110] The surface layer region of the steel plate from the surface along the center direction of the plate thickness to 100 μm strongly affects the fracture formation during blanking and high-speed deformation. By controlling the pole density of the {110}<111> orientation in this region within the range of 1.8 to 5.0, excellent toughness can be obtained after post-heating. In addition, the fracture characteristics of blanking are good, and excellent anti-delayed fracture characteristics can be obtained after post-heating. To achieve such an effect, it is necessary to make the pole density of the {110}<111> orientation 1.8 or more in the surface layer region from the surface along the center direction of the plate thickness to 100 μm. On the other hand, if the above-mentioned pole density is greater than 5.0, the reduction in strength after post-heating becomes significant, and excellent strength (ΔHV of 50 or less) after post-heating cannot be obtained. Therefore, the pole density of the {110}<111> orientation is 1.8 to 5.0 in the surface layer region from the surface along the center direction of the plate thickness to 100 μm. It is preferably 2.0 or more. In addition, it is preferably 4.0 or less, and more preferably 3.0 or less. It should be noted that the pole density of the {110}<111> orientation in the surface layer region from the surface along the center direction of the plate thickness to 100 μm is obtained by the method described in the examples.

[0111] <Manufacturing method of high-strength hot-rolled steel plate>

[0112] The high-strength hot-rolled steel plate of the present invention is manufactured by the following operations: heating a slab having the above composition at a temperature in the range of 1150 to 1300 °C and holding it in this temperature range for 0.2 to 3.5 hours. Then, when performing hot rolling, the total reduction ratio in the temperature range of 1080 °C or higher is 80 to 90%, the total reduction ratio in the temperature range of 900 °C or lower is 20% or more, and the reduction ratio in one pass at a temperature of T (°C) or lower obtained by the following formula is 25% or less. After rolling, it is air-cooled for 1.0 s or more, and then cooled at an average cooling rate of 50 °C / s or more in the temperature range up to 550 °C. The time from reaching 550 °C to the start of rapid cooling is set to 0.5 to 4.0 s. Then, it is rapidly cooled at a cooling rate of 200 °C / s or more to a coiling temperature of 100 to 250 °C and coiled at the above coiling temperature.

[0113] T (°C) = 800 + 1000[Ti] + 2500[Nb]

[0114] Wherein, [Ti] and [Nb] are the contents (mass%) of Ti and Nb respectively, and are 0 when not contained.

[0115] In addition, the total reduction ratio in the temperature range above 1080°C is calculated based on the thickness of the slab before hot rolling and is obtained by the ratio of the thickness at 1080°C to the thickness at that time. In addition, the total reduction ratio in the temperature range below 900°C is calculated based on the thickness at 900°C and is obtained by the ratio of the thickness at 900°C to the final thickness. In addition, the reduction ratio in one pass below T (°C) is obtained by the ratio of the thickness before and after rolling in each pass below T (°C).

[0116] Hereinafter, a detailed description will be given. It should be noted that the above temperature is the temperature on the surface at the center of the width of the steel plate, and the above average cooling rate and cooling rate are the average cooling rate and cooling rate on the surface at the center of the width of the steel plate, respectively. In addition, unless otherwise specified, the average cooling rate is [(cooling start temperature - cooling stop temperature) / cooling time from the cooling start temperature to the cooling stop temperature].

[0117] Heating temperature of slab: 1150 - 1300°C

[0118] If the heating temperature of the slab is less than 1150°C, the dissolution of Ti-containing precipitates becomes insufficient. As a result, a value of 0.300 or more and less than 0.800 for (dissolved Ti amount + dissolved Nb amount) / (total Ti amount + total Nb amount) and a value of 0.010 - 0.030 mass% for the total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more cannot be obtained. On the other hand, if the heating temperature of the slab is greater than 1300°C, the dissolution of Ti-containing precipitates and Nb-containing precipitates becomes excessive. As a result, a value of 0.300 or more and less than 0.800 for (dissolved Ti amount + dissolved Nb amount) / (total Ti amount + total Nb amount) and a value of 0.010 - 0.030 mass% for the total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more cannot be obtained. Therefore, the heating temperature of the slab is 1150 - 1300°C. The above heating temperature is preferably 1170°C or more, more preferably 1185°C or more. In addition, the above heating temperature is preferably 1280°C or less, more preferably 1265°C or less.

[0119] Holding time in the temperature range of 1150 - 1300°C: 0.2 - 3.5 hours

[0120] If the holding time in the temperature range of 1150 to 1300 °C is less than 0.2 hours, the dissolution of Ti-containing precipitates and Nb-containing precipitates becomes insufficient. As a result, a value of 0.300 or more and less than 0.800 for (solid solution Ti amount + solid solution Nb amount) / (total Ti amount + total Nb amount) and a value of 0.010 to 0.030 mass% for the total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more cannot be obtained. On the other hand, if the holding time in the above temperature range is more than 3.5 hours, decarburization near the surface becomes significant, and ferrite, upper bainite, residual γ, etc. are likely to be generated on the surface, and the structure of the present invention cannot be obtained. Therefore, the holding time in the above temperature range of the slab is 0.2 to 3.5 hours. The above holding time is preferably 0.4 hours or more. In addition, the above holding time is preferably 2.5 hours or less.

[0121] Total reduction ratio in the temperature range of 1080 °C or higher: 80 to 90%

[0122] By performing rolling with a total reduction ratio of 80 to 90% in the temperature range of 1080 °C or higher, the generation and growth of coarse Ti-containing precipitates and Nb-containing precipitates with a particle size of 100 nm or more can be promoted. As a result, the total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more can be made 0.010 to 0.030 mass%. If the total reduction ratio is less than 80%, the generation of precipitates with a particle size of 100 nm or more becomes insufficient, and the total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more is less than 0.010 mass%. On the other hand, if the total reduction ratio is more than 90%, the generation of precipitates with a particle size of 100 nm or more becomes excessive, and the total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more is more than 0.030 mass%. Therefore, the total reduction ratio in the temperature range of 1080 °C or higher is 80 to 90%. The above total reduction ratio is preferably 81% or more. In addition, the above total reduction ratio is preferably 88% or less.

[0123] Total reduction ratio in the temperature range of 900 °C or lower is 20% or more

[0124] If the total reduction ratio in the temperature range of 900 °C or lower is less than 20%, strain-induced precipitation is suppressed, Ti-containing precipitates and Nb-containing precipitates are reduced, and a value of 0.300 or more and less than 0.800 for (solid solution Ti amount + solid solution Nb amount) / (total Ti amount + total Nb amount) cannot be obtained. Or the development of the texture in the surface layer part becomes insufficient, and a value of 1.8 to 5.0 for the pole density of the {110}<111> orientation in the surface layer region cannot be obtained. Therefore, the total reduction ratio in the temperature range of 900 °C or lower is 20% or more. It should be noted that the upper limit of the above total reduction ratio is not particularly limited, but the above total reduction ratio is preferably 80% or less, more preferably 60% or less.

[0125] Reduction ratio in one pass below T (°C): 25% or less

[0126] If reduction greater than 25% is carried out in one pass below T (°C) obtained by the following formula, strain-induced precipitation is promoted, Ti-containing precipitates and Nb-containing precipitates increase, and a value of 0.300 or more and less than 0.800 for (solid-solution Ti amount + solid-solution Nb amount) / (total Ti amount + total Nb amount) cannot be obtained. At the same time, the texture in the surface layer part is developed, and a value of 1.8 to 5.0 for the pole density of the {110}<111> orientation in the surface layer region cannot be obtained. Therefore, the reduction ratio in one pass below the above T (°C) is 25% or less. The above reduction ratio is preferably 20% or less, more preferably 18% or less. The lower limit of the above reduction ratio is not particularly limited, but if it is 5% or less, coarse grains may sometimes be generated, so the above reduction ratio is preferably greater than 5%. The above reduction ratio is more preferably 7% or more.

[0127] It should be noted that T (°C) is obtained by the following formula.

[0128] T (°C) = 800 + 1000[Ti] + 2500[Nb]

[0129] Wherein, [Ti] and [Nb] are the contents (mass%) of Ti and Nb respectively, and are not included when sometimes being 0.

[0130] Slow cooling for 1.0 s or more

[0131] By carrying out slow cooling after rolling under the above conditions, strain can be partially released, strain-induced precipitation and precipitation on dislocations during subsequent cooling can be suppressed, and Ti-containing precipitates and Nb-containing precipitates can be reduced. In order to obtain such an effect, the slow cooling time after rolling needs to be 1.0 s or more. The above slow cooling time is preferably 1.5 s or more, more preferably 2.0 s or more, and further preferably 2.2 s or more. The upper limit of the above slow cooling time is not particularly limited, but if the above slow cooling time is 5.0 s or less, it is easy to carry out subsequent hot rolling control, so the above slow cooling time is preferably 5.0 s or less. It should be noted that slow cooling means exposure to the atmosphere (air cooling) rather than active cooling (accelerated cooling) such as water injection. It should be noted that in the present invention, hot rolling includes rough rolling and finish rolling, and the slow cooling time after the above rolling is the slow cooling time after finish rolling, that is, after hot rolling.

[0132] Cool at an average cooling rate of 50°C / s or more in the temperature range up to 550°C

[0133] After the above-mentioned cooling, it is cooled at an average cooling rate of 50 °C / s or more in the temperature range up to 550 °C. If the average cooling rate up to 550 °C is less than 50 °C / s, it will lead to excessive formation of ferrite, upper bainite, Ti-containing precipitates, Nb-containing precipitates, etc., and hinder the formation of crystal orientation in the surface layer region. As a result, the phase structure, precipitates, and the value of the pole density of {110}<111> orientation in the surface layer region of the present invention cannot be obtained. Therefore, the average cooling rate in the temperature range from the cooling start temperature to 550 °C after the above-mentioned cooling is 50 °C / s or more. The above average cooling rate is preferably 70 °C / s or more. The upper limit of the above average cooling rate is not particularly limited, but if the above average cooling rate is 500 °C / s or more, it may sometimes lead to deterioration of the steel plate shape. Therefore, the above average cooling rate is preferably less than 500 °C / s, more preferably less than 200 °C / s.

[0134] Time from reaching 550 °C to the start of quenching: 0.5 to 4.0 s

[0135] By ensuring a certain time (leaving a certain time interval) during the period from reaching 550 °C to the start of quenching (quenching at a cooling rate of 200 °C / s or more described later), bainite in the intermediate temperature region near the surface can be formed. As a result, the pole density of {110}<111> orientation in the surface layer region of the present invention can be obtained. If the time from reaching 550 °C to the start of quenching is less than 0.5 s, such an effect cannot be fully obtained, and the value of the pole density of {110}<111> orientation in the surface layer region cannot be obtained. On the other hand, if the above time is greater than 4.0 s, upper bainite is excessively generated and the phase structure of the present invention cannot be obtained. Therefore, the time during the period from 550 °C to the start of quenching is 0.5 to 4.0 s. The above time is preferably 0.7 s or more. In addition, the above time is preferably 2.0 s or less, more preferably 1.6 s or less.

[0136] Cooling rate to the coiling temperature of 100 to 250 °C: 200 °C / s or more

[0137] As described above, after leaving a time of 0.5 to 4.0 s during the period from reaching 550°C to the start of rapid cooling, rapid cooling is started. If the cooling (rapid cooling) rate to a coiling temperature of 100 to 250°C is less than 200°C / s, upper bainite and retained γ are excessively generated, or the pole density of the {110}<111> orientation in the surface layer region increases. As a result, the phase structure of the present invention and the pole density of the {110}<111> orientation in the surface layer region cannot be obtained. Therefore, the cooling rate to the coiling temperature is 200°C / s or more. The above cooling rate is preferably 250°C / s or more. It should be noted that the upper limit of the above cooling rate is not particularly limited, but from the viewpoints of shape stability and the like, the above cooling rate is preferably 1000°C / s or less, more preferably 500°C / s or less.

[0138] Coiling temperature: 100 to 250°C

[0139] By adjusting the coiling temperature in the range of 100 to 250°C, martensite and lower bainite can be moderately tempered, and other phases can be removed to obtain the structure of the present invention. If the coiling temperature is less than 100°C, such an effect cannot be obtained sufficiently, and excessive fresh martensite is generated and the structure of the present invention cannot be obtained. On the other hand, if the coiling temperature is greater than 250°C, the tempering of martensite and lower bainite becomes significant, and fresh martensite and retained γ are generated, and the structure of the present invention cannot be obtained. Therefore, the coiling temperature is 100 to 250°C. The coiling temperature is preferably 120°C or more. In addition, the coiling temperature is preferably 220°C or less.

[0140] There are no particular limitations other than the conditions of the above manufacturing method, but it is preferably manufactured by appropriately adjusting the following conditions. For example, from the viewpoint of reducing coarse grains that cause a decrease in workability, finish rolling is preferably 4 passes or more. In addition, after hot rolling, skin pass rolling may be performed for shape correction, adjustment of surface roughness, etc. When pickling is performed, it is preferably performed by immersing in an acid bath at 50 to 100°C multiple times.

[0141] The high-strength hot-rolled steel sheet of the present invention is excellent in strength, toughness, and hydrogen embrittlement resistance after post-heating. Here, as the heating temperature for post-heating, 400°C or more can be cited. In addition, the upper limit of the heating temperature for post-heating is not particularly limited, but as an example, the heating temperature for post-heating can be cited as 1150°C or less. The post-heating time (holding time at the above heating temperature) is not particularly limited, but as an example, it can be cited as greater than 0 s. In addition, the above heating time can be cited as 3600 s or less as an example.

[0142] Examples

[0143] The steel composed of the components shown in Table 1 is melted in a converter to form a slab, and then the slab is heated and hot-rolled under the conditions shown in Table 2 to manufacture a hot-rolled steel sheet (original steel sheet). Using the obtained hot-rolled steel sheet, the following test methods are carried out for microstructure observation, analysis of dissolved Ti, dissolved Nb, Ti-containing precipitates and Nb-containing precipitates, and evaluation of tensile properties. Furthermore, post-heating shown in Table 2 is performed on the above hot-rolled steel sheet, and using the hot-rolled steel sheet after post-heating, the following test methods are carried out for evaluation of hardness, toughness and stress corrosion cracking resistance. It is carried out under the conditions that the post-heating temperature is 400 °C or higher at which an improvement in stretch flangeability is observed, and the post-heating time is 3600 s or less from the viewpoint of productivity.

[0144] Microstructure observation

[0145] The area ratios of martensite and lower bainite refer to the proportions of the areas of respective microstructures in the observed area. For the area ratio of martensite, samples are cut from the obtained hot-rolled steel sheet. After grinding the cross-section parallel to the rolling direction with respect to the sheet thickness, it is etched with a 3% nitric acid ethanol etching solution, and three fields of view are respectively photographed at a magnification of 1500 times using SEM (scanning electron microscope) at the 1 / 4 position of the sheet thickness. Based on the image data of the obtained secondary electron images, the area ratios of respective microstructures are obtained using Image-Pro manufactured by Media Cybernetics, Inc., and the average area ratio of the three fields of view is taken as the area ratio of each microstructure. The determination of the microstructure can be carried out through general classification. For example, it can be determined as follows. In the image data, lower bainite is distinguished as black or dark gray or gray or light gray containing carbide with consistent orientation. Martensite is a microstructure of black to light gray containing regular but multiple-orientation carbides. Or it is observed as white or light gray without carbide. Retained austenite is observed as white or light gray without carbide. Since it is sometimes impossible to distinguish a part of martensite and retained austenite, the retained austenite is obtained by the method described later, and the area ratio of martensite is obtained by subtracting it from the total area ratio of martensite and retained austenite obtained from the SEM image. It should be noted that the stronger the tempering degree of the microstructure, the darker the contrast of the matrix image. Therefore, based on the color of the above-mentioned matrix, in the present invention, the amount of carbide, the microstructure morphology, etc. are comprehensively judged and classified into a certain microstructure with similar characteristics including the microstructures described later. Carbide is white dot-shaped or linear. In addition, as microstructures other than the above, ferrite is a microstructure that is black or dark gray and does not have lower microstructures such as carbide and lath inside, and pearlite can be distinguished as a layered or partially discontinuous layered microstructure of black and white. In addition, upper bainite can be distinguished as a microstructure that is black or dark gray and has lower microstructures such as carbide and lath inside. The amount of retained γ is obtained as follows. After grinding the hot-rolled steel sheet to 1 / 4 + 0.1 mm of the sheet thickness, it is further ground by 0.1 mm through chemical polishing, and the obtained surface is used as the measurement surface. For the above-mentioned measurement surface, using the Kα1 ray of Mo in an X-ray diffractometer, 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 (ferrite) are measured. Then, the volume ratio is obtained from the intensity ratio of the integrated reflection intensities of respective planes from fcc iron and the integrated reflection intensities of respective planes from bcc iron, and this is taken as the amount of retained γ.

[0146] Table 3 shows the microstructures that constitute the main phase with an area ratio of 50% or more of each obtained microstructure and other microstructures. It should be noted that M in Table 3 refers to martensite, LB refers to lower bainite, γ refers to retained austenite, and O refers to other phases. Other phases include one or more of ferrite, pearlite, and upper bainite.

[0147] The pole density of the {110}<111> orientation in the surface layer region from the surface along the central direction of the plate thickness to 100 μm

[0148] Samples were cut out from the obtained hot-rolled steel sheet, the cross-section of the plate thickness parallel to the rolling direction was polished, and the strain was further removed by electrolytic polishing. Then, for the surface layer region from the surface along the central direction of the plate thickness to 100 μm, crystal orientation data were obtained by the EBSD (electron backscatter diffraction) method. The measurement area was 100 μm × 100 μm, the acceleration voltage was 30 kV, and the step size was 100 nm. Three fields of view were measured for each sample. The obtained data were analyzed using OIM Analysis Ver. 7.3.0 manufactured by TSL SOLUTIONS. For the obtained data, after removing the data with a Confidence Index (CI) value of 0.1 or less, calculate The Orientation Distribution Function (ODF) with Φ = 85 - 90° and a Resolution of 5° for each, calculate the pole density of this region, and take the average value as the pole density of the {110}<111> orientation for each field of view. Then, take the average value of the pole densities of the three fields of view of each sample as the pole density of each sample.

[0149] Analysis of dissolved Ti, dissolved Nb, Ti-containing precipitates, and Nb-containing precipitates

[0150] Specimens with a width of 30 mm and a length of 30 mm were collected from the obtained hot-rolled steel sheet and subjected to constant current electrolysis in a non-aqueous solvent-based electrolyte (10% AA-based electrolyte: 10 vol% acetylacetone - 1 mass% tetramethylammonium chloride - methanol). The current density was 20 mA / cm 2, the electrolysis amount was about 0.2 g. The electrolyzed electrolyte was used as an analytical solution, and the in-liquid concentrations (mass%) of Ti, Nb, and Fe as a comparative element were measured by ICP mass spectrometry. Based on the obtained concentrations, the concentration ratios of Ti and Nb to Fe were calculated, and then multiplied by the content (mass%) of Fe in the test piece, thereby obtaining the amount of dissolved Ti (mass%) and the amount of dissolved Nb (mass%). It should be noted that the content (mass%) of Fe in the test piece was obtained by subtracting the total of the component contents other than Fe from 100 mass%. Using the obtained amount of dissolved Ti (mass%) and the amount of dissolved Nb (mass%), the ratio of the total of the amount of dissolved Ti (mass%) and the amount of dissolved Nb (mass%) to the total of the amount of contained Ti (mass%) and the amount of contained Nb (mass%) was calculated. On the other hand, the test piece with precipitates attached to the surface after electrolysis was taken out of the electrolyte and immersed in a sodium hexametaphosphate aqueous solution (500 mg / L) (hereinafter referred to as SHMP aqueous solution). Then, ultrasonic vibration was applied to peel the precipitates from the test piece and extract them into the SHMP aqueous solution. Next, the SHMP aqueous solution containing the precipitates was filtered through a filter with a pore size of 100 nm, and then the precipitates trapped on the 100 nm filter were acid-decomposed, and the decomposition solution was analyzed using an ICP emission spectroscopic analyzer to measure the absolute values of Ti and Nb in the decomposition solution. The obtained absolute values of Ti and Nb were divided by the electrolysis mass to obtain the amounts of Ti and Nb contained in the precipitates with a particle size of 100 nm or more (mass% when the total composition of the test piece is 100 mass%). Next, the total of the obtained amounts of Ti (mass%) and Nb (mass%) was divided by the total of the amount of contained Ti (mass%) and the amount of contained Nb (mass%) in the test piece to obtain the total of the amount of Ti (mass%) present as the Ti-containing precipitate with a particle size of 100 nm or more and the amount of Nb (mass%) present as the Nb-containing precipitate with a particle size of 100 nm or more. It should be noted that the electrolysis mass was obtained by measuring the mass of the test piece after the precipitates were peeled off and subtracting it from the mass of the test piece before electrolysis.

[0151] Tensile test

[0152] From the obtained hot-rolled steel sheet, a JIS No. 5 tensile test piece (JIS Z2241: 2011) was collected in the direction parallel to the rolling direction, and a tensile test was conducted according to JIS Z 2241: 2011 with a strain rate of 10 －3 / s to obtain the TS. It should be noted that in the present invention, a TS of 1180 MPa or more was rated as qualified.

[0153] Vickers hardness test

[0154] Samples were cut from the obtained hot-rolled steel plate and the hot-rolled steel plate after post-heating. After grinding the plate thickness section parallel to the rolling direction, a Vickers hardness test was carried out at the 1 / 4 position of the plate thickness with a load of 5 kg and 5 measurement points. The average (arithmetic mean) was taken as the Vickers hardness of the steel plate. When the difference in hardness (ΔHV) before and after post-heating was 50 or less, it was judged that the strength after post-heating was excellent and rated as qualified.

[0155] Charpy impact test

[0156] Test pieces with a width of 10 mm and a length of 55 mm were collected from the hot-rolled steel plate that had been subjected to post-heating treatment, and Charpy impact test pieces with a V-notch having a front end angle of 45°, a front end radius of 0.25 mm, and a depth of 2 mm were made. Then, according to JIS Z 2242:2018, 5 Charpy impact tests were carried out at -20 °C to evaluate the ductile fracture rate. When the average value of the 5 ductile fracture rates was 50% or more, it was judged that the toughness after post-heating was excellent and rated as qualified. It should be noted that the plate thickness was 2.9 mm and the notch direction was parallel to the rolling direction.

[0157] Delayed fracture test

[0158] Test pieces with a width of 30 mm and a length of 110 mm were collected from the obtained hot-rolled steel plate, and the post-heating treatment shown in Table 2 was carried out to obtain test pieces. They were subjected to 90° V-bending processing with a bending radius of 15 mm with the ridge line parallel to the rolling direction, and the part opened due to springback was fastened with bolts, and then immersed in hydrochloric acid with a pH of 3 for 96 hours to investigate the presence or absence of cracks. Test pieces without cracks were judged to have excellent delayed fracture resistance characteristics after post-heating and rated as qualified. It should be noted that the end faces of the test pieces were formed by shearing with a shear angle of 1° and a gap of 10%, and the burrs were on the outer side of the bend. It should be noted that the "delayed fracture time (hr)" in Table 3 indicates the time when cracks occurred in the test pieces. However, the "96" in the above "delayed fracture time (hr)" means that no cracks occurred in the test pieces after the 96-hour immersion.

[0159] [Table 1]

[0160]

[0161] ※ The underlined part indicates outside the scope of the present invention.

[0162] ※ The remaining part other than the above component composition is Fe and unavoidable impurities.

[0163]

[0164]

[0165] The invention examples all have a TS of 1180 MPa or more, and the strength, toughness, and stress corrosion cracking resistance after post-heating are excellent. On the other hand, the comparative examples outside the scope of the present invention do not have the desired strength (TS) or do not achieve any one or more of the desired strength, toughness, and stress corrosion cracking resistance after post-heating.

[0166] Industrial applicability

[0167] According to the present invention, a high-strength hot-rolled steel sheet having a TS of 1180 MPa or more and less than 1600 MPa and excellent strength, toughness, and stress corrosion cracking resistance after post-heating can be obtained. When the high-strength steel sheet of the present invention is used for automotive part applications, it can greatly contribute to improving the crash safety and fuel efficiency of automobiles.

Claims

1. A high-strength hot-rolled steel sheet having the following composition: Containing by mass% C：0.06～0.23％、 Si: 0.1 - 3.0%, Mn: 1.5 - 3.5%, P: greater than 0% and 0.050% or less, S: greater than 0% and 0.0050% or less, Al: greater than 0% and 1.5% or less, N: greater than 0% and 0.010% or less, O: greater than 0% and 0.003% or less, And further containing a total of 0.040 - 0.200% of Ti and Nb, The balance being composed of Fe and inevitable impurities; And the steel structure is mainly composed of martensite and / or lower bainite, and the retained austenite is less than 3% by volume fraction, The ratio of the total amount of solid-solution Ti and solid-solution Nb to the total amount of Ti content and Nb content, i.e., (solid-solution Ti amount + solid-solution Nb amount) / (total Ti amount + total Nb amount), is 0.300 or more and less than 0.800, The total amount of Ti and Nb present as precipitates with a particle size of 100 nm or more is 0.010 - 0.030 mass%, In the surface layer region from the surface along the plate thickness center direction to 100 μm, the pole density of the {110}<111> orientation is 1.8 - 5.

0.

2. The high-strength hot-rolled steel sheet according to claim 1, wherein, The above composition further contains by mass% selected from Cr：0.005～2.0％、 Ni: 0.005 - 2.0%, Mo: 0.005 - 1.0%, V：0.005～0.5％、 B：0.0002～0.0050％、 Ca: 0.0001 - 0.0050%, REM: 0.0001 - 0.0050%, Cu: 0.005 - 0.5%, Sb: 0.0010 - 0.10%, and Sn: 0.0010 - 0.10% Of one or more.

3. A method for manufacturing a high-strength hot-rolled steel sheet, which is a method for manufacturing the high-strength hot-rolled steel sheet according to claim 1 or 2, The slab having the above composition is heated in a temperature range of 1150 - 1300 °C and held in this temperature range for 0.2 - 3.5 hours, Then, when performing hot rolling, The total reduction ratio in the temperature range above 1080 °C is 80 - 90%, the total reduction ratio in the temperature range below 900 °C is 20% or more, and the reduction ratio in one pass below T obtained by the following formula is 25% or less. After rolling, it is air-cooled for 1.0 s or more, Then, it is cooled at an average cooling rate of 50 °C / s or more in the temperature range up to 550 °C, and the time from reaching 550 °C to the start of quenching is set to 0.5 - 4.0 s, Then, it is quenched at a cooling rate of 200 °C / s or more to a coiling temperature of 100 - 250 °C and coiled at the coiling temperature, T = 800 + 1000[Ti] + 2500[Nb] Among them, [Ti] and [Nb] are the contents of Ti and Nb by mass%, respectively, and are 0 when not contained, and the unit of T is °C.

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