HIGH-STRENGTH HOT ROLLED STEEL SHEET AND METHOD FOR PRODUCING THE SAME

Description

CROSS REFERENCE TO RELATED APPLICATIONS

This is the U.S. National Phase application of PCT/JP2023/041535, filed Nov. 17, 2023 which claims priority to Japanese Patent Application No. 2022-186156, filed Nov. 22, 2022, the disclosures of these applications being incorporated herein by reference in their entireties for all purposes.

FIELD OF THE INVENTION

The present invention relates to a high-strength hot rolled steel sheet and a method for producing the same and more particularly to a high-strength hot rolled steel sheet suitable for a material of automotive parts and a method for producing the same.

BACKGROUND OF THE INVENTION

From the viewpoint of improving crash safety of automobiles and their fuel economy, there is a need to increase the strength of hot rolled steel sheets used for automotive parts. However, with hot rolled steel sheets with increased strength, the occurrence of cracking due to insufficient workability is significant during pressing, and it is therefore necessary to improve the pressing method and the workability of the steel sheets. It has been attempted to improve the workability by heating a steel sheet (base sheet) used for the pressing method. In the present description, heat treatment performed when a steel sheet (base sheet) is processed into a part is referred to also as post-heating. As for the steel sheets, development in consideration of the characteristics of parts subjected to post-heating and a working process (a post-heating working process) is being carried out. One problem with high-strength steel sheets having a tensile strength (TS) of 1180 MPa or more is delayed fracture. Therefore, it is necessary to design the materials in consideration of delayed fracture resistance after post-heating.

Patent Literature 1 discloses a technique for a method in which working temperature (heating temperature of post-heating) is set to 400 to 1000° C. to improve stretch flangeability. Patent Literature 2 discloses a technique for a high-strength hot rolled steel sheet having a TS of 730 MPa or more. The hot rolled steel sheet disclosed in Patent Literature 2 has a microstructure in which bainite is a main phase and the amount of solute Ti with respect to the total amount of Ti is 80% or more. In this manner, heat treatment hardenability is obtained in which the amount of increase in YS (yield strength) and the amount of increase in TS after heat treatment including heating the steel sheet to a temperature range of 500° C. to the Acl transformation point and then holding the steel sheet for 60 minutes are 100 MPa or more. Patent Literature 3 discloses a technique for a hot rolled steel sheet having a TS of 120 kgf/mm² or more and having good delayed fracture resistance.

PATENT LITERATURE

• PTL 1: Japanese Unexamined Patent Application Publication No. 2002-113527
• PTL 2: Japanese Unexamined Patent Application Publication No. 2015-57514
• PTL 3: Japanese Unexamined Patent Application Publication No. 6-145894

SUMMARY OF THE INVENTION

However, in Patent Literature 1, no consideration is given to the performance of parts such as strength, toughness, etc. after the post-heating, and there is room for improvement. In particular, when the post-heating is performed at a high temperature exceeding 400° C., the steel microstructure is largely changed. In this case, when the strength of the steel sheet (base sheet) before the post-heating is high, i.e., 1180 MPa or more, the influence on the strength is significant. Therefore, it is necessary to design the material in consideration of the influence of the post-heating on the strength. In the steel sheet disclosed in Patent Literature 2, although no reduction in strength occurs even after the heat treatment, the heat treatment causes an excessive increase in the strength, and the toughness after the heat treatment is insufficient due to significant precipitation of fine carbides, so that there is room for improvement. In Patent Literature 2, the level of the strength of the steel sheet under consideration is limited to about 980 MPa, and there is no information or suggestion about a high strength of 1180 MPa or more. Patent Literature 3 aims to improve the workability of the base sheet to thereby improve the delayed fracture resistance without post-heating. In Patent Literature 3, deep drawing is used for evaluation in the delayed fracture test, and no consideration is given to the resistance to delayed fracture from more severe end faces, i.e., end faces processed after post-heating, so that there is room for improvement.

Aspects of the present invention have been made in view of the foregoing circumstances, and it is an object to provide a high-strength hot rolled steel sheet exhibiting high strength, good toughness, and good delayed fracture resistance after post-heating and to provide a method for producing the high-strength hot rolled steel sheet.

The inventors have focused attention on the behavior of precipitation of Ti and Nb after post-heating of hot rolled steel sheets and have achieved the improvement in the properties of the steel sheets after heating (after post-heating) by controlling the initial amounts of coarse Ti-containing precipitates, coarse Nb-containing precipitates, solute Ti, and solute Nb before post-heating. The inventors have further focused attention on the orientations of crystals and have arrived at the idea that, by forming a microstructure in which surface layer regions are aligned in a specific orientation, the delayed fracture of punched end faces after post-heating is suppressed. The inventors have thus found that a hot rolled steel sheet (base sheet) that has high strength, can maintain the strength close to that of the base sheet even after post-heating, and exhibits good toughness and delayed fracture resistance after post-heating can be obtained by adjusting the chemical components, forming martensite and/or lower bainite as a main phase, adjusting the volume fraction of retained austenite (retained γ) to less than 3%, adjusting the pole density of the {110}<111> orientation in a surface layer region extending 100 μm from the surface in a direction toward the thicknesswise center to 1.8 to 5.0, adjusting the ratio of the sum of the amount of solute Ti and the amount of solute Nb to the sum of the amount of Ti contained and the amount of Nb contained, i.e., (the amount of solute Ti+the amount of solute Nb)/(the total amount of Ti+the total amount of Nb), to 0.300 or more and less than 0.800, and adjusting the total amount of Ti and Nb present as precipitates having a diameter of 100 nm or more to 0.010 to 0.030% by mass. Thus, aspects of the invention has been completed.

In accordance with aspects of the present invention, the high strength means a tensile strength (TS) of 1180 MPa or more and less than 1600 MPa.

In accordance with aspects of the present invention, the phrase “the strength after post-heating is high” means that the reduction in the strength of a hot rolled steel sheet after post-heating with respect to the strength of the hot rolled steel sheet (base sheet) before post-heating is 50 or less in terms of Vickers hardness.

In accordance with aspects of the present invention, the phrase “the toughness after post-heating is good” means that, in a Charpy impact test using a test specimen cut from a hot rolled steel sheet after post-heating, the percent ductile fracture at −20° C. is 50% or more. The thickness of the test specimen is 0.6 to 3.0 mm. When the thickness of the hot rolled steel sheet exceeds 3.0 mm, the front and back surfaces of a test specimen cut from the hot rolled steel sheet are ground, and the resulting test specimen is used for the Charpy impact test.

In accordance with aspects of the present invention, the phrase “the delayed fracture resistance after post-heating is good” means that, when a strip-shaped test specimen cut from a steel sheet after post-heating and having sheared end faces is subjected to 90° V-bending, tightened with a bolt etc. by the amount of springback, and immersed in hydrochloric acid with a pH of 3 for 96 hours, no cracking occurs.

In accordance with aspects of the present invention, the post-heating means heat treatment in which a hot rolled steel sheet (base sheet) is heated to 400° C. or higher.

The present invention includes the following aspects.

[1]A high-strength hot rolled steel sheet having a chemical composition containing, in % by mass,

• • C: 0.06 to 0.23%,
  • Si: 0.1 to 3.0%,
  • Mn: 1.5 to 3.5%,
  • P: more than 0% and 0.050% or less,
  • S: more than 0% and 0.0050% or less,
  • Al: more than 0% and 1.5% or less,
  • N: more than 0% and 0.010% or less, and
  • O: more than 0% and 0.003% or less and
  • further containing Ti and Nb in a total amount of 0.040 to 0.200%,
  • with the balance being Fe and incidental impurities,
  • wherein the high-strength hot rolled steel sheet has a steel microstructure in which martensite and/or lower bainite is a main phase and in which a volume fraction of retained austenite is less than 3%,
  • wherein a ratio of a sum of an amount of solute Ti and an amount of solute Nb to a sum of an amount of Ti contained and an amount of Nb contained, which is given by (the amount of solute Ti+the amount of solute Nb)/(a total amount of Ti+a total amount of Nb), is 0.300 or more and less than 0.800,
  • wherein a total amount of Ti and Nb present as precipitates having a diameter of 100 nm or more is 0.010 to 0.030% by mass, and
  • wherein, in a surface layer region extending 100 μm from a surface in a direction toward a thicknesswise center, a pole density of a {110}<111> orientation is 1.8 to 5.0.

[2] The high-strength hot rolled steel sheet according to [1], wherein the chemical composition further contains, in % by mass, at least one element 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%, and
  • Sn: 0.0010 to 0.10%.

[3]A method for producing the high-strength hot rolled steel sheet according to [1] or [2], the method including:

• • heating a slab having the chemical composition in a temperature range of 1150 to 1300° C.; holding the slab in the temperature range for 0.2 to 3.5 hours;
  • then hot-rolling the slab to obtain a steel sheet under the conditions that a total rolling reduction in a temperature range of 1080° C. or higher is 80 to 90%, that a total rolling reduction in a temperature range of 900° C. or lower is 20% or more, and that a rolling reduction per pass at a temperature lower than or equal to T (° C.) determined from a formula below is 25% or less; then allowing the steel sheet to naturally cool for 1.0 s or longer,
  • then cooling the resulting steel sheet in a temperature range down to 550° C. at an average cooling rate of 50° C./s or more; setting a time from when the temperature of the steel sheet reaches 550° C. to a start of rapid cooling to 0.5 to 4.0 s;
  • then rapidly cooling the resulting steel sheet to a coiling temperature of 100 to 250° C. at a cooling rate of 200° C./s or more; and coiling the resulting steel sheet at the coiling temperature:

T ⁢ ( ° ⁢ C . ) = 800 + 1000 [ Ti ] + 2500 [ Nb ] ,

• • where [Ti] and [Nb] are a content (% by mass) of Ti and a content (% by mass) of Nb, respectively, and are each 0 when a corresponding element is not contained.

Aspects of the present invention can provide a high-strength hot rolled steel sheet exhibiting high strength, good toughness, good delayed fracture resistance after post-heating and can provide a method for producing the high-strength hot rolled steel sheet.

The high-strength hot rolled steel sheet obtained in accordance with aspects of the invention is suitable for a material of automotive parts and exhibits high strength, good toughness, and good delayed fracture resistance after post-heating or a post-heating working process.

With the high-strength hot rolled steel sheet according to aspects of the invention, high-strength products such as high-strength automotive parts can be obtained which exhibit high strength, good toughness, and good delayed fracture resistance even after heat treatment for improving workability and improving fatigue properties.

DETAILED DESCRIPTION OF EMBODIMENTS OF THE INVENTION

The details of embodiments of the high-strength hot rolled steel sheet according to aspects of the invention and a method for producing the same will be described. However, the invention is not limited to the following embodiments.

<High-Strength Hot Rolled Steel Sheet>

The high-strength hot rolled steel sheet according to aspects of the invention may be an as-hot-rolled steel sheet with mill scale or a so-called pickled hot rolled steel sheet subjected to pickling after the hot rolling. Preferably, the high-strength hot rolled steel sheet to be obtained in accordance with aspects of the invention has a thickness of 0.6 mm or more. Preferably, the high-strength hot rolled steel sheet according to aspects of the invention has a thickness of 10.0 mm or less. When the high-strength hot rolled steel sheet according to aspects of the invention is used as a material of automotive parts, the thickness is more preferably 1.0 mm or more. When the high-strength hot rolled steel sheet according to aspects of the invention is used as a material of automotive parts, the thickness is more preferably 6.0 mm or less. The width of the high-strength hot rolled steel sheet according to aspects of the invention is preferably 500 mm or more and more preferably 700 mm or more. The width of the high-strength hot rolled steel sheet according to aspects of the invention is preferably 1800 mm or less and more preferably 1400 mm or less.

The high-strength hot rolled steel sheet according to aspects of the invention has a specific chemical composition and a specific steel microstructure. The chemical composition and the steel microstructure will be described in this order.

First, the chemical composition of the high-strength hot rolled steel sheet according to aspects of the invention will be described. “%” representing a content in the chemical composition means “% by mass.”

The high-strength hot rolled steel sheet according to aspects of the invention has a chemical composition containing, in % 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%), and Ti and Nb in a total amount of 0.040 to 0.200%, with the balance being Fe and incidental impurities.

C: 0.06 to 0.23%

C is an element that increases the TS through the formation and strengthening of martensite and lower bainite, is bonded with Ti, Nb, N, etc. to form precipitates, and is thereby effective in preventing a reduction in the strength after post-heating. If the content of C is less than 0.06%, these effects are not obtained, and the TS of the steel sheet (base sheet) is not 1180 MPa or more, or the steel sheet does not exhibit high strength after post-heating. If the content of C exceeds 0.23%, the reduction in toughness after post-heating is significant, and the steel sheet does not exhibit good toughness after post-heating. Therefore, the content of C is 0.06 to 0.23%. The content of C is preferably 0.07% or more. The content of C is preferably 0.22% or less and more preferably 0.20% or less.

Si: 0.1 to 3.0%

Si is an element effective in solid solution strengthening of the steel and in preventing a reduction in the strength after post-heating. To obtain these effects, the content of Si must be 0.1% or more. If the content of Si exceeds 3.0%, the amount of polygonal ferrite formed is excessively large, so that the steel microstructure in accordance with aspects of the invention is not obtained. Therefore, the content of Si is 0.1 to 3.0%. The content of Si is preferably 0.2% or more. The content of Si is preferably 2.0% or less and more preferably 1.5% or less.

Mn: 1.5 to 3.5%

Mn is an element effective in suppressing the formation of ferrite and upper bainite to thereby allow the formation of lower bainite and martensite. If the content of Mn is less than 1.5%, this effect is not obtained sufficiently, and polygonal ferrite, upper bainite, etc. are formed, so that the microstructure in accordance with aspects of the invention is not obtained. If the content of Mn exceeds 3.5%, the toughness and delayed fracture resistance decrease significantly, and the toughness and delayed fracture resistance after post-heating are not good. Therefore, the content of Mn is 1.5 to 3.5%. The content of Mn is preferably 1.6% or more. The content of Mn is preferably 3.0% or less and more preferably 2.5% or less.

P: More than 0% and 0.050% or Less

P causes deterioration in toughness and delayed fracture resistance after post-heating, and it is therefore preferable to reduce the amount of P as much as possible. In accordance with aspects of the present invention, the allowable content of P is 0.050%. Therefore, the content of P is 0.050% or less. The content of P is preferably 0.030% or less. No particular limitation is imposed on the lower limit, and the content of P may be more than 0%. However, if the content of P is less than 0.001%, the production efficiency is low. Therefore, the content of P is preferably 0.001% or more.

S: More than 0% and 0.0050% or Less

S causes deterioration in toughness and delayed fracture resistance after post-heating, and it is therefore preferable to reduce the amount of S as much as possible. In accordance with aspects of the present invention, the allowable content of S is 0.0050%. Therefore, the content of S is 0.0050% or less. The content of S is preferably 0.0030% or less, more preferably 0.0020% or less, and still more preferably 0.0015%. No particular limitation is imposed on the lower limit, and the content of S may be more than 0%. However, if the content of S is less than 0.0002%, the production efficiency is low. Therefore, the content of S is preferably 0.0002% or more.

Al: More than 0% and 1.5% or Less

Al serves as a deoxidizing agent, and it is preferable to add Al in a deoxidization step. The content of Al may be more than 0%. However, from the viewpoint of using Al as a deoxidizing agent, the content of Al is preferably 0.01% or more. If a large amount of Al is contained, a large amount of polygonal ferrite is formed, and the steel microstructure in accordance with aspects of the invention is not obtained. In accordance with aspects of the present invention, the allowable content of Al is 1.5%. Therefore, the content of Al is 1.5% or less. The content of Al is preferably 0.50% or less and more preferably 0.20% or less.

N: More than 0% and 0.010% or Less

N forms TiN and NbC and inhibits the precipitation of fine TiC, NbC, etc., and it is therefore preferable to reduce the amount of N as much as possible. In accordance with aspects of the present invention, the allowable content of N is 0.010%. Therefore, the content of N is 0.010% or less. The content of N is preferably 0.007% or less. No particular limitation is imposed on the lower limit, and the content of N may be more than 0%. However, if the content of N is less than 0.0005%, the production efficiency is low. Therefore, the content of N is preferably 0.0005% or more.

O: More than 0% and 0.003% or Less

O causes deterioration in toughness and delayed fracture resistance after post-heating, and it is therefore preferable to reduce the amount of O as much as possible. In accordance with aspects of the present invention, the allowable content of O is 0.003%. Therefore, the content of O is 0.003% or less. The content of O is preferably 0.002% or less. No particular limitation is imposed on the lower limit, and the content of O may be more than 0%. However, if the content of O is less than 0.0002%, the production efficiency is low. Therefore, the content of O is preferably 0.0002% or more.

Total Amount of Ti and Nb: 0.040 to 0.200%

Ti and Nb are the most important elements in accordance with aspects of the invention and are elements necessary to appropriately form fine precipitates such as TiC and NbC after post-heating to thereby obtain high strength, good toughness, and good delayed fracture resistance after post-heating. If the total content of Ti and Nb is less than 0.040%, these effects are not obtained sufficiently, and the strength after post-heating is not high. If the total content of Ti and Nb exceeds 0.200%, the amount of coarse precipitates containing Ti and Nb increases. In this case, the delayed fracture resistance after post-heating deteriorates, and the amount of precipitates after post-heating is excessively large, so that the steel sheet does not exhibit good toughness after post-heating. 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. The total content of Ti and Nb is preferably 0.160% or less and more preferably 0.120% or less. It is only necessary that the total content of Ti and Nb be within the above range, and the content of one of them may be 0%.

The above components are the basic components of the high-strength hot rolled steel sheet according to aspects of the invention. The high-strength hot rolled steel sheet according to aspects of the invention may have a chemical composition containing the above components with the balance being Fe and incidental impurities.

The high-strength hot rolled steel sheet according to aspects of the invention may further contain, in addition to the components described above, at least one element 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%, and Sn: 0.0010 to 0.10%.

Cr: 0.005 to 2.0%

Cr is an element effective in the formation of lower bainite and martensite while the formation of ferrite is suppressed. When Cr is contained, the content of Cr contained is preferably 0.005% or more in order to obtain the above effect. If the content of Cr exceeds 2.0%, a significant reduction in corrosion resistance may occur. Therefore, when Cr is contained, the content of Cr is preferably 2.0% or less. The content of Cr is more preferably 0.1% or more. The content of Cr is more preferably 0.8% or less.

Ni: 0.005 to 2.0%

Ni is an element effective in the formation of lower bainite and martensite while the formation of ferrite is suppressed. When Ni is contained, the content of Ni is preferably 0.005% or more in order to obtain the above effect. If the content of Ni exceeds 2.0%, a large amount of retained γ is formed, and this may cause deterioration in toughness after post-heating. Therefore, when Ni is contained, the content of Ni is preferably 2.0% or less. The content of Ni is more preferably 0.05% or more. The content of Ni is more preferably 0.8% or less and still more preferably 0.5% or less.

Mo: 0.005 to 1.0%

Mo is an element effective in increasing the hardenability of the steel sheet and in the formation of lower bainite and martensite. When Mo is contained, the content of Mo is preferably 0.005% or more in order to obtain these effects. If the content of Mo exceeds 1.0%, the formation of Mo-based precipitates is significant, and this may cause deterioration in the toughness after post-heating. Therefore, when Mo is contained, the content of Mo is preferably 1.0% or less. The content of Mo is more preferably 0.05% or more. The content of Mo is more preferably 0.50% or less.

V: 0.005 to 0.5%

V is an element effective in increasing the hardenability of the steel sheet and in the formation of lower bainite and martensite. When V is contained, the content of V is preferably 0.005% or more in order to obtain these effects. If the content of V exceeds 0.5%, an excessively large number of V-based precipitates are formed, and this may cause deterioration in the toughness after post-heating. Therefore, when V is contained, the content of V is preferably 0.5% or less. The content of V is more preferably 0.01% or more. The content of V is more preferably 0.1% or less.

B: 0.0002 to 0.0050%

B is an element effective in increasing the hardenability of the steel sheet and in the formation of lower bainite and martensite. When B is contained, the content of B is preferably 0.0002% or more in order to obtain these effects. If the content of B exceeds 0.0050%, the amount of B-based compounds increases, and this may cause deterioration in the toughness and delayed fracture resistance after post-heating. Therefore, when B is contained, the content of B is preferably 0.0050% or less. The content of B is more preferably 0.0005% or more. The content of B is more preferably 0.0040% or less.

Ca: 0.0001 to 0.0050% and REM: 0.0001 to 0.0050%

Ca and REM (rare earth metal) are elements effective in improving the toughness and delayed fracture resistance after post-heating through shape control of inclusions. When Ca and REM are contained, their contents are each preferably 0.0001% or more in order to obtain the above effect. If the contents of Ca and REM each exceed 0.0050%, the influence of the increase in the amount of inclusions is excessively large, and this may cause deterioration in the toughness and delayed fracture resistance after post-heating. Therefore, when Ca and REM are contained, the contents of Ca and REM are each preferably 0.0050% or less. The content of Ca is more preferably 0.0005% or more. The content of Ca is more preferably 0.0030% or less. The content of REM is more preferably 0.0005% or more. The content of REM is more preferably 0.0030% or less. The REM is a generic term for Sc, Y, and 15 elements from lanthanum (La) with an atomic number of 57 to lutetium (Lu) with an atomic number of 71. The content of REM mentioned here is the total content of these elements.

Cu: 0.005 to 0.5%, Sb: 0.0010 to 0.10%, and Sn: 0.0010 to 0.10%

Cu, Sb, and Sn are each an element effective in delaying a corrosion reaction and in improving the delayed fracture resistance after post-heating. When Cu, Sb, and Sn are contained, it is preferable that the content of Cu is 0.005% or more, that the content of Sb is 0.0010% or more, and that the content of Sn is 0.0010% or more, respectively, in order to obtain these effects. If the content of Cu exceeds 0.5%, an excessively large number of Cu precipitates are formed, and this may cause deterioration in the toughness after post-heating. Therefore, when Cu is contained, the content of Cu is preferably 0.5% or less. If the contents of Sb and Sn each exceed 0.10%, the grain boundary embrittlement effect is excessively high, and this may cause deterioration in the delayed fracture resistance. Therefore, when Sb and Sn are contained, the contents of Sb and Sn are each preferably 0.10% or less. The content of Cu is more preferably 0.05% or more. The content of Cu is more preferably 0.3% or less. The content of Sb is more preferably 0.0050% or more. The content of Sb is more preferably 0.050% or less. The content of Sn is more preferably 0.0050% or more. The content of Sn is more preferably 0.050% or less.

Even when the contents of Cr, Ni, Mo, V, B, Ca, REM, Cu, Sb, and Sn are lower than the respective lower limits, the effects according to aspects of the invention are not impaired. Therefore, when the contents of these components are lower than the above lower limits, these elements are regarded as the content of incidental impurities. In accordance with aspects of the present invention, the steel sheet may contain, in addition to the chemical composition described above, in % by mass, one or two or more of Mg, As, W, Ta, Pb, Zr, Hf, Te, Bi, and Se in the range of a total amount of 0.3% or less. Preferably, the content of each of these elements is limited to 0.03% or less.

Next, the steel microstructure of the high-strength hot rolled steel sheet according to aspects of the invention will be described.

The high-strength hot rolled steel sheet according to aspects of the invention has a steel microstructure including martensite and/or lower bainite as a main phase and further including retained γ at a volume fraction of less than 3%.

Main Phase: Martensite and/or Lower Bainite

In accordance with aspects of the present invention, to obtain high strength and to obtain good toughness and delayed fracture resistance after post-heating, the microstructure includes martensite and/or lower bainite as a main phase. If ferrite, pearlite, retained γ, etc. is a main phase, it is difficult to achieve high strength and good toughness and delayed fracture resistance after post-heating simultaneously. Therefore, the steel microstructure includes martensite and/or lower bainite as a main phase. The martensite may be auto-tempered martensite or tempered martensite, but fresh martensite containing no carbides inside is excluded. The lower bainite may be tempered lower bainite. In accordance with aspects of the present invention, the main phase is a phase whose area fraction is 50% or more. The area fraction of the main phase is preferably 60% or more and more preferably 75% or more. In accordance with aspects of the present invention, the main phase may be martensite or may be lower bainite. The main phase may be a combination of martensite and lower bainite. No particular limitation is imposed on the upper limit of the area fraction of the main phase, and the area fraction may be 100%. The area fraction of the main phase may be, for example, less than 100% or 98% or less.

Amount of Retained Austenite (Retained γ): Less than 3%

Retained austenite (retained γ) is a microstructure that transforms to pearlite after post-heating, causing significant deterioration in strength and toughness. Therefore, it is preferable to reduce the amount of retained austenite as much as possible. In accordance with aspects of the present invention, the allowable volume fraction of retained γ is less than 3%. Therefore, the volume fraction of the retained γ is less than 3%. The volume fraction of the retained γ is preferably less than 2% and more preferably less than 1%. No particular limitation is imposed on the lower limit of the volume fraction of the retained γ, and the volume fraction of the retained γ may be 0%.

Phases other than martensite, lower bainite, and retained γ (the other phases) include one or two or more of ferrite, pearlite, and upper bainite. The total area fraction of the other phases is preferably 30% or less and more preferably 25% or less. No particular limitation is imposed on the lower limit of the area fraction of the other phases, and the total area fraction of the other phases may be 0%.

(Amount of Solute Ti+Amount of Solute Nb)/(Total Amount of Ti+Total Amount of Nb): 0.300 or More and Less than 0.800

If the ratio of the sum of the amount of solute Ti and the amount of solute Nb to the sum of the amount of Ti contained and the amount of Nb contained, i.e., [(the amount of solute Ti+the amount of solute Nb)/(the total amount of Ti+the total amount of Nb)], is less than 0.300, the amounts of solute Ti and solute Nb that form precipitates during the post-heating are insufficient to compensate for the reduction in strength. In this case, the strength after post-heating is not good. If the ratio is 0.800 or more, the increase in strength due to precipitation of precipitates after post-heating is excessively large, so that the steel sheet does not exhibit good toughness after post-heating. Therefore, (the amount of solute Ti+the amount of solute Nb)/(the total amount of Ti+the total amount of Nb) is 0.300 or more and less than 0.800. The ratio is preferably 0.350 or more. The ratio is preferably 0.700 or less. (The amount of solute Ti+the amount of solute Nb)/(the total amount of Ti+the total amount of Nb) can be determined by a method described in Examples.

Total Amount of Ti and Nb Present as Precipitates Having Diameter of 100 nm or More: 0.010 to 0.030% by Mass

When Ti-containing precipitates and Nb-containing precipitates having a diameter of 100 nm or more are contained in a certain amount or more, the growth of these precipitates competes with new precipitation of TiC, NbC, etc. during post-heating. In this case, the precipitation of fine TiC, NbC, etc. is appropriately suppressed, and an excessive increase in strength and an excessive reduction in toughness can be prevented. To obtain these effects, the total amount of Ti and Nb present as precipitates having a diameter of 100 nm or more must be 0.010% by mass or more. If the total amount of Ti and Nb exceeds 0.030% by mass, the reduction in toughness due to coarse precipitates is significant. Therefore, the total amount of Ti and Nb present as precipitates having a diameter of 100 nm or more must be 0.030% by mass or less. Thus, the total amount of Ti and Nb present as precipitates having a diameter of 100 nm or more is 0.010 to 0.030% by mass. The total amount is preferably 0.013% by mass or more. The total amount is preferably 0.027% by mass or less. The total amount of Ti and Nb present as precipitates having a diameter of 100 nm or more can be determined by a method described in Examples.

Pole Density of {110}<111> Orientation in Surface Layer Region Extending 100 μm from Surface in Direction Toward Thicknesswise Center: 1.8 to 5.0

A surface layer region extending 100 μm from the surface of the steel sheet in a direction toward the thicknesswise center has a strong influence on the formation of a fracture surface during punching 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, good toughness can be obtained after post-heating. Moreover, the fracture surface formed by punching is well shaped, and the delayed fracture resistance after post-heating is good. To obtain these effects, the pole density of the {110}<111> orientation in the surface layer region extending 100 μm from the surface in the direction toward the thicknesswise center must be 1.8 or more. If the pole density exceeds 5.0, the reduction in strength after post-heating is significant, and good strength (a ΔHV of 50 or less) is not obtained after post-heating. Therefore, the pole density of the {110}<111> orientation in the surface layer region extending 100 μm from the surface in the direction toward the thicknesswise center is 1.8 to 5.0. The pole density is preferably 2.0 or more. The pole density is preferably 4.0 or less and more preferably 3.0 or less. The pole density of the {110}<111> orientation in the surface layer region extending 100 μm from the surface in the direction toward the thicknesswise center can be determined by a method described in Examples.

<Method for Producing High-Strength Hot Rolled Steel Sheet>

The high-strength hot rolled steel sheet according to aspects of the invention is produced as follows. A slab having the chemical composition described above is heated in a temperature range of 1150 to 1300° C., held in this temperature range for 0.2 to 3.5 hours, and hot-rolled under the conditions that the total rolling reduction in a temperature range of 1080° C. or higher is 80 to 90%, that the total rolling reduction in a temperature range of 900° C. or lower is 20% or more, and that the rolling reduction per pass at T (° C.) or lower determined from a formula below is 25% or less. Then the steel sheet is allowed to naturally cool for 1.0 s or longer and then cooled in a temperature range down to 550° C. at an average cooling rate of 50° C./s or more. Then the time from when the temperature of the steel sheet reaches 550° C. to the start of rapid cooling is set to 0.5 to 4.0 s, and the steel sheet is rapidly cooled to a coiling temperature of 100 to 250° C. at a cooling rate of 200° C./s or more and coiled at the coiling temperature.

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

Here, [Ti] and [Nb] are the content (% by mass) of Ti and the content (% by mass) of Nb, respectively, and are each 0 when the corresponding element is not contained.

The total rolling reduction in the temperature range of 1080° C. or higher is determined from the ratio of the thickness at 1080° C. with respect to the thickness of the slab before the hot rolling. The total rolling reduction in the temperature range of 900° C. or lower is determined from the ratio of the final thickness of the sheet with respect to the thickness at 900° C. The rolling reduction per pass at T (° C.) or lower is determined from the ratio of the thickness after a rolling pass to the thickness before the rolling pass at T (° C.) or lower.

A detailed description will be given below. The temperature described above is the temperature of the surface of the steel sheet at the widthwise center, and the average cooling rate and the cooling rate described above are the average cooling rate and the cooling rate, respectively, at the surface of the steel sheet at the widthwise center. The average cooling rate is [(cooling start temperature−cooling stop temperature)/cooling time from cooling start temperature to cooling stop temperature], unless otherwise specified.

Heating Temperature of Slab: 1150 to 1300° C.

If the heating temperature of the slab is lower than 1150° C., the Ti-containing precipitates do not dissolve sufficiently. In this case, the value of (the amount of solute Ti+the amount of solute Nb)/(the total amount of Ti+the total amount of Nb) does not fall within the range of 0.300 or more and less than 0.800, and the total amount of Ti and Nb present as precipitates having a diameter of 100 nm or more does not fall within the range of 0.010 to 0.030% by mass. If the heating temperature of the slab exceeds 1300° C., the Ti-containing precipitates and Nb-containing precipitates dissolve excessively. In this case, the value of (the amount of solute Ti+the amount of solute Nb)/(the total amount of Ti+the total amount of Nb) does not fall within the range of 0.300 or more and less than 0.800, and the total amount of Ti and Nb present as precipitates having a diameter of 100 nm or more does not fall within the range of 0.010 to 0.030% by mass. Therefore, the heating temperature of the slab is 1150 to 1300° C. The heating temperature is preferably 1170° C. or higher and more preferably 1185° C. or higher. The heating temperature is preferably 1280° C. or lower and more preferably 1265° C. or lower.

Holding Time in Temperature Range of 1150 to 1300° C.: 0.2 to 3.5 Hours

If the holing time in the temperature range of 1150 to 1300° C. is shorter than 0.2 hours, the Ti-containing precipitates and the Nb-containing precipitates do not dissolve sufficiently. In this case, the value of (the amount of solute Ti+the amount of solute Nb)/(the total amount of Ti+the total amount of Nb) does not fall within the range of 0.300 or more and less than 0.800, and the total amount of Ti and Nb present as precipitates having a diameter of 100 nm or more does not fall within the range of 0.010 to 0.030% by mass. If the holding time in the temperature range exceeds 3.5 hours, decarburization near the surface layer is significant, and ferrite, upper bainite, retained γ, etc. are likely to be formed in the surface layer, so that the microstructure in accordance with aspects of the invention is not obtained. Therefore, the holding time of the slab in the temperature range is 0.2 to 3.5 hours. The holding time is preferably 0.4 hours or longer. The holding time is preferably 2.5 hours or shorter.

Total Rolling Reduction in Temperature Range of 1080° C. Or Higher: 80 to 90%

By performing rolling in the temperature range of 1080° C. or higher at a total rolling reduction of 80 to 90%, the nucleation and growth of coarse Ti-containing precipitates and coarse Nb-containing precipitates larger than or equal to 100 nm can be facilitated. In this manner, the total amount of Ti and Nb present as precipitates having a diameter of 100 nm or more can be adjusted to 0.010 to 0.030% by mass. If the total rolling reduction is less than 80%, the formation of the precipitates having a diameter of 100 nm or more is insufficient, and the total amount of Ti and Nb present as precipitates having a diameter of 100 nm or more is less than 0.010% by mass. If the total rolling reduction exceeds 90%, the precipitates having a diameter of 100 nm or more are formed excessively, and the total amount of Ti and Nb present as the precipitates having a diameter of 100 nm or more is more than 0.030% by mass. Therefore, the total rolling reduction in the temperature range of 1080° C. or higher is 80 to 90%. The total rolling reduction is preferably 81% or more. The total rolling reduction is preferably 88% or less.

Total Rolling Reduction in Temperature Range of 900° C. Or Lower: 20% or More

If the total rolling reduction at 900° C. or lower is less than 20%, strain-induced precipitation is suppressed, and the amount of the Ti-containing precipitates and the Nb-containing precipitates is reduced, so that the (the amount of solute Ti+the amount of solute Nb)/(the total amount of Ti+the total amount of Nb) does not fall within the range of 0.300 or more and less than 0.800. In some cases, the development of the texture in the surface layer portion is insufficient, and the value of the pole density of the {110}<111> orientation in the surface layer region does not fall within the range of 1.8 to 5.0. Therefore, the total rolling reduction in the temperature range of 900° C. or lower is 20% or more. No particular limitation is imposed on the upper limit of the total rolling reduction. However, the total rolling reduction is preferably 80% or less and more preferably 60% or less.

Rolling Reduction Per Pass at T (° C.) or Lower: 25% or Less

When the rolling is performed at lower than or equal to T (° C.) determined by a formula below at a rolling reduction of more than 25% per pass, strain induced precipitation is facilitated, and the amount of the Ti-containing precipitates and Nb-containing precipitates increases, so that the value of (the amount of solute Ti+the amount of solute Nb)/(the total amount of Ti+the total amount of Nb) does not fall within the range of 0.300 or more and less than 0.800. At the same time, the texture in the surface layer portion develops, and the pole density of the {110}<111> orientation in the surface layer region does not fall within the range of 1.8 to 5.0. Therefore, the rolling reduction per pass at T (° C.) or lower is 25% or less. The rolling reduction is preferably 20% or less and more preferably 18% or less. No particular limitation is imposed on the lower limit of the rolling reduction. However, if the rolling reduction is 5% or less, coarse grains may be formed. Therefore, the rolling reduction is preferably more than 5%. The rolling reduction is more preferably 7% or more.

T (° C.) is determined from the following formula.

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

Here, [Ti] and [Nb] are the content (% by mass) of Ti and the content (% by mass) of Nb, respectively, and are each 0 when the corresponding element is not contained.

Natural Cooling for 1.0 s or Longer

After the rolling under the conditions described above, the steel sheet is allowed to naturally cool. In this manner, strain is partially released, and the occurrence of strain induced precipitation and precipitation on dislocations during the subsequent cooling is suppressed, so that the amount of the Ti-containing precipitates and the Nb-containing precipitates can be reduced. To obtain this effect, the natural cooling time after the rolling must be 1.0 s or longer. The natural cooling time is preferably 1.5 s or longer, more preferably 2.0 s or longer, and still more preferably 2.2 s or longer. No particular limitation is imposed on the upper limit of the natural cooling time. However, when the natal cooling time is 5.0 s or shorter, the subsequent hot rolling can be easily controlled. Therefore, the natural cooling time is preferably 5.0 s or shorter. The natural cooling means that the steel sheet is exposed to the atmosphere (air cooling) without active cooling (accelerated cooling) such as injection of water. In accordance with aspects of the present invention, the hot rolling includes rough rolling and finish rolling, and the natural cooling time after the rolling is a cooling time after the hot rolling, i.e., after the finish rolling.

Cooling in Temperature Range Down to 550° C. At Average Cooling Rate of 50° C./s or More

After the natural cooling, the steel sheet is cooled in a temperature range down to 550° C. at an average cooling rate of 50° C./s or more. If the average cooling rate to 550° C. is less than 50° C./s, excessively large amounts of ferrite, upper bainite, Ti-containing precipitates, Nb-containing precipitates, etc. are formed, and the orientations of crystals formed in the surface layer region are disturbed. In this case, the phase structure and the precipitates in accordance with aspects of the invention are not obtained, and the pole density of the {110}<111> orientation in the surface layer region does not fall within the range of 1.8 to 5.0. Therefore, the average cooling rate in the temperature range from the cooling start temperature down to 550° C. after the natural cooling is 50° C./s or more. The average cooling rate is preferably 70° C./s or more. No particular limitation is imposed on the upper limit of the average cooling rate. However, if the average cooling rate is 500° C./s or more, the shape of the steel sheet may deteriorate. Therefore, the average cooling rate is preferably less than 500° C./s and more preferably less than 200° C./s.

Time from when Temperature Reaches 550° C. To Start of Rapid Cooling: 0.5 to 4.0 s

By providing a certain time period (a certain time interval) between the time when the temperature reaches 550° C. and the time at which rapid cooling (rapid cooling at a cooling rate of 200° C./s or more described later) is started, medium temperature bainite can be formed near the surface layer. In this case, the pole density of the {110}<111> orientation in accordance with aspects of the invention can be obtained in the surface layer region. If the time from when the temperature reaches 550° C. to when the rapid cooing is started is shorter than 0.5 s, the above effect is not obtained sufficiently, and the pole density of the {110}<111> orientation in the surface layer region does not fall within the range of 1.8 to 5.0. If the above time exceeds 4.0 s, an excessively large amount of upper bainite is formed, and the phase structure in accordance with aspects of the invention is not obtained. Therefore, the time from 550° C. to the start of rapid cooling is 0.5 to 4.0 s. The time is preferably 0.7 s or longer. The time is preferably 2.0 s or shorter and more preferably 1.6 s or shorter.

Cooling Rate to Coiling Temperature of 100 to 250° C.: 200° C./s or More

The time from when the temperature reaches 550° C. to when the rapid cooling is started is set to 0.5 to 4.0 s as described above, and then the rapid cooling is started. If the cooling (rapid cooling) rate to the coiling temperature of 100 to 250° C. is less than 200° C./s, excessively large amounts of upper bainite and retained γ are formed, and the pole density of the {110}<111> orientation in the surface layer region increases. In this case, the phase structure in accordance with aspects of the invention and the pole density of the {110}<111> orientation in accordance with aspects of the invention are not obtained. Therefore, the cooling rate to the coiling temperature is 200° C./s or more. The cooling rate is preferably 250° C./s or more. No particular limitation is imposed on the upper limit of the cooling rate. However, from the viewpoint of shape stability etc., the cooling rate is preferably 1000° C./s or less and more preferably 500° C./s or less.

Coiling Temperature: 100 to 250° C.

By adjusting the coiling temperature within the range of 100 to 250° C., martensite and lower bainite can be appropriately tempered, and other phases are eliminated, so that the microstructure in accordance with aspects of the invention can be obtained. If the coiling temperature is lower than 100° C., the above effect is not obtained sufficiently, and an excessively large amount of fresh martensite is formed, so that the microstructure in accordance with aspects of the invention is not obtained. If the coiling temperature exceeds 250° C., martensite and lower bainite are tempered significantly, and fresh martensite and retained γ are formed, so that the microstructure in accordance with aspects of the invention is not obtained. Therefore, the coiling temperature is 100 to 250° C. The coiling temperature is preferably 120° C. or higher. The coiling temperature is preferably 220° C. or lower.

No particular limitation is imposed on the production method except for the conditions described above. However, it is preferable to produce the steel sheet by appropriately controlling conditions as described below. For example, it is preferable to perform the finish rolling in at least four passes, from the viewpoint of reducing the number of coarse grains that cause a reduction in workability etc. After the hot rolling, temper rolling may be performed for the purpose of shape correction and surface roughness adjustment. When pickling is performed, it is preferable to immerse the steel sheet a plurality of times in a pickling bath at 50 to 100° C.

The high-strength hot rolled steel sheet according to aspects of the invention exhibits high strength, good toughness, and good delayed fracture resistance after post-heating. The heating temperature of the post-heating may be 400° C. or higher. No particular limitation is imposed on the upper limit of the heating temperature of the post-heating. The heating temperature of the post-heating is, for example, 1150° C. or lower. No particular limitation is imposed on the heating time of the post-heating (the holding time at the heating temperature). The heating time is, for example, longer than 0 s. The heating time is, for example, 3600 s or shorter.

EXAMPLES

Steel material having chemical compositions shown in Table 1 was produced using converters and formed into slabs. Then the slabs were heated and hot-rolled under conditions shown in Table 2 to produce hot rolled steel sheets (base sheets). The hot rolled steel sheets obtained were used to perform microstructure observation, to analyze solute Ti, solute Nb, Ti-containing precipitates, and Nb-containing precipitates, and to evaluate tensile properties according to test methods described below. The hot rolled steel sheets were subjected to post-heating as shown in Table 2, and the hot rolled steel sheets subjected to the post-heating were used to evaluate hardness, toughness, and delayed fracture resistance according to test methods described below. The temperature of the post-heating was 400° C. or higher at which the improvement in the stretch flangeability was found, and the post-heating time was 3600 s or shorter from the viewpoint of productivity.

Microstructure Observation

The area fractions of martensite and lower bainite are the ratios of the areas of these microstructures in the observation area. The area fraction of martensite was determined as follows. A sample was cut from one of the obtained hot rolled steel sheets, and its thicknesswise cross section parallel to the rolling direction was polished and etched with 3% nital. Then, at a position ¼ of the thickness, photographs were taken in three viewing areas using an SEM (scanning electron microscope) at a magnification of 1500×. The image data of the obtained secondary electron images was used to determine the area fractions of the microstructures using Image-Pro available from Media Cybernetics, and the average area fractions in the three viewing areas were used as the area fractions of the microstructures. To identify the microstructures, a general classification method may be used, and the following method, for example, may be used for identification. In the image data, black, dark gray, gray, or light gray regions containing uniformly oriented carbides are identified as lower bainite. Martensite is observed as a black to light gray microstructure containing a plurality of regularly arranged carbides with different orientations or a white or light gray microstructure containing no carbides. Retained austenite is observed as white or light gray regions containing no carbide. Part of martensite and retained austenite cannot be distinguished from each other in some cases. Therefore, the area fraction of retained austenite was determined by a method described later, and the area fraction of retained austenite was subtracted from the total area fraction of martensite and retained austenite determined from the SEM images to thereby determine the area fraction of martensite. As the degree of tempering increases, the black contrast of the matrix in the microstructure image increases, and therefore the color of the matrix is merely a guide. In accordance with aspects of the present invention, the identification was made in a comprehensive manner based on the amount of carbides, the form of the microstructures, etc., and each of the microstructures including microstructures described later was classified into a microstructure having characteristics similar thereto. Carbides appear as white dots or lines. Microstructures other than those described above include the following microstructures. Ferrite is a black or dark gray microstructure containing no carbides and no substructures such as laths. Pearlite can be identified as a black and white lamellar or partially discontinuous lamellar microstructure. Upper bainite can be identified as a black or dark gray microstructure containing carbides and substructures such as laths. The amount of retained γ is determined as follows. A hot rolled steel sheet is ground to a position ¼ of the thickness+0.1 mm and then further polished by 0.1 mm by chemical polishing, and the resulting surface is used as a measurement surface. On the measurement surface, the MoKα1 line from an X-ray diffractometer is used to measure the integrated reflection intensities from {200}, {220}, and {311}planes of fcc iron (austenite) and {200}, {211}, and {220}planes of bcc iron (ferrite). The volume fraction of the retained γ is determined from the intensity ratios of the integrated reflection intensities from the above planes of fcc iron to the integrated reflection intensities from the above planes of bcc iron and is used as the amount of the retained γ.

Microstructures forming the main phase with an area fraction of 50% or more were determined using the obtained area fractions of the microstructures, and the main phase and the other microstructures are shown in Table 3. In Table 3, M represents martensite, and LB represents lower bainite. γ represents retained austenite, and O represents the other phases. The other phases include one or two or more of ferrite, pearlite, and upper bainite.

Pole Density of {110}<111> Orientation in Surface Layer Region Extending 100 μm from Surface in Direction Toward Thicknesswise Center

Samples were cut from one of the obtained hot rolled steel sheets, and the thicknesswise cross section of each sample parallel to the rolling direction was polished and then electropolished to remove strain. Then an EBSD (electron back-scatter diffraction) method was used to obtain crystal orientation data from a surface layer region extending 100 μm from the surface in a direction toward the thicknesswise center. The measurement region was a 100 μm×100 μm region. The acceleration voltage was 30 kV, and the step size of 100 nm. The measurement was performed at three viewing areas in each sample. The data obtained was analyzed using OIMAnalysis Ver. 7.3.0 available from TSL SOLUTIONS. In the obtained data, data values with a Confidence Index (CI) value of 0.1 or less were cut. The Orientation Distribution Function (ODF) was computed using φ1=30 to 40°, φ2=45°, and φ=85 to 90° with each of the Resolution set to 5°. The pole density of each of the regions was computed, and the average value was used as the pole density of the {110}<111> orientation in the viewing areas. The average of the pole densities in the three viewing areas in each of the sample was used as the pole density of the samples.

Analysis of Solute Ti, Solute Nb, Ti-Containing Precipitates, and Nb-Containing Precipitates

A test specimen having a width of 30 mm and a length of 30 mm was cut from one of the obtained hot rolled steel sheets 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 set to 20 mA/cm², and the amount of electrolysis was about 0.2 g. The electrolyte after electrolysis was used as an analysis solution, and the concentrations (% by mass) of Ti, Nb, and Fe serving as a comparison element in the solution were measured by ICP mass spectrometry. The concentration ratios of Ti and Nb with respect to Fe were computed based on the obtained concentrations and multiplied by the content (% by mass) of Fe in the test specimen to thereby obtain the amount (% by mass) of solute Ti and the amount (% by mass) of solute Nb. The content (% by mass) of Fe in the test specimen was determined by subtracting the total content (% by mass) of components other than Fe from 100% by mass. The amount (% by mass) of solute Ti and the amount (% by mass) of solute Nb were used to compute the ratio of the sum of the amount (% by mass) of solute Ti and the amount (% by mass) of solute Nb to the sum of the amount of (% by mass) Ti contained and the amount (% by mass) of Nb contained. After the electrolysis, the test specimen with precipitates adhering to the surface was removed from the electrolyte and immersed in an aqueous sodium hexametaphosphate solution (500 mg/L) (hereinafter referred to as an aqueous SHMP solution). Then ultrasonic vibrations were applied to the test specimen, and the precipitates were thereby separated from the test specimen and extracted into the aqueous SHMP solution. Next, the aqueous SHMP solution containing the precipitates was filtered using a filter with a pore size of 100 nm, and then the precipitates collected on the 100 nm filter were decomposed using an acid. The decomposition solution was analyzed using an IPC emission spectrometer 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 electrolyzed mass to obtain the amounts of Ti and Nb (% by mass with respect to 100% by mass of all the components in the test specimen) contained in precipitates having a diameter of 100 nm or more. Next, the sum of the obtained Ti and Nb amounts (% by mass) was divided by the sum of the amounts (% by mass) of Ti and Nb contained in the test specimen to obtain the sum of the amount (% by mass) of Ti present as Ti-containing precipitates with a diameter of 100 nm or more and the amount (% by mass) of Nb present as Nb-containing precipitates with a diameter of 100 nm or more. The electrolyzed mass was determined by measuring the mass of the test specimen from which the precipitates had been removed and subtracting the determined mass from the mass of the test specimen before the electrolysis.

Tensile Test

One of the obtained hot rolled steel sheets was cut in a direction parallel to the rolling direction to obtain a JIS No. 5 tensile test piece (JIS Z 2241: 2011), and a tensile test was performed at a strain rate of 10⁻³/s according to the specifications of JIS Z 2241:2011 to determine the TS. In accordance with aspects of the present invention, a steel sheet with a TS of 1180 MPa or more was rated pass.

Vickers Hardness Test

Samples were cut from the obtained hot rolled steel sheets and from the hot rolled steel sheets subjected to post-heating. A thicknesswise cross section of each sample parallel to the rolling direction was polished, and a Vickers hardness test with a load of 5 kg was performed at a position ¼ of the thickness. The measurement was performed at 5 points, and the average (arithmetic mean) was used as the Vickers hardness of the steel sheet. When the difference in hardness (ΔHV) before and after the post-heating was 50 or less, the strength after the post-heating was judged as good, and the steel sheet was rated pass.

Charpy Impact Test

The obtained hot rolled steel sheets were subjected to post-heating treatment. Test specimens having a width of 10 mm and a length of 55 mm were cut from one of the treated hot rolled steel sheets, and Charpy impact test specimens having a V-notch with an included angle of 45°, a tip radius of 0.25 mm, and a depth of 2 mm were produced. Then the Charpy impact test was performed five times at −20° C. according to JIS Z 2242: 2018 to evaluate the percent ductile fracture. When the average percent ductile fracture of the five measurements was 50% or more, the toughness after post-heating was judged as good, and the steel sheet was rated pass. The sheet thickness was 2.9 mm, and the direction of the notch was parallel to the rolling direction.

Delayed Fracture Test

A test specimen having a width of 30 mm and a length of 110 mm was cut from one of the obtained hot rolled steel sheets and subjected to post-heating treatment shown in Table 2 to thereby obtain a test specimen. The test specimen was subjected to 90° V-bending at a bending radius of 15 mm such that the ridge line was parallel to the rolling direction. The bent test specimen was tightened with a bolt by the amount of springback and immersed in hydrochloric acid with a pH of 3 for 96 hours, and the present or absence of cracking was checked. A test specimen with no cracking was judged to have good delayed fracture resistance after post-heating, and the steel sheet was rated pass. The end faces of the test specimen were formed by shearing with a shearing angle of 1° and a clearance of 10% such that the burrs were on the outer side of the bend to be formed. In Table 3, “Delayed fracture time (hr)” is the time at which cracking occurred in the test specimen. “96” in the “Delayed fracture time (hr)” column means that no cracking occurred in the test specimen even after immersion for 96 hours.

	TABLE 1
	Chemical composition (% by mass)

Steel	C	Si	Mn	P	S	Al	N	O	Ti	Nb	Ti + Nb	Others	Remarks
A	0.09	0.55	2.0	0.018	0.0013	0.033	0.006	0.001	0.072	—	0.072	B: 0.0009	Suitable
													steel
B	0.07	0.92	3.2	0.031	0.0022	0.033	0.002	0.003	0.148	—	0.148	—	Suitable
													steel
C	0.12	0.29	2.1	0.015	0.0029	0.037	0.002	0.001	0.051	0.013	0.064	Mo: 0.20	Suitable
													steel
D	0.15	0.66	1.5	0.009	0.0014	0.032	0.004	0.002	0.035	0.045	0.080	Ni: 0.60,	Suitable
												Cr: 0.20	steel
E	0.18	1.45	2.4	0.010	0.0004	0.024	0.005	0.002	0.025	0.089	0.114	Cu: 0.10,	Suitable
												Sn: 0.03	steel
F	0.21	1.88	2.5	0.011	0.0003	0.121	0.003	0.001	—	0.133	0.133	REM: 0.0020,	Suitable
												Sb: 0.020	steel
G	0.08	0.73	1.7	0.016	0.0021	0.035	0.003	0.001	0.082	—	0.082	Cr: 0.30,	Suitable
												B: 0.0015,	steel
												Sb: 0.010
H	0.05	0.38	2.3	0.019	0.0015	0.018	0.002	0.002	0.012	0.039	0.051	V: 0.15,	Comparative
												Ni: 0.30	steel
I	0.25	0.95	1.7	0.010	0.0008	0.066	0.004	0.002	0.068	0.022	0.090	B: 0.0012	Comparative
													steel
J	0.09	0.02	2.2	0.009	0.0013	0.028	0.003	0.001	0.052	0.030	0.082	Mo: 0.20,	Comparative
												Ca: 0.0020	steel
K	0.09	3.20	1.8	0.007	0.0010	0.030	0.003	0.003	0.006	0.080	0.086	Cr: 0.40,	Comparative
												Mo: 0.10	steel
L	0.08	0.77	1.2	0.025	0.0017	0.043	0.002	0.002	0.041	0.025	0.066	Ni: 0.20	Comparative
													steel
M	0.12	1.15	3.6	0.009	0.0020	0.049	0.004	0.001	0.093	—	0.093	—	Comparative
													steel
N	0.10	0.33	2.2	0.013	0.0006	0.030	0.003	0.001	0.031	0.006	0.037	Mo: 0.10,	Comparative
												B: 0.0022	steel
O	0.10	0.59	2.2	0.011	0.0008	0.029	0.004	0.001	0.178	0.025	0.203	V: 0.10	Comparative
													steel
* Underlines mean outside the range of the invention.
* The balance other than the above components is Fe and incidental impurities.

	TABLE 2
	Hot rolling

Heating of slab

Total

Total

Maximum

Natural

Average

		Slab		rolling	rolling	rolling	cooling	cooling
		heating	Slab	reduction at	reduction at	reduction per pass	time	rate down
Steel		temper-	heating	1080° C.	900° C.	at T° C.	after	to 550°
sheet		ature	time	or higher	or lower	or lower	rolling	C. (°
No.	Steel	(° C.)	(hr)	(%)	(%)	(%)	(s)	C./s)
1	A	1250	1.0	90	26	17	2.2	50
2		1250	1.0	90	26	17	2.2	50
3		1080	1.0	90	26	17	2.2	50
4	B	1200	1.2	82	31	24	1.5	60
5		1320	1.2	82	31	24	1.5	60
6		1200	1.2	77	34	25	1.5	60
7	C	1160	0.6	85	29	15	1.2	100
8		1160	0.6	85	29	15	0.8	100
9		1160	0.6	85	15	15	1.2	100
10	D	1280	1.0	83	31	24	1.2	80
11		1280	4.0	83	31	24	1.2	80
12		1280	1.0	83	31	24	1.2	80
13	E	1270	0.5	84	26	24	1.8	80
14		1270	0.5	84	26	24	1.8	30
15		1270	0.5	84	26	24	1.8	80
16	F	1290	2.5	87	34	23	1.2	60
17		1290	2.5	87	34	23	1.2	60
18		1290	2.5	87	34	31	1.2	60
19		1290	0.1	87	34	23	1.2	60
20	G	1230	1.0	87	45	20	1.6	60
21		1230	1.0	91	45	20	1.6	60
22		1230	1.0	87	45	20	1.6	60
23	H	1220	0.8	87	29	17	1.0	70
24	I	1200	1.5	87	28	23	1.2	60
25	J	1200	0.6	87	28	23	1.2	60
26	K	1200	1.0	87	29	23	1.2	60
27	L	1200	1.0	87	29	16	1.2	60
28	M	1200	1.0	87	29	16	1.2	60
29	N	1200	1.0	87	29	15	1.2	70
30	O	1280	1.0	87	29	24	1.2	70

Hot rolling

Time from

Cooling

Post-heating

			550° C.	rate to			Post-	Post-
			to start	coiling	Coiling		heating	heating
	Steel		of rapid	temper-	temper-		temper-	holding
	sheet		cooling	ature	ature	T*1	ature	time
	No.	Steel	(s)	(° C./s)	(° C.)	(° C.)	(° C.)	(s)	Remarks
	1	A	2.2	280	200	872	550	3600	Inventive
									Example
	2		2.2	280	80	872	550	3600	Comparative
									Example
	3		2.2	280	200	872	550	3600	Comparative
									Example
	4	B	2.0	280	200	948	700	1	Inventive
									Example
	5		2.0	280	200	948	700	1	Comparative
									Example
	6		2.0	280	200	948	700	1	Comparative
									Example
	7	C	2.2	250	120	884	780	180	Inventive
									Example
	8		2.2	250	120	884	780	180	Comparative
									Example
	9		2.2	250	120	884	780	180	Comparative
									Example
	10	D	2.2	250	240	948	850	10	Inventive
									Example
	11		2.2	250	240	948	850	10	Comparative
									Example
	12		0.3	250	240	948	450	10	Comparative
									Example
	13	E	1.2	200	180	1048	500	1	Inventive
									Example
	14		1.2	200	180	1048	500	1	Comparative
									Example
	15		1.2	200	330	1048	500	1	Comparative
									Example
	16	F	2.2	250	220	1133	450	1	Inventive
									Example
	17		5.0	250	220	1133	450	1	Comparative
									Example
	18		2.2	250	220	1133	450	1	Comparative
									Example
	19		2.2	250	220	1133	450	1	Comparative
									Example
	20	G	3.0	250	180	882	800	1	Inventive
									Example
	21		3.0	250	180	882	800	1	Comparative
									Example
	22		3.0	100	180	882	650	1	Comparative
									Example
	23	H	2.2	280	200	910	400	1	Comparative
									Example
	24	I	2.2	250	200	923	400	1	Comparative
									Example
	25	J	2.2	250	200	927	500	1	Comparative
									Example
	26	K	2.2	250	200	1006	750	10	Comparative
									Example
	27	L	2.2	250	200	904	600	1	Comparative
									Example
	28	M	2.2	250	200	893	400	600	Comparative
									Example
	29	N	2.2	280	200	846	600	1	Comparative
									Example
	30	O	2.2	280	200	1041	600	1	Comparative
									Example
	* Underlines mean outside the range of the invention.
	*1T(° C.) = 800 + 1000 [Ti] + 2500 [Nb]. In the formula, [Ti] is the content (% by mass) of Ti, and [Nb] is the content (% by mass) of Nb.

	TABLE 3
	Steel microstructure

	(Amount of	Total amount
	solute Ti +	of Ti and
	amount of	Nb present	Pole

solute Nb)/

as precipitates

density of

Properties after post-heating

				(total amount	equal to	{110}<111>			Percent	Delayed
	V(M + LB)	V(γ)	V(O)	of Ti +	or larger	orientation	Property		ductile	fracture
Steel	(area	(volume	(area	total amount	than 100 nm	in surface	TS		fracture	time
sheet No.	fraction)	fraction)	fraction)	of Nb)	(% by mass)	layer region	(MPa)	ΔHV	(%)	(hr)	Remarks

1	88	1	11	0.611	0.015	2.9	1209	7	95	96	Inventive
											Example
2	45	1	54	0.583	0.016	2.8	1268	19	45	56	Comparative
											Example
3	87	0	13	0.292	0.035	2.9	1217	60	40	96	Comparative
											Example
4	98	0	2	0.757	0.028	3.3	1184	18	100	96	Inventive
											Example
5	99	0	1	0.845	0.008	3.3	1182	33	5	96	Comparative
											Example
6	99	0	1	0.797	0.009	3.3	1189	25	15	96	Comparative
											Example
7	88	1	11	0.500	0.014	2.6	1311	38	70	96	Inventive
											Example
8	94	0	6	0.297	0.015	2.2	1320	57	70	96	Comparative
											Example
9	90	1	9	0.641	0.014	1.7	1303	42	45	72	Comparative
											Example
10	94	0	6	0.588	0.019	3.0	1342	5	70	96	Inventive
											Example
11	45	5	50	0.613	0.022	2.3	1183	35	30	24	Comparative
											Example
12	96	0	4	0.563	0.018	5.1	1355	51	70	96	Comparative
											Example
13	98	2	0	0.667	0.026	2.5	1512	32	80	96	Inventive
											Example
14	79	2	19	0.395	0.025	1.7	1468	29	40	32	Comparative
											Example
15	88	10	2	0.684	0.025	2.7	1446	52	20	96	Comparative
											Example
16	93	2	5	0.669	0.029	4.5	1554	48	60	96	Inventive
											Example
17	75	6	19	0.684	0.028	4.1	1502	45	15	96	Comparative
											Example
18	96	1	3	0.286	0.030	5.2	1557	60	30	48	Comparative
											Example
19	95	1	4	0.496	0.056	4.0	1540	49	30	96	Comparative
											Example
20	76	1	23	0.707	0.020	3.1	1220	21	95	96	Inventive
											Example
21	76	0	24	0.573	0.032	3.1	1225	37	40	96	Comparative
											Example
22	69	5	26	0.683	0.023	3.0	1182	51	30	96	Comparative
											Example
23	99	1	0	0.588	0.012	2.0	1145	23	100	96	Comparative
											Example
24	98	2	0	0.667	0.026	2.6	1535	47	0	96	Comparative
											Example
25	85	0	15	0.512	0.021	2.5	1214	57	90	96	Comparative
											Example
26	49	6	45	0.140	0.028	1.5	903	33	80	96	Comparative
											Example
27	21	2	77	0.061	0.022	2.1	720	2	100	96	Comparative
											Example
28	99	1	0	0.742	0.021	2.0	1296	34	0	3	Comparative
											Example
29	97	1	2	0.459	0.011	2.4	1254	52	90	96	Comparative
											Example
30	95	0	5	0.714	0.042	2.4	1262	9	30	52	Comparative
											Example
* Underlines mean outside the range of the invention.
* M: martensite, LB: lower bainite, γ: retained austenite, O: other phases

In all Inventive Examples, the TS was 1180 MPa or more, and the strength, toughness, and delayed fracture resistance after post-heating were good. However, in Comparative Examples outside the ranges of the invention, the desired strength (TS) was not obtained, or at least one of the desired strength, the desired toughness, and the desired delayed fracture resistance after post-heating was not obtained.

INDUSTRIAL APPLICABILITY

According to aspects of the present invention, a high-strength hot rolled steel sheet having a TS of 1180 MPa or more and less than 1600 MPa and exhibiting high strength, good toughness, and good delayed fracture resistance after post-heating can be obtained. When the high-strength steel sheet according to aspects of the invention is used for automotive part applications, the steel sheet can significantly contribute to improvement in crash safety of the automobiles and improvement in their fuel economy.