JP5861434B2 - High-strength hot-rolled steel sheet excellent in punchability and manufacturing method thereof - Google Patents

Description

  The present invention relates to a high-strength hot-rolled steel sheet having high tensile strength (TS): 900 MPa or more and excellent workability (particularly punchability) useful for the use of automobile members and a method for producing the same.

In recent years, from the viewpoint of global environmental conservation, the automobile industry as a whole has been aimed at improving the fuel efficiency of automobiles for the purpose of regulating CO ₂ emissions. In order to improve the fuel consumption of automobiles, it is most effective to reduce the weight of automobiles by thinning the members used. In recent years, the amount of high-strength hot-rolled steel sheets used as materials for automobile parts is increasing. On the other hand, many automotive parts made of steel sheets are formed by pressing, burring, etc., so steel sheets for automobile parts have high workability (such as punchability and stretch flangeability) in addition to high strength. It is also required to have.

  However, in general, as steel materials increase in strength, ductility decreases and workability deteriorates. For this reason, a steel plate with a tensile strength of 900 MPa or higher has various problems when it is formed into a desired part shape. For example, when a punching process is performed on a steel sheet having a tensile strength of 900 MPa or more, cracks, steps, turning, peeling, etc. are prominent on the end face of the punched hole, and the fatigue characteristics and dimensional accuracy of the parts are reduced. In addition, such deterioration of the properties of the end face of the punched hole adversely affects stretch flangeability.

Therefore, when applying high-strength hot-rolled steel sheets to automobile parts, etc., it is essential to develop high-strength hot-rolled steel sheets that have workability such as punchability and stretch flangeability, and various technologies have been proposed to date. ing.
For example, in Patent Document 1, the steel sheet composition is mass%, C: 0.010 to 0.200%, Si: 0.01 to 1.50%, Mn: 0.25 to 3.00%, B: 0.0002 to 0.0030%, and P: 0.05% It is limited to the following, and further contains any one or more of Ti: 0.03 to 0.20%, Nb: 0.01 to 0.20%, V: 0.01 to 0.20%, Mo: 0.01 to 0.20%, and the balance Has been proposed that has a composition comprising Fe and unavoidable impurities, and the sum of the amount of C segregated and the amount of B segregated to the large-angle grain boundaries of ferrite ranges from 4 to 10 atms / nm ² . And according to this technology, the precipitation strengthening elements Ti, V, Nb, Mo are added to precipitate carbides, etc., so that the steel sheet has a tensile strength of 690 MPa or more, and B and C grain boundary segregation. By controlling the amount, damage to the end face can be surely prevented even when punching is performed under extremely severe conditions.

  Moreover, in patent document 2, a steel plate composition is the mass%, C: 0.015-0.06%, Si: Less than 0.5%, Mn: 0.1-2.5%, P <= 0.10%, S <= 0.01%, Al: 0.005-0.3% , N ≦ 0.01%, Ti: 0.01 to 0.30%, B: 2 to 50 ppm, the balance consisting of Fe and unavoidable impurities, and further specifying the atomic ratio of carbide-forming elements and C, The content of Si, Mn, B, and Mo, which are elements that control the γ / α transformation temperature, is specified so as to satisfy a predetermined relationship, and the total area ratio of one or both of ferrite and bainitic ferrite is 90 %, And a cementite structure with a cementite area ratio of 5% or less has been proposed. And according to this technique, all of stretch flange formability, punching cracking resistance and surface condition are good, and a high strength hot rolled steel sheet having a tensile strength of 690 MPa or more can be stably manufactured at low cost. It is said that.

JP 2008-266726 A JP 2007-302992 A

  However, in the technique proposed in Patent Document 1, since it is necessary to lower the coiling temperature for the purpose of controlling the grain boundary segregation amount of B and C, the precipitation amount of carbides precipitated in ferrite is not sufficient. A high-strength steel sheet with a tensile strength of 900 MPa or more cannot be obtained. Further, in the technique proposed in Patent Document 2, as shown in the examples, it is necessary to add 0.5% or more of Mn, which is a solid solution strengthening element, in order to obtain desired steel plate characteristics. The center segregation is likely to occur, and the state of the punched end face at the center segregation portion is extremely inferior. Furthermore, as specified in Patent Document 2, with the technique proposed in Patent Document 2, when the tensile strength of the steel sheet exceeds 850 MPa, it is extremely difficult to suppress damage to the punched end face.

As described above, in all of the conventional techniques, since the steel sheet having a tensile strength of 900 MPa or more uses Mn solid solution strengthening, deterioration of the properties of the punched end face due to center segregation cannot be suppressed.
The present invention has been made in view of such circumstances, and an object thereof is to provide a high-strength hot-rolled steel sheet having a tensile strength of 900 MPa or more and excellent workability, particularly punchability.

  In order to solve the above-mentioned problems, the present inventors focused on hot-rolled steel sheets having a ferrite single-phase structure with good workability, and various factors affecting the strengthening and workability of the hot-rolled steel sheets, particularly the punchability. We studied earnestly. As a result, Mn, which has been considered to be extremely effective as a solid solution strengthening element for increasing the strength of steel sheets, has been actively incorporated into high-strength hot-rolled steel sheets. It was found that it adversely affects sex. Further, it has been found that the central segregation of Mn contained in the steel sheet is a cause of deterioration such as punchability, and that the Mn content needs to be less than 0.5% in order to suppress the segregation.

  On the other hand, a reduction in steel sheet strength accompanying the suppression of the Mn content, which is a solid solution strengthening element, is inevitable. Therefore, the present inventors adopted precipitation strengthening by Nb carbide, V carbide and Nb and V composite carbide as a strengthening mechanism instead of solid solution strengthening by Mn, and these carbides are converted into a ferrite phase which is a matrix of the steel sheet. An attempt was made to obtain a desired steel plate strength (tensile strength: 900 MPa or more) by fine precipitation. Moreover, since the effect of significantly improving the steel sheet strength can be expected as the carbides precipitated in the ferrite phase are finer and the precipitation amount is larger, a means for achieving a finer carbide and a means for ensuring a sufficient precipitation amount were sought.

  As a result, the carbide is precipitated in the cooling process after the hot rolling is finished at the time of manufacturing the steel sheet, but the carbide precipitated in a temperature range higher than the steel sheet winding temperature is likely to be coarsened, while the carbide is in the steel sheet winding temperature range. It became clear that fine carbides can be obtained by precipitation. Nb carbides, V carbides, and Nb and V composite carbides precipitate almost simultaneously with the austenite → ferrite transformation of steel during the cooling process after hot rolling. It has been found that fine carbides can be obtained by adjusting the temperature range.

  Here, Mn is an element that has the effect of lowering the austenite → ferrite transformation point of steel. However, as described above, when a large amount of Mn (0.5% or more) is contained in the steel sheet, the punchability and stretch flangeability deteriorate. End up. Therefore, the present inventors examined a means for reducing the austenite → ferrite transformation point of steel while suppressing the Mn content to less than 0.5%. As a result, it was found that the steel austenite → ferrite transformation point can be lowered without impairing various properties of the steel sheet by adding a predetermined amount of B to the steel sheet and further specifying the manufacturing conditions of the steel sheet. did. In addition to the above, it has been found that when the average grain size of the ferrite phase is 10 μm or less, the punchability of the steel sheet is further improved.

  Furthermore, in order to precipitate the carbides precipitated in the ferrite phase crystal grains, which are the matrix of the steel sheet, sufficiently to obtain a desired steel sheet strength (tensile strength: 900 MPa or more), carbide forming elements (Nb and It has been found that the total content of V, or even Ti, W, and Mo) needs to be a predetermined amount or more. In addition, in order to sufficiently precipitate fine carbides, it has been found that it is extremely effective to add not only Nb alone but also Nb and V as carbide forming elements to be added to the steel sheet.

  As described above, as a result of intensive studies by the present inventors, it has become clear that it is important to suppress the occurrence of center segregation of Mn in order to stably suppress defects on the punched end face. However, conventionally, reducing the Mn content of the steel sheet coarsens the carbide particle diameter, or the solid solution strengthening amount decreases and the steel sheet softens, so the Mn content can only be reduced to about 0.5%. No steel sheet having a high strength of 900 MPa or more and good punchability could not be obtained.

  For such problems of the prior art, the present inventors have made the steel sheet matrix contain a predetermined amount of B even when the Mn content of the steel sheet is less than 0.5%. It has been found that very fine carbides with an average particle diameter of less than 10 nm can be precipitated in the crystal grains of a ferrite phase. Further, in addition to the above B, further, the total content of carbide forming elements (Nb and V, or even Ti, W, Mo) is set to a predetermined amount or more, so that the fine carbide having an average particle diameter of less than 10 nm can be obtained. It was found that the amount of precipitation could be secured sufficiently. And, as described above, the steel plate strength is ensured without actively utilizing the solid solution strengthening of Mn, and the generation of center segregation of Mn can be prevented by reducing the amount of Mn. Furthermore, it was found that a high-strength hot-rolled steel sheet with excellent punchability can be obtained by adding B in order to maximize the production of fine carbides.

The present invention has been completed based on the above findings, and the gist thereof is as follows.
[1] By mass%
C: 0.05% or more and 0.11% or less, Si: 0.3% or less,
Mn: less than 0.5%, P: 0.03% or less,
S: 0.005% or less, Al: 0.1% or less,
N: 0.01% or less, B: 0.0005% or more and 0.005% or less,
Nb: 0.01% or more and 0.25% or less, V: 0.05% or more and 0.4% or less, so that Nb and V satisfy the following formula (1), and the balance is composed of Fe and inevitable impurities, and the ferrite phase Having an area ratio of 95% or more, an average crystal grain size of the ferrite phase of 10 μm or less, a structure in which the average grain size of carbide in the crystal grains of the ferrite phase is less than 10 nm, and a tensile strength of 900 MPa A high-strength hot-rolled steel sheet excellent in punchability, characterized by the above.
Record
[Nb] / 93 + [V] / 51 ≧ 0.0043 (1)
([Nb], [V]: Content of each element (mass%))

[2] In the above [1], in addition to the above composition, any one of Ti: 0.01% to 0.13%, W: 0.01% to 1.0%, Mo: 0.01% to 0.5% in addition to the composition A high-strength hot-rolled steel sheet excellent in punchability, characterized by containing the above so as to satisfy the following formula (2) instead of the formula (1).
Record
[Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96 ≥ 0.0043 (2)
([Nb], [V], [Ti], [W], [Mo]: Content of each element (mass%))

[3] A high-strength hot-rolled steel sheet excellent in punchability, wherein the composition in [1] or [2] satisfies the following formula (3):
Record
0.8 ≦ ([C] / 12) / ([Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96) ≦ 1.5 (3)
([C], [Nb], [V], [Ti], [W], [Mo]: Content of each element (mass%))

[4] In any one of the above [1] to [3], in addition to the composition, in addition to mass, REM, Zr, As, Cu, Ni, Sn, Pb, Ta, Cr, Sb, Mg, Ca A high-strength hot-rolled steel sheet excellent in punchability, comprising one or more of Co, Se, Zn, Cs, Ga, Ba, and Sr in a total amount of 1.0% or less.

[5] A high-strength hot-rolled steel sheet having excellent punchability, which has a plating layer on the steel sheet surface in any one of [1] to [4].

[6] A high-strength hot-rolled steel sheet having excellent punchability, wherein the plating layer is a galvanization layer in [5].

[7] The high-strength hot-rolled steel sheet having excellent punchability according to [5], wherein the plated layer is an alloyed galvanized layer.

[8] Heating the steel material, subjecting it to hot rolling consisting of rough rolling and finish rolling, cooling after completion of finish rolling, winding, and hot rolling steel sheet,
The steel material in mass%,
C: 0.05% or more and 0.11% or less, Si: 0.3% or less,
Mn: less than 0.5%, P: 0.03% or less,
S: 0.005% or less, Al: 0.1% or less,
N: 0.01% or less, B: 0.0005% or more and 0.005% or less,
Nb: 0.01% or more and 0.25% or less, V: 0.05% or more and 0.4% or less, so that Nb and V satisfy the following formula (1), and the balance is composed of Fe and inevitable impurities,
The heating temperature of the heating is 1100 ° C. or more and 1350 ° C. or less, the finish rolling temperature of the finish rolling is 850 ° C. or more, the cooling is started within 3 seconds after finishing rolling, and the cooling rate of the cooling is 20 ° C. / s or more, and by setting the winding temperature of the winding to be 550 ° C. or more and 700 ° C. or less, the area ratio of the ferrite phase is 95% or more, the average grain size of the ferrite phase is 10 μm or less, and the ferrite phase A high-strength hot-rolled steel sheet having excellent punchability, characterized in that a hot-rolled steel sheet having a structure in which the average particle size of carbides in the crystal grains is less than 10 nm and having a tensile strength of 900 MPa or more is obtained . Production method.
Record
[Nb] / 93 + [V] / 51 ≧ 0.0043 (1)
([Nb], [V]: Content of each element (mass%))

[9] In the above [8], in addition to the above composition, any one of Ti: 0.01% to 0.13%, W: 0.01% to 1.0%, Mo: 0.01% to 0.5% in addition to the composition The method for producing a high-strength hot-rolled steel sheet excellent in punchability, characterized by containing the above so as to satisfy the following formula (2) instead of the formula (1).
Record
[Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96 ≥ 0.0043 (2)
([Nb], [V], [Ti], [W], [Mo]: Content of each element (mass%))

[10] A method for producing a high-strength hot-rolled steel sheet excellent in punchability, wherein the composition in [8] or [9] satisfies the following formula (3):
Record
0.8 ≦ ([C] / 12) / ([Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96) ≦ 1.5 (3)
([C], [Nb], [V], [Ti], [W], [Mo]: Content of each element (mass%))

[11] In any one of the above [8] to [10], in addition to the composition, the REM, Zr, As, Cu, Ni, Sn, Pb, Ta, Cr, Sb, Mg, and Ca are further contained in mass%. A method for producing a high-strength hot-rolled steel sheet having excellent punchability, comprising one or more of Co, Se, Zn, Cs, Ga, Ba, and Sr in a total amount of 1.0% or less.

[12] In any one of [8] to [11], a method for producing a high strength hot rolled steel sheet having excellent punchability, wherein a plating layer is formed on a surface of the hot rolled steel sheet.

[13] A method for producing a high-strength hot-rolled steel sheet having excellent punchability, wherein the plating layer is a galvanized layer in [12].

[14] A method for producing a high-strength hot-rolled steel sheet having excellent punchability, wherein the plating layer is an alloyed galvanized layer in [12].

  According to the present invention, it is possible to obtain a high-strength steel plate having a tensile strength of 900 MPa or more and excellent in punchability, which is suitable for the use of structural members of automobiles, and enables weight reduction of automobile members and automobile member molding. The effect is remarkable. In addition, according to the present invention, a high strength hot rolled steel sheet with a workability (punchability, stretch flangeability) and a tensile strength of 900 MPa or more can be obtained. Thus, there is a remarkable industrial effect.

It is a figure which shows the relationship between the punching end surface property (punching property) of a hot-rolled steel plate, and a hole expansion rate (stretch flangeability).

Hereinafter, the present invention will be described in detail.
First, the structure of the steel sheet of the present invention and the reasons for limiting the carbide will be described.
The hot-rolled steel sheet of the present invention has a structure in which the ferrite phase area ratio is 95% or more, the average crystal grain size of the ferrite phase is 10 μm or less, and the average carbide grain size in the ferrite phase crystal grains is less than 10 nm. Have

Area ratio of ferrite phase: 95% or more It is preferable that the matrix of the hot-rolled steel sheet has a ferrite single-phase structure excellent in workability. As described above, in the present invention, a desired strength of the steel sheet is ensured by precipitating fine carbides in the ferrite crystal grains which are the matrix of the hot-rolled steel sheet. Therefore, in the present invention, it is necessary to precipitate C added to the steel material as fine carbides. If coarse carbides such as cementite are present, the amount of C forming fine carbides is reduced, and the steel sheet strength is reduced. . Furthermore, since microvoids are likely to be generated at the interface between coarse carbides such as cementite and ferrite, the properties of the punched end face are deteriorated.

  Moreover, when bainite and martensite are formed, the hardness difference between the structures with the ferrite phase becomes large, and microvoids are easily generated between the structures, so that the punchability is reduced as described above. For the above reasons, when the area ratio of the ferrite phase is less than 95%, the adverse effect due to the generation of microvoids becomes obvious and the punching end face properties deteriorate, or the amount of fine carbides precipitated in the ferrite crystal grains is insufficient. Steel plate strength (tensile strength: 900 MPa) cannot be obtained. Therefore, it is necessary that the hot rolled steel sheet structure of the present invention has a ferrite single phase, that is, a ferrite phase area ratio of 95% or more. Preferably it is 98% or more.

  In the hot rolled steel sheet of the present invention, examples of the structure other than the ferrite phase that can be contained in the matrix include cementite, pearlite, bainite phase, martensite phase, and the like. If these structures are present in a large amount in the matrix, the steel sheet properties (punchability, stretch flangeability, etc.) deteriorate. Therefore, it is preferable to reduce these structures as much as possible, but it is acceptable if the total area ratio with respect to the entire matrix structure is 5% or less. Preferably it is 2% or less.

Average crystal grain size of ferrite phase: 10 μm or less If the average crystal grain size of the ferrite phase exceeds 10 μm, the deformation at the time of punching becomes non-uniform, and cracks at the punched end face tend to occur, resulting in a decrease in punchability. Therefore, the upper limit of the average crystal grain size of the ferrite phase is set to 10 μm. Preferably, it is 8 μm or less.

Carbide in ferrite grains Therefore, it is not expected to improve the strength of the steel sheet by solid solution strengthening. Therefore, in the hot-rolled steel sheet of the present invention, it is essential to finely precipitate carbide in the ferrite phase crystal grains in order to ensure strength. Examples of the carbide finely precipitated in the ferrite phase grains in the present invention include Nb carbide, V carbide, and Nb and V composite carbide, or those containing Ti, W, and Mo in the carbide. Most of these carbides are carbides that precipitate at the interface at the same time as the austenite → ferrite transformation in the cooling process after finishing rolling in the hot rolled steel sheet manufacturing process.

Carbide average particle diameter in ferrite crystal grains: less than 10 nm In the steel of the present invention, strengthening is achieved by finely dispersing the carbides such as Nb and V described above. When the carbides are coarsened, the number of carbides that hinder the movement of dislocations that are generated when deformation is applied to the steel sheet decreases, so that the strength of the steel sheet increases as the carbides become finer. In order to obtain a high-strength hot-rolled steel sheet having a tensile strength of 900 MPa or more, it is necessary that the average particle diameter of the carbide is less than 10 nm. Preferably it is 6 nm or less.

Next, the reason for limiting the component composition of the hot rolled steel sheet according to the present invention will be described. In addition,% showing the following component composition shall mean the mass% (mass%) unless there is particular notice.
C: 0.05% or more and 0.11% or less
C combines with Nb, V, or Ti, Mo, W, and is finely dispersed in the steel sheet as a carbide. That is, C is an element that forms fine carbides and remarkably strengthens the ferrite structure, and is an essential element for strengthening the hot-rolled steel sheet. In order to obtain a high-strength hot-rolled steel sheet having a tensile strength of 900 MPa or more, the C content needs to be at least 0.05%. On the other hand, when the C content exceeds 0.11%, a large amount of cementite precipitates. And since a micro void is easy to produce | generate at a cementite and a matrix (ferrite) interface, a punching end surface property will deteriorate remarkably. Therefore, the C content is 0.05% or more and 0.11%. Preferably they are 0.06% or more and 0.10% or less.

Si: 0.3% or less
Si is positively contained in conventional high-strength steel sheets as an effective element for improving the steel sheet strength without reducing ductility (elongation). However, Si tends to be concentrated on the surface of the steel sheet, and the steel sheet surface is partially cured by this concentration, so that burrs are easily generated when the steel sheet is punched, and the punchability is deteriorated. Therefore, in the present invention, it is desirable to reduce the Si content as much as possible, but up to 0.3% is acceptable, so the upper limit of the Si content is set to 0.3%. Preferably it is 0.1% or less. Note that the Si content may be reduced to the impurity level.

Mn: Less than 0.5%
The Mn content is one of the important requirements in the present invention. Mn is a solid solution strengthening element and, like Si, is actively contained in conventional high-strength steel sheets. However, Mn generates inevitable center segregation during casting. And since this center segregation part is very hard and is inferior to ductility, a crack generate | occur | produces at the time of a punching process, and end surface property deteriorates. For this reason, in the present invention, it is necessary to reduce the Mn content as much as possible to less than 0.5% in order to ensure good punched end face properties. Preferably it is 0.4% or less. Note that the Mn content may be reduced to the impurity level.

P: 0.03% or less
P is a harmful element that segregates at the grain boundary and becomes the starting point of grain boundary cracking during processing, and deteriorates the properties of the punched end face, so it is preferable to reduce it as much as possible. Therefore, in the present invention, the P content is set to 0.03% or less in order to avoid the above problems. Preferably it is 0.02% or less.

S: 0.005% or less
S exists as an inclusion such as MnS in steel. This inclusion causes a remarkable adverse effect on the punchability because the deformation is not uniform in the wedge shape when the steel sheet is punched. Therefore, in the present invention, it is preferable to reduce the S content as much as possible, and it is 0.005% or less. Preferably it is 0.003% or less.

Al: 0.1% or less
Al is an element that acts as a deoxidizer. In order to acquire such an effect, it is desirable to contain 0.02% or more, but if the Al content exceeds 0.1%, the adverse effect on punchability due to inclusions such as alumina becomes obvious. Therefore, the Al content is 0.1% or less. Preferably it is 0.08% or less.

N: 0.01% or less
N combines with carbide forming elements Nb and V at the steel making stage to form coarse nitrides and inhibits the formation of fine carbides, so that the steel sheet strength is significantly reduced. Coarse nitrides also reduce the punchability of the steel sheet. Therefore, the N content is preferably reduced as much as possible, and is 0.01% or less. Preferably it is 0.008% or less.

B: 0.0005% or more and 0.005% or less
B is an important element in the present invention. As described above, in the present invention, the austenite → ferrite transformation point of steel in the cooling process after the end of hot rolling is adjusted to the temperature range of the coiling temperature described later for the purpose of refining the carbides precipitated in the ferrite crystal grains. Since it is necessary, B is added. In order to lower the austenite → ferrite transformation point to 700 ° C., which is the upper limit of the coiling temperature, for the steel with a reduced Mn content, the B content needs to be 0.0005% or more. On the other hand, if the B content exceeds 0.005%, the effect of controlling the transformation point is saturated. Therefore, the B content is 0.0005% or more and 0.005% or less. Preferably it is 0.0008% or more and 0.003% or less.

Nb: 0.01% or more and 0.25% or less
Nb inhibits the recrystallization of austenite during hot rolling, so that the ferrite grains after austenite → ferrite transformation are refined in the cooling and winding process following hot rolling, and within the ferrite grains after transformation. By forming fine carbides, it contributes to increasing the strength of the steel sheet. In order to obtain such an effect, the Nb content needs to be 0.01% or more. On the other hand, when the Nb content exceeds 0.25%, when manufacturing a hot-rolled steel sheet, coarse Nb carbides are not completely dissolved during heating of the steel material before hot rolling, and finally obtained (after winding) (Ii) Coarse Nb carbide remains on the hot-rolled steel sheet. And this coarse Nb carbide causes micro void generation at the time of punching and causes the punchability of the steel sheet to deteriorate. Therefore, the Nb content is 0.01% or more and 0.25% or less. Preferably they are 0.05% or more and 0.2% or less.

V: 0.05% or more and 0.4% or less
V, like Nb, is an element that contributes to increasing the strength of the steel sheet by forming carbides with C. In addition, V is effective in increasing the strength of the steel sheet because it combines with Nb to form fine composite carbide. In order to ensure the desired steel plate strength (tensile strength: 900 MPa or more), the V content needs to be 0.05% or more. On the other hand, if the V content exceeds 0.4%, when manufacturing a hot-rolled steel sheet, the carbide tends to be coarsened during winding after the hot rolling, and the hot-rolled steel sheet is softened. Therefore, the V content is 0.05% or more and 0.4% or less. This is preferably 0.05% or more and 0.35% or less.
Thus, in the present invention, it is important to effectively add N having a higher solubility in addition to using Nb having a lower solubility in steel. In the present invention, a hot-rolled steel sheet having a tensile strength of 900 MPa or more can be obtained by adding Nb and V in combination, and at least by adding Nb alone, a sufficient precipitation amount of fine carbides can be secured to obtain a desired steel sheet strength. Can not do it.

The hot-rolled steel sheet of the present invention contains Nb and V so as to satisfy the above-described range and the expression (1).
[Nb] / 93 + [V] / 51 ≧ 0.0043 (1)
([Nb], [V]: Content of each element (mass%))
The above formula (1) is a requirement that must be satisfied in order to make the tensile strength of the hot-rolled steel sheet 900 MPa or more, and is an important index in the present invention.

  As described above, in the present invention, the desired strength of the steel sheet is ensured by precipitating fine carbides in the ferrite phase crystal grains which are the matrix of the hot-rolled steel sheet. Here, since the steel sheet strength is improved as the amount of fine carbide precipitation increases, in the present invention, the total content of carbide forming elements Nb and V is a predetermined amount or more, and the steel sheet has a tensile strength of 900 MPa or more. Therefore, it is necessary to secure a sufficient amount of fine carbide precipitation for imparting.

  Therefore, the present inventors examined atomic% of carbide-forming elements contained in the steel sheet, and if the total ([Nb] / 93 + [V] / 51) is 0.0043 or more, the tensile strength is 900 MPa or more. It has been found that a hot-rolled steel sheet can be obtained. Therefore, in the present invention, the value of [Nb] / 93 + [V] / 51 is set to 0.0043 or more. Preferably it is 0.0045 or more. However, if the value of [Nb] / 93 + [V] / 51 exceeds 0.011, coarse carbides cannot be dissolved by heating the steel material before hot rolling as described above, and Nb and V are added. Since there is a concern that the effect of increasing strength is saturated, the above value is preferably set to 0.011 or less.

The above is the basic composition in the present invention. In addition to the above basic composition, Ti: 0.01% or more and 0.13% or less, W: 0.01% or more and 1.0% or less, Mo: 0.01% or more and 0.5% or less More than one seed may be contained so as to satisfy the following formula (2) instead of the above formula (1).
[Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96 ≥ 0.0043 (2)
([Nb], [V], [Ti], [W], [Mo]: Content of each element (mass%))

  Ti, W, and Mo are elements that contribute to increasing the strength of the steel sheet by forming a fine composite carbide by combining with Nb and V together with C. Furthermore, carbides containing Ti, W, and Mo are difficult to coarsen, and an effect of suppressing the softening of the steel sheet can be expected in a coiling temperature range of 700 ° C. or lower as will be described later. In order to obtain such an effect, it is preferable that the Ti content is 0.01% or more, the W content is 0.01% or more, and the Mo content is 0.01% or more. On the other hand, if the W content exceeds 1.0% and the Mo content exceeds 0.5%, the austenite → ferrite transformation is not completed during coil winding during the hot-rolled steel sheet manufacturing and cooling process after hot rolling, and the ferrite area It becomes difficult to make the rate 95% or more, and the punchability of the hot-rolled steel sheet may be lowered. On the other hand, if the Ti content exceeds 0.13%, coarse Ti carbides and carbides containing Ti and Nb cannot be dissolved during heating of the steel material before hot rolling, and the hot-rolled steel sheet strength does not increase. Therefore, it is preferable that the Ti content is 0.01% or more and 0.13% or less, the W content is 0.01% or more and 1.0% or less, and the Mo content is 0.01% or more and 0.5% or less. More preferably, the Ti content is 0.01% to 0.1%, the W content is 0.01% to 0.3%, and the Mo content is 0.01% to 0.35%.

  Further, as described above, in order to secure a sufficient amount of fine carbide precipitation for imparting a tensile strength of 900 Pma or more to the hot-rolled steel sheet, the total atomic% of carbide forming elements needs to be 0.0043 or more. Therefore, when one or more of Ti, W and Mo is contained as a carbide forming element in addition to Nb and V, the total of these atomic%, that is, [Nb] / 93 + [V] / 51 + [Ti ] / 48 + [W] / 184 + [Mo] / 96 is set to 0.0043 or more. Preferably it is 0.0045 or more. However, if the value of [Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96 exceeds 0.011, coarse carbides will be dissolved during heating before hot rolling. The above value is preferably 0.011 or less because there is a concern that the ferrite transformation may not be completed in the winding process.

In addition, the hot-rolled steel sheet of the present invention preferably contains C, Nb and V, or further Ti, W, and Mo within the above-described range and satisfying the expression (3).
0.8 ≦ ([C] / 12) / ([Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96) ≦ 1.5 (3)
([C], [Nb], [V], [Ti], [W], [Mo]: Content of each element (mass%))

  C atomic% ([C] / 12) with respect to total atomic% ([Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96) of carbide forming elements contained in hot-rolled steel sheet ) Ratio of less than 0.8, the carbide-forming element may not be sufficiently precipitated as carbide, and there is a concern that a tensile strength of 900 MPa or more cannot be obtained. On the other hand, the C atom% ([C] to the total atom% ([Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96) of carbide forming elements contained in the hot-rolled steel sheet) If the ratio of / 12) exceeds 1.5, C that does not bind to the carbide-forming element may produce coarse cementite. And this coarse cementite tends to be a starting point of microvoids when the steel plate is deformed, and therefore, it may cause cracks when punching, and good punchability may not be obtained.

  Therefore, the value of ([C] / 12) / ([Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96) is preferably 0.8 or more and 1.5 or less. More preferably, it is 0.9 or more and 1.4 or less. In calculating the value of ([C] / 12) / ([Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96) in the above equation (3), When the hot-rolled steel sheet does not contain Ti, W, and Mo, the values of [Ti], [W], and [Mo] are calculated as zero.

  In addition to the above basic composition, any one of REM, Zr, As, Cu, Ni, Sn, Pb, Ta, Cr, Sb, Mg, Ca, Co, Se, Zn, Cs, Ga, Ba, Sr One or more kinds may be contained in total of 1% or less. From the viewpoint of punchability, the total content of these elements is preferably 0.5% or less. Components other than the above are Fe and inevitable impurities.

Tensile strength: 900MPa or more
In a steel plate having a tensile strength of 900 MPa or more, the amount of sagging during punching is smaller than that of a steel plate having a tensile strength of less than 900 MPa. Therefore, in the present invention, the lower limit of the tensile strength of the hot-rolled steel sheet is set to 900 MPa.

  The hot-rolled steel sheet of the present invention has little material fluctuation even when subjected to heat treatment up to 700 ° C., which is the upper limit of the coiling temperature described later. Therefore, for the purpose of imparting corrosion resistance to the steel sheet, the hot-rolled steel sheet of the present invention can be plated, and a plating layer can be provided on the surface thereof. Since the plated steel sheet can be produced even when the heating temperature in the plating process is 700 ° C. or less, the effect of the present invention described above is not impaired even if the hot-rolled steel sheet of the present invention is plated. The type of the plating layer is not particularly limited, and any of an electroplating layer and an electroless plating layer can be applied. Further, the alloy component of the plating layer is not particularly limited, and a hot dip galvanized layer, an alloyed hot dip galvanized layer and the like can be mentioned as suitable examples, but of course, it is not limited to these, and any conventionally known one can be applied. is there.

Next, the manufacturing method of the hot rolled steel sheet of the present invention will be described.
In the present invention, the steel material (steel slab) having the above composition is heated, subjected to hot rolling consisting of rough rolling and finish rolling, cooled after completion of finish rolling, and wound into a hot-rolled steel sheet. At this time, the heating temperature of the heating is 1100 ° C. or more and 1350 ° C. or less, the finishing rolling temperature of the finish rolling is 850 ° C. or more, the cooling is started within 3 seconds after finishing rolling, and the average cooling rate of the cooling Is 20 ° C./s or more, and the winding temperature of the winding is 550 ° C. or more and 700 ° C. or less.

  In the present invention, the method for melting steel is not particularly limited, and a known melting method such as a converter or an electric furnace can be employed. Further, secondary refining may be performed in a vacuum degassing furnace. Thereafter, the slab (steel material) is preferably formed by a continuous casting method from the viewpoint of productivity and quality, but the slab may be formed by a known casting method such as an ingot-bundling rolling method or a thin slab continuous casting method. . In the present invention, since Mn is reduced, the drawability at 700 ° C. or higher is good, and therefore production by continuous casting becomes easy.

Heating temperature of steel material: 1100 ° C or higher and 1350 ° C or lower The steel material obtained as described above is subjected to rough rolling and finish rolling. In the present invention, the steel material is heated prior to rough rolling to be substantially homogeneous. The austenite phase needs to be dissolved and coarse carbides need to be dissolved. When the heating temperature of the steel material is below 1100 ° C, coarse carbides do not dissolve, so the amount of carbides that are finely dispersed in the cooling and winding process after hot rolling is reduced, and the hot rolling finally obtained is reduced. The strength of the steel sheet is significantly reduced. On the other hand, when the heating temperature exceeds 1350 ° C., the scale bites and deteriorates the surface properties of the steel sheet.

  For the above reasons, the heating temperature of the steel material is set to 1100 ° C or higher and 1350 ° C or lower. Preferably they are 1150 degreeC or more and 1300 degrees C or less. However, when hot rolling the steel material, if the steel material after casting is in the temperature range of 1100 ° C or higher and 1350 ° C or lower, or if the carbide of the steel material is dissolved, the steel material is heated. Direct rolling may be performed without any problem. The rough rolling conditions are not particularly limited.

Finishing rolling temperature: 850 ° C or higher When the finishing rolling temperature is lower than 850 ° C, the ferrite transformation starts during finish rolling, and the ferrite grains are expanded, and the mixed grain structure in which the ferrite grains partially grow Therefore, the punchability of the hot-rolled steel sheet is significantly reduced. Therefore, the finish rolling temperature is 850 ° C. or higher. Preferably it is 870 degreeC or more. The upper limit of the finish rolling temperature is not particularly defined, but the finish rolling temperature is naturally determined by the heating temperature before hot rolling, the sheet feeding speed, and the steel plate thickness. From this viewpoint, the finish rolling temperature is substantially 980 ° C. or lower.

Time until start of forced cooling after finish rolling: within 3 s seconds In a high-temperature steel sheet immediately after finish rolling, carbides are generated due to strain-induced precipitation because the strain energy accumulated in the austenite phase is large. Since this carbide precipitates at a high temperature and is easy to coarsen, when strain-induced precipitation occurs, it becomes difficult to obtain a fine precipitate. Therefore, in the present invention, for the purpose of suppressing strain-induced precipitation, forced cooling needs to be started immediately after the end of hot rolling, and cooling is started within at least 3 seconds after the end of finish rolling. Preferably it is within 2 s.

Average cooling rate: 20 ° C./s or more As described above, the longer the time during which the steel sheet is maintained at a high temperature after finishing rolling, the more easily the coarsening of the carbide by strain-induced precipitation proceeds. In the present invention, the steel sheet contains a predetermined amount of B to suppress the austenite → ferrite transformation. However, if the cooling rate is low, the ferrite transformation starts at a high temperature, and the carbide tends to become coarse. Therefore, it is necessary to rapidly cool after finish rolling, and to avoid the above problem, it is necessary to cool at an average cooling rate of 20 ° C./s or more. Preferably it is 40 ° C./s or more. However, if the cooling rate after finishing rolling is excessively increased, it may be difficult to control the coiling temperature.

Winding temperature: 550 ° C. or higher and 700 ° C. or lower When the winding temperature is lower than 550 ° C., a sufficient amount of precipitates cannot be obtained, and the hot-rolled steel sheet strength decreases. On the other hand, when the coiling temperature exceeds 700 ° C., the precipitated carbide is coarsened, so that the strength of the hot rolled steel sheet is lowered. Accordingly, the coiling temperature range is 550 ° C. or more and 700 ° C. or less. Preferably they are 580 degreeC or more and 680 degrees C or less.

  As described above, according to the present invention, by optimizing the steel sheet composition and production conditions, the ferrite phase area ratio is 95% or more, the average crystal grain size of the ferrite phase is 10 μm or less, A hot-rolled steel sheet having a structure in which the carbide average particle diameter in the crystal grains is less than 10 nm is obtained. In the present invention, the Mn content of the steel sheet is reduced to less than 0.5% for the purpose of improving punchability, and the strength is increased by a particle dispersion strengthening mechanism instead of solid solution strengthening by Mn. It can be set as the high intensity | strength hot-rolled steel plate excellent in the flange property. Furthermore, in the present invention, the content of carbide forming elements (Nb and V, or even Ti, W, Mo) contained in the steel sheet is optimized, and a predetermined amount of B having an effect of suppressing the coarsening of the carbide is contained. In addition, the production conditions for hot-rolled steel sheets are specified. As a result, the carbide having the average particle diameter of less than 10 nm can be sufficiently precipitated in the ferrite crystal grains, and the tensile strength of the hot-rolled steel sheet reaches 900 MPa or more while maintaining the workability such as punchability. Can be increased.

  For the strengthening by the particle dispersion strengthening mechanism, a tensile strength of 900 MPa cannot be obtained unless the component composition, manufacturing method, and the like are appropriately selected. The amount of strengthening by the particle dispersion strengthening mechanism is related to the particle size and volume fraction of carbides. To maximize the amount of particle dispersion strengthening, the particle size of the precipitate is made as small as possible and the volume fraction is increased. This is because it becomes important. Here, the present invention suppresses the content of Mn to less than 0.5% for the purpose of improving punchability and the like, but since Mn is a solid solution strengthening element and has an effect of suppressing carbide coarsening, Mn Reducing the content is disadvantageous in securing the steel sheet strength. However, in the present invention, a predetermined amount of B, which is effective in suppressing the coarsening of precipitates, is contained, Nb having a low solubility in steel and V having a high solubility are effectively combined as carbide forming elements, and further hot-rolled steel plate By optimizing the production conditions, the particle dispersion strengthening mechanism is utilized to the maximum extent. Therefore, a high-strength hot-rolled steel sheet having excellent punchability and a tensile strength of 900 MPa or more can be obtained.

  In addition, the hot rolled steel sheet of the present invention after hot rolling has its characteristics changed even if the scale is attached to the surface or the scale is removed by pickling. No, the above-described excellent characteristics are exhibited in any state. In the present invention, the hot-rolled steel sheet after winding may be plated to form a plating layer on the surface of the hot-rolled steel sheet. The type of the plating treatment is not particularly limited, and both electroplating treatment and electroless plating treatment are applicable. For example, a hot dip galvanizing process can be performed as a plating process to form a hot dip galvanized layer. Alternatively, after the hot dip galvanizing treatment, an alloying treatment may be further performed to form an alloyed hot dip galvanized layer. In addition to zinc, other metals and alloys such as aluminum or aluminum alloy can be plated for hot dipping.

  The hot-rolled steel sheet obtained by the present invention does not change the state of precipitates in the temperature range up to 700 ° C. or less. Therefore, for example, a continuous plating line with an annealing temperature of 700 ° C. or less can be passed. Examples of the method for attaching the plating layer include a method in which a steel sheet is immersed in a plating bath and pulled up. Examples of the alloying method include a method performed in a furnace capable of heating the steel sheet surface such as a gas furnace after plating.

A steel material having a thickness of 250 mm having the composition shown in Table 1 was formed into a hot rolled steel sheet having a thickness of 1.2 to 3.2 mm under the hot rolling conditions shown in Table 2. In addition, the cooling rate described in Table 2 is an average cooling rate from the finish rolling temperature to the winding temperature. Moreover, about a part of obtained hot-rolled steel sheet, it passes through the hot dip galvanizing line of annealing temperature 700 degreeC, and is then immersed in a 460 degreeC plating bath (plating composition: Zn-0.13mass% Al), A hot-dip galvanized material (GI material) was used. Further, a part of the hot dip galvanized material (GI material) was subjected to alloying treatment at 520 ° C. after immersion in the plating bath to obtain an alloyed hot dip galvanized material (GA material). The plating adhesion amount was 45 g / m ² per side for both GI and GA materials.
In addition, except for steel plates Nos. 3 to 7 and 16 to 18 in Table 2, it was separately confirmed that transformation from austenite to ferrite did not occur during cooling until winding.

  Samples are taken from the hot-rolled steel sheet (hot-rolled steel sheet, GI material, GA material) obtained as described above, and the structure observation and tensile test are performed. The area ratio of the ferrite phase, the type and area ratio of the structure other than the ferrite phase The average crystal grain size of the ferrite phase, the average particle size of the carbide, the yield strength, the tensile strength, the elongation, and the hole expansion rate were determined. In addition, a test piece is taken from the hot-rolled steel sheet obtained as described above, a punching test and a hole expansion test are performed, and a property evaluation of the punched end surface, an evaluation of stretch flange property, and confirmation of the presence or absence of segregation at the center of the plate thickness went. The test method was as follows.

(I) Microstructure observation The area ratio of the ferrite phase was evaluated by the following method. About the central part of the plate thickness with a cross section parallel to the rolling direction, the corrosion appearance structure by 5% nital was magnified 400 times with a scanning optical microscope and photographed for 10 fields of view. The ferrite phase is a structure having a form in which corrosion marks and cementite are not observed in the grains. Further, the area ratio and particle size were determined using polygonal ferrite, bainitic ferrite, acicular ferrite and granular ferrite as ferrite.
The area ratio of the ferrite phase was determined by separating the ferrite phase from those other than the ferrite phase such as bainite and martensite by image analysis and determining the area ratio of the ferrite phase with respect to the observation field. At this time, the grain boundary observed as a linear form was counted as a part of the ferrite phase.
The average crystal grain size of the ferrite phase was obtained by enlarging the above 400 times, drawing three horizontal lines and three vertical lines for each of the three representative photographs, and determining the final grain size according to ASTM E 112-10. Table 3 shows the average value of the three sheets.
The average particle size of the carbides in the ferrite phase grains was measured by a transmission electron microscope (magnification: 120,000 times) by preparing a sample from the center of the thickness of the obtained hot-rolled steel sheet using a thin film method. It calculated | required by the average of the particle diameter of the precipitate more than a point. In calculating the particle size of the precipitate, coarse cementite and nitride having a particle size of 1.0 μm or more were not included.

(Ii) Tensile test JIS No. 13 B tensile test specimen was produced from the obtained hot-rolled steel sheet in the direction perpendicular to the rolling direction, and the tensile test was conducted five times in accordance with the provisions of JIS Z 2241 (2011). (YS), tensile strength (TS), and total elongation (El) were determined. The crosshead speed in the tensile test was 10 mm / min.

(Iii) Punching test (evaluation of punched end face properties)
Each of the obtained hot-rolled steel sheets was punched at 50 points in the longitudinal direction of the steel sheet, and the presence or absence of defects on the end face was visually observed. The evaluation was “X” when an abnormal part was observed on the end face such as a crack, a step, a turn, or a peel on the end face, and “◯” when the abnormal part was not observed.

(Iv) Hole expansion test (stretch flangeability evaluation)
A hole expansion test was performed using the test pieces (50 per each hot-rolled steel sheet) in which the punched end face properties were observed, and the stretch flangeability was evaluated. The test conditions were in accordance with the Japan Iron and Steel Federation Standard (T1001-1996), punching of a diameter of 10 mm with a clearance of 12% was performed at the center of a 100 W × 100 L mm sample, and a punch with a truncated cone was used. The hole expansion ratio (λ) represented by the following formula was calculated for each test piece, and the average value of 50 test pieces was determined for each hot-rolled steel sheet. In the following equation, the “post-test hole diameter” means that a punch with a truncated cone is inserted into the initial hole (diameter 10 mm) obtained by punching, the hole is expanded, and a crack is formed on the hot-rolled steel sheet (test piece). It is the diameter of the hole when penetrating.
(Hole expansion ratio λ%) = (Pore diameter after test−Initial hole diameter (10 mm)) / (Initial hole diameter) × 100

(V) Observation of center segregation A section parallel to the rolling direction was cut out from the test piece after the hole expansion test (50 pieces per hot-rolled steel sheet), and the section was analyzed by line analysis in the thickness direction using EPMA. It was investigated whether central segregation was observed in the center of the thickness. If the X-ray intensity of Mn at the center of the plate thickness is 1.5 times or more of the X-ray intensity of Mn at the 1/4 thickness portion, the evaluation is “×” if there is center segregation, and if it is less than 1.5 times, “ ○ ”.
The obtained results are shown in Table 3.

  Each of the inventive examples is a hot-rolled steel sheet having a tensile strength of TS: 900 MPa or more, excellent punched end face properties, and having both strength and workability. On the other hand, in the comparative example that is out of the scope of the present invention, a predetermined high strength cannot be ensured, or good punched end face properties and a sufficient hole expansion rate are not obtained.

  FIG. 1 illustrates the relationship between the hole expansion rate and the presence or absence of defects in the punched end face of the steel plate No. 17 shown in Table 3. The sample No. on the horizontal axis in FIG. 1 corresponds to 50 test pieces in the above (iii) punching test (evaluation of punched end surface properties). Moreover, the hole expansion rate λ (%) on the vertical axis in FIG. 1 corresponds to the hole expansion rate obtained in the above (iv) hole expansion test (elongation flange property evaluation), and is punched in the test of (iii) above. “●” indicates that a defect was observed on the end face, and “■” indicates that no defect was observed on the punched end face. As shown in FIG. 1, the hole expansion rate was reduced at the portion where the punched end face had a defect. Further, center segregation was observed in the portion where the punched end face had a defect.

  Table 4 shows the relationship between the hole expansion ratio and the center segregation with respect to the presence or absence of defects in the punched end face of the steel plate No. 17 shown in Table 3. “Sample No.” in Table 4 corresponds to 50 test pieces in the above (iii) punching process test (evaluation of punched end face properties). Further, “hole expansion rate λ (%)” in Table 4 corresponds to the hole expansion rate obtained in the above (iv) hole expansion test (elongation flange property evaluation). Further, “Presence / absence of center segregation” in Table 4 represents the result of the above (v) observation of center segregation. As shown in Table 4, center segregation was observed in the portion having a defect on the punched end face, and the hole expansion rate was lowered. As the above result shows, if the punchability is good, the stretch flangeability is also improved. Moreover, it can be said that it is important to completely prevent center segregation in order to obtain a good end face property stably in the longitudinal direction of the steel sheet.

Claims (14)

% By mass
C: 0.05% or more and 0.11% or less, Si: 0.3% or less,
Mn: less than 0.5%, P: 0.03% or less,
S: 0.005% or less, Al: 0.1% or less,
N: 0.01% or less, B: 0.0005% or more and 0.005% or less,
Nb: 0.01% or more and 0.25% or less, V: 0.05% or more and 0.4% or less, so that Nb and V satisfy the following formula (1), and the balance is composed of Fe and inevitable impurities, and the ferrite phase Having an area ratio of 95% or more, an average crystal grain size of the ferrite phase of 10 μm or less, a structure in which the average grain size of carbide in the crystal grains of the ferrite phase is less than 10 nm, and a tensile strength of 900 MPa A high-strength hot-rolled steel sheet excellent in punchability, characterized by the above.
Record
[Nb] / 93 + [V] / 51 ≧ 0.0043 (1)
([Nb], [V]: Content of each element (mass%)) In addition to the above-mentioned composition, any one or more of Ti: 0.01% or more and 0.13% or less, W: 0.01% or more and 1.0% or less, Mo: 0.01% or more and 0.5% or less in terms of mass% is expressed by the above formula (1). Instead, the high strength hot-rolled steel sheet having excellent punchability according to claim 1, which is contained so as to satisfy the following formula (2).
Record
[Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96 ≥ 0.0043 (2)
([Nb], [V], [Ti], [W], [Mo]: Content of each element (mass%)) The high-strength hot-rolled steel sheet having excellent punchability according to claim 1 or 2, wherein the composition satisfies the following formula (3).
Record
0.8 ≦ ([C] / 12) / ([Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96) ≦ 1.5
(3)
([C], [Nb], [V], [Ti], [W], [Mo]: Content of each element (mass%))   In addition to the above-mentioned composition, in mass%, among REM, Zr, As, Cu, Ni, Sn, Pb, Ta, Cr, Sb, Mg, Ca, Co, Se, Zn, Cs, Ga, Ba, Sr A high-strength hot-rolled steel sheet having excellent punchability according to any one of claims 1 to 3, wherein one or more of the above are contained in total of 1.0% or less.   The high-strength hot-rolled steel sheet having excellent punchability according to any one of claims 1 to 4, wherein the steel sheet surface has a plating layer.   The high-strength hot-rolled steel sheet having excellent punchability according to claim 5, wherein the plated layer is a galvanized layer.   The high-strength hot-rolled steel sheet having excellent punchability according to claim 5, wherein the plated layer is an alloyed galvanized layer. The steel material is heated, subjected to hot rolling consisting of rough rolling and finish rolling, and after finishing rolling, cooled, wound, and hot rolled steel sheet,
The steel material in mass%,
C: 0.05% or more and 0.11% or less, Si: 0.3% or less,
Mn: less than 0.5%, P: 0.03% or less,
S: 0.005% or less, Al: 0.1% or less,
N: 0.01% or less, B: 0.0005% or more and 0.005% or less,
Nb: 0.01% or more and 0.25% or less, V: 0.05% or more and 0.4% or less, so that Nb and V satisfy the following formula (1), and the balance is composed of Fe and inevitable impurities,
The heating temperature of the heating is 1100 ° C. or more and 1350 ° C. or less, the finish rolling temperature of the finish rolling is 850 ° C. or more, the cooling is started within 3 seconds after finishing rolling, and the cooling rate of the cooling is 20 ° C. / s or more, and by setting the winding temperature of the winding to be 550 ° C. or more and 700 ° C. or less, the area ratio of the ferrite phase is 95% or more, the average grain size of the ferrite phase is 10 μm or less, and the ferrite phase A high-strength hot-rolled steel sheet having excellent punchability, characterized in that a hot-rolled steel sheet having a structure in which the average particle size of carbides in the crystal grains is less than 10 nm and having a tensile strength of 900 MPa or more is obtained . Production method.
Record
[Nb] / 93 + [V] / 51 ≧ 0.0043 (1)
([Nb], [V]: Content of each element (mass%)) In addition to the above-mentioned composition, any one or more of Ti: 0.01% or more and 0.13% or less, W: 0.01% or more and 1.0% or less, Mo: 0.01% or more and 0.5% or less in terms of mass% is expressed by the above formula (1). Instead, it contains so that the following formula (2) may be satisfied, The method for producing a high-strength hot-rolled steel sheet having excellent punchability according to claim 8.
Record
[Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96 ≥ 0.0043 (2)
([Nb], [V], [Ti], [W], [Mo]: Content of each element (mass%)) The method for producing a high-strength hot-rolled steel sheet having excellent punchability according to claim 8 or 9, wherein the composition satisfies the following formula (3).
Record
0.8 ≦ ([C] / 12) / ([Nb] / 93 + [V] / 51 + [Ti] / 48 + [W] / 184 + [Mo] / 96) ≦ 1.5
(3)
([C], [Nb], [V], [Ti], [W], [Mo]: Content of each element (mass%))   In addition to the above-mentioned composition, in mass%, among REM, Zr, As, Cu, Ni, Sn, Pb, Ta, Cr, Sb, Mg, Ca, Co, Se, Zn, Cs, Ga, Ba, Sr The method for producing a high-strength hot-rolled steel sheet having excellent punchability according to any one of claims 8 to 10, wherein one or more of the above are contained in a total of 1.0% or less.   The method for producing a high-strength hot-rolled steel sheet with excellent punchability according to any one of claims 8 to 11, wherein a plating layer is formed on the surface of the hot-rolled steel sheet.   The method for producing a high-strength hot-rolled steel sheet having excellent punchability according to claim 12, wherein the plated layer is a galvanized layer.   The method for producing a high-strength hot-rolled steel sheet having excellent punchability according to claim 12, wherein the plated layer is an alloyed galvanized layer.

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