JP5971281B2 - Method for producing high-strength hot-rolled steel sheet with excellent workability and toughness - Google Patents

Description

  The present invention relates to a high-strength hot-rolled steel sheet having high tensile strength (TS): 980 MPa or more, excellent workability and toughness, suitable for materials for automobile parts and the like, and a method for producing the same.

In order to reduce CO ₂ emissions from the viewpoint of global environmental conservation, the automobile industry as a whole is required to improve automobile fuel efficiency. The most effective way to improve the fuel efficiency of automobiles is to reduce the thickness of various parts of the automobile to reduce the weight of the automobile body. Recently, from the viewpoint of ensuring the safety of passengers in the event of a collision, it is required to increase the strength of parts.

  For these reasons, high-strength steel sheets are actively used as materials for automotive parts. In particular, in recent years, higher strength has been promoted, and a high strength steel sheet having a tensile strength of 980 MPa or more is required.

  On the other hand, many automotive parts made of steel sheets are manufactured by performing complex molding by pressing, burring, etc., so high-strength steel sheets for automotive parts have strength, ductility, It is required to have excellent workability such as stretch flangeability.

  In addition, it is necessary to consider that the automobile is used in various environments. For example, in a low-temperature environment such as a cold district, there is a possibility that the automobile parts may be brittlely broken. Therefore, as a high-strength steel sheet for automobile parts, there is a demand for a high-strength steel sheet with excellent toughness to be compatible with use in cold regions.

  As described above, a high-strength steel sheet as a material for automobile parts is required to have various properties such as workability and toughness in addition to high strength. Under such circumstances, various steel plates have been studied so far.

  For example, in Patent Document 1, the content and form of sulfide are controlled by defining the contents of Ti and S so that (Ti / 48) / (S / 32) is 30 or more. By controlling the texture so that the {211} plane X-ray random strength ratio is 2.4 or less, a high-strength hot-rolled steel sheet with a tensile strength of 780 MPa or more and excellent hole expansibility and fracture characteristics is obtained. Technology is disclosed. Also, in Patent Document 2, the tensile strength is 590 MPa or more by the texture control by refining the crystal grains and further making the X-ray random intensity ratio of the {211} <011> orientation parallel to the rolling direction 2.5 or less. A technique for producing a high-strength hot-rolled steel sheet having excellent low-temperature toughness and hole expansibility is disclosed.

  Furthermore, in Patent Document 3, the total grain boundary number density of solute C and solute B, the cementite grain size precipitated at the grain boundaries in the steel sheet, and the average crystal grain size at the center of the sheet thickness are controlled. In addition, by controlling the texture so that the {211} random strength ratio at the center of the plate thickness is 2 or less, and further controlling the size and density of TiC precipitates in the crystal grains, the tensile strength can be reduced. A technique for producing a high-strength hot-rolled steel sheet of 780 MPa or more and excellent burring properties is disclosed. Further, in Patent Document 4, a high strength thin film having excellent elongation and local ductility is obtained by forming a mixed structure mainly composed of precipitation strengthened ferrite and containing residual austenite, and precipitating carbonitride in the ferrite by interphase precipitation. A technique of making a steel plate (hot rolled steel plate) is disclosed.

  All of these technologies utilize the fine precipitation of Ti and Nb to achieve both the strength and workability of hot-rolled steel sheets.

Japanese Patent No. 4842413 JP 2012-136773 A JP 2012-001776 A JP 2011-225941 A

  However, none of the techniques disclosed in Patent Document 1, Patent Document 2 and Patent Document 3 has obtained a hot-rolled steel sheet having a tensile strength of 980 MPa or more. On the other hand, according to the technique disclosed in Patent Document 4, the tensile strength of the hot-rolled steel sheet can be set to 980 MPa or more. However, in the technique disclosed in Patent Document 4, the tensile strength cannot always be set to 980 MPa or more as shown in the examples, and may be significantly lower than 980 MPa. Further, in the technique disclosed in Patent Document 4, there is no room for improvement because the stretch flangeability and low temperature toughness of the hot-rolled steel sheet are not studied at all.

  As described above, there is no known technique that can achieve both high workability and good low-temperature toughness for hot-rolled steel sheets having a tensile strength of 980 MPa or more.

  The present invention has been made in view of the above circumstances, has a tensile strength of 980 MPa or more, good workability such as stretch flangeability and ductility (elongation), and excellent low temperature toughness. An object is to provide a high-strength hot-rolled steel sheet and a method for producing the same.

  In order to solve the above-mentioned problems, the present inventors diligently investigated means for improving workability and low-temperature toughness while ensuring a tensile strength of 980 MPa or more for hot-rolled steel sheets.

  It is known that both the strength and workability (particularly ductility and stretch flangeability) of a steel sheet can be achieved by using ferrite as the main phase structure of the steel sheet and utilizing precipitation strengthening by fine precipitates. Therefore, the inventors first tried to increase the strength of the steel sheet to a tensile strength of 980 MPa or more by utilizing fine precipitates.

  As a result, it was found that when Ti and V are used as the precipitation elements, Ti and V are combined, and precipitate in the steel finely and with high density, thereby obtaining extremely high steel sheet strength. And, the content of C, Ti, V, which is a carbide forming element, is set within a predetermined range, the amount of dissolved Ti in the hot-rolled steel sheet is suppressed, and further, fine carbides containing Ti and V are precipitated. It was found that a steel plate strength of 980 MPa or more was obtained. Furthermore, for the hot-rolled steel sheet with a tensile strength of 980 MPa or more obtained as described above, the precipitation form of fine carbides containing Ti and V was investigated, and the shape when observed from the [001] direction of the ferrite phase was It was confirmed that the film had a substantially rectangular shape with an average length of 8 nm or less and an average width of 1 nm or less.

  On the other hand, since precipitation strengthening usually degrades the toughness of steel, it is generally difficult to achieve both strength and toughness of a steel sheet by precipitation strengthening. And when the carbide | carbonized_material containing Ti and V was deposited as mentioned above and the tensile strength of the steel plate was made into 980 MPa or more, the tendency for toughness to fall was shown.

  Based on these matters, the present inventors have further studied to improve the toughness of precipitation strengthened steel. As a result, for Ti and V in steel, when the amount of Ti that can be precipitated as carbide is larger than the amount of V that can be precipitated as carbide, that is, the addition amount of the alloy element is ((Ti / 48)-(N / 14)-(S / 32)) / (V / 51)> 1.0 (Ti, V, N, S is the content of each element in steel (mass%)) It was found that it improved greatly.

  Moreover, in order to clarify the reason why the low temperature toughness is greatly improved, the structure of the hot rolled steel sheet having a main phase structure of ferrite having various chemical compositions was observed. As a result, the Ti, V, N, and S contents (mass%) of the hot-rolled steel sheet satisfy ((Ti / 48)-(N / 14)-(S / 32)) / (V / 51)> 1.0. When it does, it became clear that a ferrite particle size becomes small. That is, it was confirmed that when the Ti, V, N, and S contents (mass%) of the hot-rolled steel sheet satisfy the above formula, the ferrite structure is refined and the toughness of the hot-rolled steel sheet is improved.

  When the Ti, V, N, and S contents (mass%) of the hot-rolled steel sheet satisfy ((Ti / 48)-(N / 14)-(S / 32)) / (V / 51)> 1.0 The reason why the ferrite structure becomes finer is not clear. However, Ti is easier to form carbides than V, so if the amount of Ti that can be precipitated as carbide is larger than the amount of V that can be precipitated as carbide, V during the cooling process after finish rolling is finished during hot-rolled steel sheet production. It is presumed that Ti, which did not form a composite carbide, segregates at the austenite grain boundaries to form a carbide, and the carbide easily becomes a preferential nucleation site of ferrite.

  As described above, the toughness of the hot-rolled steel sheet can be improved by defining the Ti and V contents in the hot-rolled steel sheet according to the above formula. However, as the Ti and V contents are regulated in this way, it has become a new problem that it becomes difficult to ensure the stability of hot-rolled steel sheet strength.

  Fine carbides (carbides containing Ti and V) that contribute to increasing the strength of hot-rolled steel sheets precipitate during the hot-rolled steel sheet manufacturing and cooling and winding processes after finishing rolling (mainly during austenite → ferrite transformation) To do. In order to increase the strength of the hot-rolled steel sheet, it is necessary to maintain the fine carbide thus precipitated in a fine state even after the hot-rolled steel sheet is wound into a coil. However, if the Ti and V contents in the hot-rolled steel sheet are specified according to the above formula, the size of carbides precipitated in the hot-rolled steel sheet is likely to be coarsened after winding, and the desired hot-rolled steel sheet strength is not necessarily obtained. Not found out.

  In response to the above problems, the present inventors have made Ti, V of hot-rolled steel sheet to satisfy ((Ti / 48)-(N / 14)-(S / 32)) / (V / 51)> 1.0. Even when N, S content (% by mass) is specified, a means for maintaining the strength stability of the hot-rolled steel sheet and ensuring a tensile strength of 980 MPa or more was sought. As a result, for C, Ti, and V in steel, the amount of C that can be precipitated as carbide is larger than the total amount of V and Ti that can be precipitated as carbide, that is, the amount of alloying element added is (C / 12) / ((Ti / 48) − (N / 14) − (S / 32) + (V / 51))> 1.0 (C, Ti, V, N, S are the contents of each element in steel ( It has been found that securing 0.022% or more by mass% of C that satisfies the mass%)) and not precipitated as a carbide is an extremely effective means. In addition, if these conditions are satisfied, the carbide precipitated in the hot-rolled steel sheet is refined regardless of the carbide composition, and a high-strength hot-rolled steel sheet with a tensile strength of 980 MPa or more can be stably obtained. I found out. It was also found that the toughness of the hot-rolled steel sheet is not impaired even when the tensile strength is as high as 980 MPa or higher.

  The C, Ti, V, N, and S contents (% by mass) of the hot-rolled steel sheet are (C / 12) / ((Ti / 48) − (N / 14) − (S / 32) + (V / If 51))> 1.0 is satisfied, the reason why the carbide precipitated in the hot-rolled steel sheet becomes finer is not necessarily clear. However, if the mole fraction of C that can be precipitated as carbide is higher than the mole fraction of Ti and V that can precipitate as carbide, C concentrates on the surface of the carbide containing Ti and V, resulting in the size of the carbide. It is estimated that the stability is improved.

In addition to Ti and V, the present inventors have changed Mo, Nb, and W to the following formulas (1) and (2) ((Ti / 48) − (N / 14) − (S / 32) + ( Mo / 96) + (Nb / 93) + (W / 184)) / (V / 51)> 1.0… (1)
(C / 12) / ((Ti / 48)-(N / 14)-(S / 32) + (V / 51) + (Mo / 96) + (Nb / 93) + (W / 184))> 1.0 (2)
It has been newly found out that, when contained in such a manner that fine carbides are precipitated, the strength can be stably improved, and the ferrite grain size can be reduced to improve the toughness.

The present invention has been completed based on the above findings, and the gist thereof is as follows .

[1 ] A steel material is heated, subjected to hot rolling consisting of rough rolling and finish rolling, and after finishing rolling, cooled, wound, and made into a hot-rolled steel sheet, the steel material in mass%,
C: 0.07% or more and 0.12% or less, Si: 0.2% or less,
Mn: 0.1% to 1.2%, P: 0.025% or less,
S: 0.005% or less, Al: 0.1% or less,
N: 0.008% or less, Ti: 0.14% or more and 0.20% or less,
V: 0.05% or more and 0.15% or less, C, Ti, V, S and N are contained so as to satisfy the following formulas (1) and (2), and the balance is composed of Fe and inevitable impurities,
The heating temperature of the heating is 1150 ° C. or more and 1350 ° C. or less, the finish rolling finish temperature of the finish rolling is 880 ° C. or more, the winding temperature of the winding is 550 ° C. or more and 700 ° C. or less, and the heat after winding The coil is water-cooled ,
The hot-rolled steel sheet has the above-mentioned composition, and solute Ti is 0.04% or less by mass%, and C not precipitated as carbide is 0.022% or more by mass%,
The area ratio of the ferrite phase with respect to the entire structure is 95% or more, the average crystal grain size of the ferrite phase is 5.0 μm or less, and is a carbide containing Ti and V, which was observed from the [001] direction of the ferrite phase. When the shape has a structure in which carbide having a rectangular shape with an average length of 8 nm or less and an average width of 1 nm or less is precipitated,
A method for producing a high-strength hot-rolled steel sheet excellent in workability and toughness, characterized by having a tensile strength of 980 MPa or more .
((Ti / 48)-(N / 14)-(S / 32) + (Mo / 96) + (Nb / 93) + (W / 184)) / (V / 51)> 1.0 (1)
(C / 12) / ((Ti / 48)-(N / 14)-(S / 32) + (V / 51) + (Mo / 96) + (Nb / 93) + (W / 184))> 1.0 (2)
(Here, C, Ti, V, S, N, Nb, Mo, W: content of each element (mass%))

[ 2 ] In [ 1 ], in place of the winding and hot-rolled steel sheet, after forming the wound and hot-rolled steel sheet, further forming the plating layer on the surface of the hot-rolled steel sheet, the winding of the winding a temperature of 400 ° C. or higher 700 ° C. or less, and water-cooled hot-rolled coil after the winding, subjected to continuous annealing and plating, after the plating process, an average cooling rate: 20 ° C. / s of 0.99 ° C. or less or more A method for producing a high-strength hot-rolled steel sheet, characterized by cooling to a temperature.
[ 3 ] In [ 2 ], following the plating treatment, alloying treatment of the plating layer is performed, and after the alloying treatment, cooling is performed at an average cooling rate of 20 ° C./s to 150 ° C. A method for producing a high-strength hot-rolled steel sheet.
[ 4 ] In any one of [ 1 ] to [ 3 ], in addition to the above composition, in addition to mass, Mo: 0.02% to 0.30%, Nb: 0.02% to 0.10%, W: 0.02% to 0.10 % Or less. A method for producing a high-strength hot-rolled steel sheet, characterized in that the composition contains one or more selected from the following.

  According to the present invention, excellent workability (ductility, stretch flangeability) suitable for automobile steel sheets and the like, which has a tensile strength (TS) of 980 MPa or more and can be formed into a complicated shape by pressing or burring. In addition, a high-strength hot-rolled steel sheet having a low-temperature toughness sufficient to suppress brittle fracture that is a concern in a low-temperature environment such as a cold region is obtained. Specifically, a hot-rolled steel sheet having a tensile strength of 980 MPa or more, a total elongation of 15% or more, a hole expansion ratio of 40% or more, and a ductile brittle transition temperature of −20 ° C. or less is obtained. The effect of.

In addition, the high-strength hot-rolled steel sheet according to the present invention maintains the above effects unless heat treatment that significantly breaks the structure (for example, heat treatment for a long time in a temperature range of Ac ₁ point or higher) is performed. Can do. That is, the same effect as described above can be obtained also in a plated steel sheet using the high-strength hot-rolled steel sheet of the present invention as an original sheet.

  Hereinafter, the present invention will be described in detail.

  First, the reason for limiting the structure of the hot-rolled steel sheet of the present invention will be described.

  The hot-rolled steel sheet of the present invention has a ferrite phase area ratio of 95% or more with respect to the entire structure, the ferrite crystal has an average crystal grain size of 5.0 μm or less, and is a carbide containing Ti and V, wherein the ferrite The phase as observed from the [001] direction of the phase has a structure in which carbides having a rectangular shape with an average length of 8 nm or less and an average width of 1 nm or less are precipitated.

Ferrite phase: 95% or more in area ratio with respect to the entire structure In order to secure the elongation characteristics of the steel sheet, it is essential to form a ferrite phase having a larger elongation than the bainite phase or the martensite phase. Moreover, in order to improve the stretch flangeability of a steel plate, it is effective to make the steel plate structure a single phase. Therefore, in the present invention, it is desirable that the hot-rolled steel sheet structure is a ferrite single phase. However, even if it is not a complete ferrite single phase, if the ferrite single phase is substantially a ferrite phase, that is, if the area ratio to the whole structure is 95% or more, the above effect is sufficiently exhibited. Therefore, the area ratio of the ferrite phase to the entire structure is 95% or more.

  In the hot-rolled steel sheet of the present invention, examples of the structure other than the ferrite phase include martensite, pearlite, bainite, and grain boundary cementite. These totals are acceptable if the area ratio to the whole tissue is about 5% or less.

Average crystal grain size of ferrite: 5.0 μm or less The crystal grain size of the ferrite phase affects the toughness of the hot-rolled steel sheet. In order to ensure the toughness of the hot-rolled steel sheet, it is effective to set the average crystal grain size of ferrite to 5.0 μm or less. Preferably it is 4.0 micrometers or less.

Carbides containing Ti and V In order to increase the strength of hot-rolled steel sheets while maintaining excellent elongation characteristics and workability such as stretch flangeability, the main phase of the hot-rolled steel sheets is the ferrite phase as described above, and the ferrite phase It is effective to strengthen the material by precipitation of fine carbides. Note that the finer the carbide, the greater the number of particles that inhibit the movement of dislocations when the hot-rolled steel sheet is loaded, so the amount of strengthening obtained by precipitating the carbide increases.

  Here, in the present invention, in order to make the tensile strength of the hot rolled steel sheet 980 MPa or more, it is necessary to precipitate fine carbides containing Ti and V in the ferrite phase. The reason is not clear, but carbides composed only of Ti and carbides composed only of V are easily coarsened, and it is difficult to stably secure the steel plate strength. Therefore, in this invention, the carbide | carbonized_material containing Ti and V is applied as a carbide | carbonized_material for aiming at precipitation strengthening. In addition, examples of the “carbide containing Ti and V” referred to here include Ti—V based composite carbide, and composite carbide containing other carbide forming elements (Mo, Nb, W, etc.) together with Ti and V. It is done.

Shapes and dimensions of carbides containing Ti and V In the present invention, carbides (fine carbides containing Ti and V) that contribute to precipitation strengthening are plate-like carbides, and the ferrite phase that is the parent phase of hot-rolled steel sheets. Precipitate with a specific orientation relationship. Therefore, the shape and size may not be correctly determined depending on the orientation in which the carbide is observed.

  In the present invention, the plate-like carbide contributing to precipitation strengthening is such that the plate surface of the plate is parallel to the [001] direction of the ferrite phase (that is, the normal of the plate surface of the plate is perpendicular to the [001] direction of the ferrite phase). To precipitate). Therefore, when the carbide contributing to precipitation strengthening in the present invention is observed from the [001] direction of the ferrite phase, the shape is substantially rectangular. And the thickness dimension of the said plate-shaped carbide | carbonized_material substantially corresponds with the dimension of the substantially rectangular width | variety (The shorter side of 2 sets of opposing sides of a substantially rectangular shape). Thus, in the present invention, the carbide contributing to precipitation strengthening can be accurately determined by observing from the [001] direction of the ferrite phase.

  Further, the finer the carbide, the higher the amount of precipitation strengthening is obtained, and particularly when the shape when observed from the [001] direction of the ferrite phase is a rectangular shape having an average length of 8 nm or less and an average width of 1 nm or less, The steel sheet is remarkably increased in strength. Therefore, in the present invention, the carbides precipitated on the hot-rolled steel sheet, the shape when observed from the predetermined direction is a carbide (including Ti and V) having an average length of 8 nm or less and an average width of 1 nm or less. Carbide). By defining in this way, the tensile strength of the hot-rolled steel sheet can be set to 980 MPa or more.

  In addition, in order to increase the strength of the hot-rolled steel sheet, not only the size of carbides precipitated in the ferrite phase but also the amount of precipitation is important, and the precipitation strengthening amount increases as more fine carbides are precipitated. As will be described later, in the present invention, the content of C, Ti, V, which are carbide forming elements, is specified, and the Ti amount contributing to the formation of carbide is ensured by limiting the amount of Ti dissolved in the steel. Thus, a fine carbide precipitation amount sufficient to make the tensile strength of the hot-rolled steel sheet 980 MPa or more is obtained.

  Next, the reason for limiting the component composition of the hot-rolled steel sheet of the present invention will be described. In addition,% showing the following component composition shall mean the mass% unless there is particular notice.

C: 0.07% to 0.12%
C is Ti and V, etc., and nano-sized carbides (plate-like carbides having a rectangular shape with an average length of 8 nm or less and an average width of 1 nm or less when observed from the [001] direction of the ferrite phase. Is an important element that contributes to precipitation strengthening. In order to set the tensile strength of the hot-rolled steel sheet to 980 MPa or more, the C content needs to be 0.07% or more. On the other hand, if the C content exceeds 0.12%, during production of hot-rolled steel sheets, during the cooling process after finish rolling, coarse TiC precipitates before winding and stretch flangeability of the finally obtained hot-rolled steel sheets Decreases. Therefore, the C content is 0.07% or more and 0.12% or less.

Si: 0.2% or less
Si has the effect of strengthening the steel by solid solution strengthening, but at the same time causes a reduction in the toughness of the steel. Therefore, the Si content is 0.2% or less. Preferably it is 0.1% or less.

Mn: 0.1% or more and 1.2% or less
Mn has an effect of lowering the austenite → ferrite transformation temperature, and thus is an element that contributes to improving toughness by refining the crystal grains of the hot-rolled steel sheet. As will be described later, the hot-rolled steel sheet of the present invention is manufactured by subjecting it to finish rolling in the austenite region, and transforming austenite to ferrite in the cooling and winding process after finish rolling to make a ferrite single phase structure substantially. The Therefore, when the austenite → ferrite transformation temperature is low, the finish rolling finish temperature can be lowered, and the growth of austenite grains during hot rolling can be suppressed. In addition, the size of the ferrite grains obtained by the transformation of austenite to ferrite in the cooling and winding process after finish rolling depends on the size of the austenite grains before transformation, and the austenite grains before transformation are fine. The crystal grains of the ferrite phase obtained by the austenite → ferrite transformation are also refined.

  In order to obtain such an effect, the Mn content needs to be 0.1% or more. On the other hand, if the Mn content exceeds 1.2%, Mn segregates in the center of the plate thickness to form a heterogeneous structure, thereby reducing the stretch flangeability of the hot-rolled steel sheet. Therefore, the Mn content is 0.1% or more and 1.2% or less. Preferably they are 0.5% or more and 1.1% or less.

P: 0.025% or less
P is a solid solution strengthening element, but when the P content exceeds 0.025%, segregation becomes prominent, leading to a reduction in stretch flangeability and toughness of the hot-rolled steel sheet. Therefore, the P content is 0.025% or less. In addition, although it is preferable to reduce P content from a viewpoint of a hot-rolled steel plate characteristic, reducing P content extremely will cause the manufacturing cost to rise. Regarding the P content, a practical lower limit value that does not significantly increase the manufacturing cost of the hot-rolled steel sheet is about 0.001%.

S: 0.005% or less
S forms sulfides with Mn and Ti, and causes a decrease in workability (elongation and stretch flangeability) of the hot-rolled steel sheet. Therefore, in the present invention, it is desirable to reduce the S content as much as possible, and the content is made 0.005% or less. However, extremely reducing the S content causes an increase in manufacturing costs. With respect to the S content, a practical lower limit that does not significantly increase the manufacturing cost is about 0.0005%.

Al: 0.1% or less
Al is an element that acts as a deoxidizer. However, if the Al content exceeds 0.1%, a large amount of inclusions are generated, leading to elongation of the hot-rolled steel sheet and deterioration of stretch flangeability. Therefore, the Al content is 0.1% or less. Preferably it is 0.05% or less. In addition, although there is no lower limit in particular of Al content, in order to fully acquire the effect as a deoxidizer, it is preferable to make Al content 0.01% or more.

N: 0.008% or less
N is a harmful element in the present invention and is desirably reduced as much as possible. In particular, if the N content exceeds 0.008%, the stretch flangeability of the hot-rolled steel sheet decreases due to the formation of coarse nitrides in the steel. Therefore, the N content is 0.008% or less.

Ti: 0.14% to 0.20%
Ti is one of the important elements in the present invention. Ti, together with V, is a nano-sized carbide, that is, a plate-like carbide, and has a rectangular shape with an average length of 8 nm or less and an average width of 1 nm or less when observed from the [001] direction of the ferrite phase. A certain carbide is formed and contributes to precipitation strengthening of the parent phase (substantially a ferrite single phase in the present invention). In addition, Ti that does not form composite carbide with V, etc., mainly segregates at austenite grain boundaries during the hot rolling production and in the cooling process after finish rolling to form carbides, but these carbides are ferrite nucleation sites. This contributes to the refinement of ferrite grains. And with the refinement | miniaturization of a ferrite grain, the intensity | strength and toughness of a hot-rolled steel plate improve.

  In order to exhibit such an effect and secure desired hot-rolled steel sheet strength (tensile strength of 980 MPa or more) and toughness, the Ti content needs to be 0.14% or more. On the other hand, if the Ti content exceeds 0.20%, the carbides are coarsened and the shape tends to be spherical, and the hot-rolled steel sheet strength and stretch flangeability are lowered. Therefore, the Ti content is 0.14% or more and 0.20% or less. Preferably it is 0.15% or more and 0.18% or less.

V: 0.05% or more and 0.15% or less
V is one of the important elements in the present invention. As described above, V forms nano-sized carbides together with Ti and the like and contributes to precipitation strengthening of the parent phase. In order to exhibit such an effect and secure desired hot-rolled steel sheet strength (tensile strength of 980 MPa or more), the V content needs to be 0.05% or more. On the other hand, if the V content exceeds 0.15%, the ferrite grain size increases and the toughness of the hot-rolled steel sheet decreases. Therefore, the V content is 0.05% or more and 0.15% or less. Preferably they are 0.08% or more and 0.14% or less.

  The above-mentioned components are basic components, but one type selected from Mo: 0.02% to 0.30%, Nb: 0.02% to 0.10%, W: 0.02% to 0.10% as necessary Or 2 or more types can be selected and contained.

Mo: 0.02% or more and 0.30% or less, Nb: 0.02% or more and 0.10% or less, W: 0.02% or more and 0.10% or less
Mo, Nb, and W are all elements that contribute to strengthening the parent phase by forming nano-sized composite carbides together with Ti and V, as well as contributing to toughness improvement by making ferrite grains finer. Can be selected and contained. In order to secure such an effect, it is preferable to set the Mo content: 0.02% or more, the Nb content: 0.02% or more, and the W content: 0.02% or more. On the other hand, if the Mo content: 0.30%, the Nb content: 0.10%, and the W content: 0.10%, ductility tends to decrease. For this reason, when it contains, it is preferable to limit Mo content to 0.02% or more and 0.30% or less, Nb content to 0.02% or more and 0.10% or less, and W content to 0.02% or more and 0.10% or less.

The hot-rolled steel sheet of the present invention contains C, Ti, V, S, and N, or Mo, Nb, and W in the above-described range and satisfying the expressions (1) and (2).
((Ti / 48) − (N / 14) − (S / 32) + (Mo / 96) + (Nb / 93) + (W / 184)) / (V / 51)> 1.0 (1)
(C / 12) / ((Ti / 48)-(N / 14)-(S / 32) + (V / 51) + (Mo / 96) + (Nb / 93) + (W / 184))> 1.0 (2)
(Here, C, Ti, V, S, N, Mo, Nb, W: content of each element (mass%))
The above formulas (1) and (2) are requirements that must be satisfied in order to increase the strength and toughness of hot-rolled steel sheets by refining each of carbides and ferrite grains containing Ti and V, and are extremely important in the present invention. It is a good index. In addition, when calculating the value on the left side of the above formulas (1) and (2), the elements that are not included among the elements displayed in the formula are calculated as zero.

((Ti / 48) − (N / 14) − (S / 32) + (Mo / 96) + (Nb / 93) + (W / 184)) / (V / 51)> 1.0 (1)
As a means for improving the toughness of a hot-rolled steel sheet, it is effective to make ferrite grains finer. Therefore, in the present invention, the amount of Ti that can precipitate as carbide ((Ti / 48)-(N / 14)-(S / 32)) rather than the amount of V that can precipitate as carbide (V / 51), When Mo, Nb, and W are contained, the total amount of Ti, Mo (Mo / 96), Nb (Nb / 93), and W (W / 184) is increased. That is, in the present invention, ((Ti / 48) − (N / 14) − (S / 32) + (Mo / 96) + (Nb / 93) + (W / 184)) / (V / 51) Needs to be more than 1.0, and is preferably 1.1 or more. Thereby, the ferrite grain of a hot-rolled steel sheet is refined | miniaturized, a brittle transition temperature falls to -20 degrees C or less, and toughness can be improved greatly. However, the value of ((Ti / 48)-(N / 14)-(S / 32) + (Mo / 96) + (Nb / 93) + (W / 184)) / (V / 51) is excessive. If it becomes large, the carbide tends to be coarsened and the strength of the hot-rolled steel sheet may be lowered.
(C / 12) / ((Ti / 48)-(N / 14)-(S / 32) + (V / 51) + (Mo / 96) + (Nb / 93) + (W / 184))> 1.0 (2)
As described above, the toughness of the hot-rolled steel sheet is improved only by setting each of C, V, Ti, or Mo, Nb, and W to a predetermined content and satisfying the formula (1). , Carbides containing Ti and V are likely to be coarsened, and the tensile strength may be less than 980 MPa. Therefore, in the present invention, in addition to the above, the contents of C, Ti, V, N, S, or Mo, Nb, and W are further defined so as to satisfy the expression (2), as will be described later. By adding 0.022% or more of C that has not been precipitated as carbides, the strength of the hot-rolled steel sheet is stabilized.

  In the present invention, Ti and V that can be precipitated as carbides, or Ti and V are contained by increasing the molar fraction of C that can be precipitated as carbides, rather than the sum of the molar fractions of Mo, Nb, and W. Stable precipitation of carbide in the desired form (carbide that is a plate-like carbide and has a rectangular shape with an average length of 8 nm or less and an average width of 1 nm or less when observed from the [001] direction of the ferrite phase) Can be made. Therefore, in the present invention, (C / 12) / ((Ti / 48) − (N / 14) − (S / 32) + (V / 51) + (Mo / 96) + (Nb / 93) + ( W / 184)) must be greater than 1.0. Preferably it is 1.1 or more. Thereby, a hot-rolled steel sheet having a tensile strength of 980 MPa or more can be realized stably. However, (C / 12) / ((Ti / 48)-(N / 14)-(S / 32) + (V / 51) + (Mo / 96) + (Nb / 93) + (W / 184) When the value of) is excessively high, C forms grain boundary cementite and the stretch flangeability of the hot-rolled steel sheet may be lowered.

  As will be described later, when producing the hot-rolled steel sheet of the present invention, the steel material is heated to the austenite region during hot rolling, but Ti in the steel is easily N or S in the temperature range of 1300 ° C or lower. Precipitates by chemical bonding. Therefore, in the above formulas (1) and (2), the amount of Ti that can be precipitated as carbides is the amount of Ti / 48 to N bonded (N / 14) and S bonded (S / 32). The value was subtracted.

Solid solution Ti: 0.04% or less In order to increase the strength of a hot-rolled steel sheet, it is important not only to refine carbides precipitated in the ferrite phase, but also to ensure a sufficient amount of precipitation. In order to obtain a sufficient precipitation strengthening amount, the contents of C, Ti, and V, which are the precipitate forming elements, are respectively C: 0.07% or more and 0.12% or less, Ti: 0.14% or more and 0.20% or less, V: 0.05% or more It is necessary to limit it to 0.15% or less, and to limit Ti existing in a solid solution state to 0.04% or less in the total amount of Ti contained in the hot-rolled steel sheet. Thus, the tensile strength of 980 MPa or more can be ensured by restricting the amount of solid solution Ti and precipitating the remaining Ti as carbides that contribute to improving the strength and toughness of the hot-rolled steel sheet. In the present invention, it is preferable to reduce the amount of dissolved Ti as much as possible.

C not precipitated as carbide: 0.022% or more The reason is not clear, but C not precipitated as carbide is a coarsening of carbides (fine carbides containing Ti and V) that contributes to increasing the strength of hot-rolled steel sheets. It has a suppressing effect. Therefore, in the present invention, C not precipitated as carbide is 0.022% or more. As a result, the carbide containing Ti and V is in a desired form (a plate-like carbide having a rectangular shape with an average length of 8 nm or less and an average width of 1 nm or less when observed from the [001] direction of the ferrite phase. Can be stably precipitated.

  Further, C that is not precipitated as a carbide increases the grain boundary strength by segregating at the grain boundary, and contributes to increasing the strength of the hot-rolled steel sheet and also has an effect of improving toughness. Also from such a viewpoint, it is preferable that C not precipitated as carbide is 0.022% or more.

  In addition, when C which has not precipitated as carbides becomes excessive, the amount of precipitation of carbides that contribute to increasing the strength of the hot-rolled steel sheet decreases. Therefore, C not precipitated as carbide is preferably 0.050% or less.

  In the hot rolled steel sheet of the present invention, the balance other than the above is Fe and inevitable impurities. Inevitable impurities include Sb, Cu, Ni, As, Sn, and Pb. If the total content thereof is 0.2% or less, the above-described effects of the present invention are not affected. In addition, it is considered that Zr, Ta, Cr, Co, Se, Zn or Ca, REM, Mg, Cs is contained, but if the total content of different elements or more of the above elements is 1.0% or less, The above-described effects of the present invention are not affected.

  It is good also as a steel plate (plated steel plate) which forms a plating layer in the surface of an above-mentioned hot-rolled steel plate, and has a plating layer. In addition, as a plating layer, a hot dip galvanization layer and an alloying hot dip galvanization layer can be illustrated.

  Next, the manufacturing method of the hot rolled steel sheet of the present invention will be described.

  In the present invention, a steel material having a predetermined 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 steel material is set to 1150 ° C or higher and 1350 ° C or lower, the finish rolling finish temperature is set to 880 ° C or higher, the winding temperature is set to 550 ° C or higher and 700 ° C or lower, and the hot-rolled coil after winding is water-cooled. Features.

  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 used. Moreover, after melting, it is preferable to use a slab (steel material) by a continuous casting method because of problems such as segregation, but the slab (steel) can be obtained by a known casting method such as an ingot-bundling rolling method or a thin slab continuous casting method. Material).

Heating temperature: 1150 ° 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 to an austenite single phase region before rough rolling. It is necessary to dissolve the carbide in the steel material. In the present invention containing Ti and V or further Mo, Nb, and W as carbide forming elements, it is necessary to reheat the steel material to a temperature of 1150 ° C. or higher. On the other hand, if the heating temperature of the steel material is extremely high, exceeding 1350 ° C., the energy required for heating and the load on the equipment increase. Therefore, the heating temperature of the steel material is set to 1350 ° C or less.

  However, if the steel material before casting and before rough rolling maintains a temperature above the specified temperature, and the carbide in the steel material is sufficiently dissolved, the step of heating the steel material before rough rolling can be omitted. And direct rolling may be performed. The rough rolling conditions are not particularly limited.

Finish rolling end temperature: 880 ° C or higher If the finish rolling end temperature is less than 880 ° C, recrystallization does not occur and rolling is performed in the non-recrystallization temperature range, so the rolling load increases significantly and hot rolling becomes difficult. Become. Accordingly, the finish rolling finish temperature is set to 880 ° C. or higher. Preferably it is 890 degreeC or more. On the other hand, if the finish rolling finish temperature is excessively high, the ferrite grain size of the finally obtained hot-rolled steel sheet becomes coarse, which causes a decrease in toughness. Therefore, the finish rolling finish temperature is desirably 1020 ° C. or lower.

Winding temperature: 550 ° C or higher and 700 ° C or lower Carbide (fine carbides containing Ti and V) that can contribute to increasing the strength of hot-rolled steel sheets is mainly precipitated in the winding process.・ The form changes according to the coiling temperature. Therefore, optimization of the coiling temperature is important in the present invention.

  When the coiling temperature exceeds 700 ° C., carbides containing Ti and V become large, and the shape when observed from the [001] direction of the ferrite phase is a plate-like form having an average length of more than 8 nm or an average width of more than 1 nm, Or it becomes a spheroid. Therefore, the tensile strength of the hot rolled steel sheet cannot be made 980 MPa or more. On the other hand, when the coiling temperature is less than 550 ° C., the solid solution Ti in the hot-rolled steel sheet exceeds 0.04%, the precipitation amount of carbide containing Ti and V is insufficient, and the tensile strength of 980 MPa or more cannot be secured. Therefore, the winding temperature is set to 550 ° C. or more and 700 ° C. or less. Preferably they are 570 degreeC or more and 670 degrees C or less.

  The average cooling rate from the finish rolling finish temperature to the coiling temperature after finish rolling is 15 ° C / s or more and 50 ° C / s or less and normal water-cooling conditions have little effect on the steel structure. preferable. Further, the finish rolling end temperature and the coiling temperature are temperatures on the steel sheet surface.

Hot-rolled coil after winding: Water-cooling The hot-rolled steel sheet of the present invention is intended to improve toughness by containing 0.022% or more of C not precipitated as carbide. However, the ferrite which is the main phase of the hot rolled steel sheet of the present invention hardly dissolves carbon. Therefore, if the hot-rolled coil after winding is allowed to cool or air cool, C does not contribute to the formation of carbides (carbides containing Ti and V, or carbides that contribute to the refinement of ferrite grains). Before cooling, it combines with Fe and precipitates as cementite.

  As described above, when the hot-rolled coil after winding is allowed to cool or air cool, C not precipitated as carbide cannot be made 0.022% or more. Therefore, in the present invention, it is necessary to water-cool the wound hot rolled coil. As a method of water-cooling the hot-rolled coil, for example, there is a method of immersing the hot-rolled coil in a water tank attached to the hot rolling line. The water cooling of the hot-rolled coil is desirably performed immediately after winding, and is preferably water-cooled within 30 minutes after winding. Moreover, it is desirable to perform water cooling of the hot-rolled coil until the temperature of the entire hot-rolled coil becomes 200 ° C. or less.

  In the present invention, instead of using the above-described winding and hot-rolled steel plate, a plating process for forming a plating layer on the surface of the hot-rolled steel plate may be performed after winding and hot-rolling steel plate. The plating layer is preferably a hot dip galvanizing layer, but it is needless to say that the plating layer is not limited thereto. Hereinafter, a case where the plated layer is a hot dip galvanized layer will be described.

  When forming a plating layer (hot dip galvanized layer) on the surface of the hot rolled steel sheet after winding and hot rolling steel sheet, the winding temperature of the winding should be 400 ° C or higher and 700 ° C or lower. The water cooling of the hot rolled coil may or may not be performed.

Winding temperature: 400 ° C or higher and 700 ° C or lower If the winding temperature is lower than 550 ° C, the amount of fine carbides that contribute to strengthening in the hot-rolled steel sheet will be insufficient and the strength will be reduced. When heating a hot-rolled steel sheet by continuous annealing performed after winding, the precipitation amount of fine carbides increases and the strength increases. For this reason, when performing a plating process, coiling temperature is good also as 400 degreeC or more. When the coiling temperature is less than 400 ° C, a bainite phase and a martensite phase containing a high dislocation density are formed, and elongation and hole expansibility decrease. A bainite phase is generated even at 400 ° C or higher and lower than 550 ° C, but the dislocation density of the bainite phase generated in this temperature range decreases due to continuous annealing. For this reason, elongation and hole expansibility are not reduced. On the other hand, when the coiling temperature is higher than 700 ° C., the carbide becomes coarse and the strength decreases. In addition, when it is allowed to cool after winding, cementite precipitates and the amount of C not precipitated as carbide is less than 0.022%. However, even when precipitated as cementite, the cementite dissolves during heating during continuous annealing. Therefore, the water cooling treatment after winding may or may not be performed.

  After winding, pickling is performed to remove the surface scale, and it is preferable to continuously perform a continuous annealing treatment and a plating treatment, preferably using a conventional continuous annealing apparatus. The continuous annealing treatment is performed at an annealing temperature of 600 ° C. or higher and 750 ° C. or lower.

Annealing temperature: 600 ° C or more and 750 ° C or less If the annealing temperature of continuous annealing treatment is less than 600 ° C, the cementite precipitated during winding or the cementite precipitated during heating does not dissolve and the amount of C not precipitated as carbide is 0.022% Less than. On the other hand, if the annealing temperature is higher than 750 ° C., the fine carbides become coarse and the strength decreases. For this reason, it is preferable to limit the annealing temperature of a continuous annealing process to 600 degreeC or more and 750 degrees C or less.

  The plating treatment is preferably a conventional hot dip galvanizing treatment. The steel sheet that has been subjected to the continuous annealing treatment is immersed in a plating bath (hot dip galvanizing bath) maintained at a predetermined temperature (400 to 500 ° C.), and a desired amount of plating layer (hot galvanized layer) is applied to the surface. Adhere. In addition, it cools after a plating process.

Cooling rate after plating treatment: 20 ° C / s or more If the coiling temperature and annealing temperature are in the proper ranges, cementite is almost not precipitated, so the amount of C not precipitated as carbide at the time of plating treatment is 0.022% That's it. Therefore, if the cooling rate after the plating treatment is less than 20 ° C./s, the amount of C that precipitates as cementite during cooling and does not precipitate as carbide becomes less than 0.022%. Therefore, the cooling rate after the plating treatment is set to 20 ° C./s or more on average. Note that the upper limit of the cooling rate is not particularly limited, and cooling may be performed within the range of equipment capacity as long as it is 20 ° C./s or more. The cooling after the plating treatment is preferably performed at a cooling rate described above to a temperature of 150 ° C. or lower.

  After performing the plating treatment, the plating layer may be further alloyed. As the alloying treatment of the plating layer, it is preferable to heat and cool to a usual heating temperature: 450 to 550 ° C. As the cooling, the average cooling rate is preferably 20 ° C./s or more.

Cooling rate after alloying treatment: 20 ° C / s or more If the coiling temperature and annealing temperature are in the appropriate ranges described above, and the cooling rate after plating treatment is within the above-mentioned ranges, cementite will remain at the time of alloying treatment. Almost no precipitation, C content not precipitated as carbide is 0.022% or more. When the cooling rate is less than 20 ° C / s, cementite precipitates during cooling and the amount of C not precipitated as carbides is less than 0.022%. For this reason, the cooling rate after alloying is preferably 20 ° C./s or more on average. Note that the upper limit of the cooling rate is not particularly limited, and cooling may be performed within the range of equipment capacity as long as it is 20 ° C./s or more. The cooling after the alloying treatment is preferably performed at a cooling rate described above to a temperature of 150 ° C. or lower.

Example 1
Molten steel having the composition shown in Table 1 was melted in a converter and continuously cast into a slab (steel material) having a thickness of 250 mm. These slabs were heated, roughly rolled, finish-rolled, then water-cooled, wound into a coil, and made into a hot-rolled steel sheet having a thickness of 2.6 mm. Some of the wound coils were water-cooled by being immersed in a water tank within 30 minutes after winding (immersion time: 2 hours or more). The remaining coils were allowed to cool without water cooling. Table 2 shows the heating temperature of the slab, the finish rolling end temperature, the winding temperature, and whether water cooling after winding is performed.

Samples are taken from the hot-rolled steel sheet obtained by the above method, and the structure observation, chemical analysis, tensile test, hole expansion test, impact test are performed by the following method, the ferrite phase area ratio, the average of the ferrite phase Grain size, carbide shape and size including Ti and V, solid solution Ti content, C content not precipitated as carbide, tensile strength, total elongation, hole expansion rate (stretch flangeability), ductile brittle transition temperature Asked.
(I) Microstructure observation A specimen is taken from the obtained hot-rolled steel sheet (center part in the thickness direction), the cross section in the rolling direction of the specimen is mechanically polished and corroded with nital, and then a scanning electron microscope (SEM) ) For 10 fields of view at 3000x magnification. And using the obtained structure | tissue photograph (SEM photograph), the kind of structure | tissue other than a ferrite phase and a ferrite phase, and the area ratio of the ferrite phase were calculated | required with the image-analysis apparatus. Further, the average crystal grain size of the ferrite phase was determined by a cutting method based on the provisions of JIS G 0551, using the structure photograph (SEM photograph) obtained as described above.

  Moreover, the thin film (test piece) produced from the obtained hot-rolled steel sheet was observed with a transmission electron microscope (magnification: 100,000 times or more), and the shape and size of the carbide containing Ti and V were determined.

  Observation with a transmission electron microscope is carried out from the [001] crystal direction with respect to the ferrite phase, which is the parent phase, that is, the [001] crystal direction of the ferrite phase, and the shape and size of the carbides detected for Ti and V by EDS analysis. Asked. The reason for limiting the observation direction is that carbides containing Ti and V are precipitated with a specific orientation relative to the ferrite phase.

  Carbides containing Ti and V are plate-like carbides, so that the plate surface of the plate is parallel to the [001] direction of the ferrite phase (that is, the normal of the plate surface of the plate is perpendicular to the [001] direction of the ferrite phase) To precipitate). Therefore, when the carbide containing Ti and V is observed from the [001] direction of the ferrite phase, it is observed in a rectangular shape.

For each carbide observed in a rectangular shape by the above-described method, the length of the portion with the largest dimension was defined as the length of the rectangle, and the length of the portion with the smallest dimension as the width of the rectangle. The length and width of the rectangle were measured for 100 or more carbides, and the arithmetic average value of each was taken as the average length of the rectangle and the average width of the rectangle.
(Ii) Chemical analysis Non-patent literature (Tetsufumi Shiroshiro, 3 others, “Development of solid solution determination method for alloy elements in high-strength steels”, Materials and Processes, Japan Iron and Steel Institute, 2010 , Vol.23, No.2, p.1348). That is, a test piece collected from the obtained hot-rolled steel sheet (central part in the plate thickness direction) was electrolyzed in a 10% acetylacetone-1% tetramethylammonium chloride-methanol solution, and the electrolyte immediately after electrolysis was collected, and EDTA After adding the aqueous solution, the iron concentration (Cfe), Ti concentration (Cti) and V concentration (Cv) were measured with an ICP mass spectrometer. In addition, elements other than iron contained in the test piece were quantified by spark discharge optical emission spectrometry, and the total value thereof was subtracted from 100% to calculate the Fe concentration (Fe%) of the test piece. Using these values, the amount of solid solution Ti and the amount of solid solution V were determined from the following equations.

Solid Ti content = (Cti / Cfe) x Fe%
Solid solution V = (Cv / Cfe) x Fe%
In addition, since the size of the Ti-V composite carbide contained in the hot-rolled steel sheet of the present invention is extremely fine, generally used precipitation amount analysis, that is, electrolyzing the steel sheet, and collecting the residue collected on the filter In the method of quantifying, the carbide leaks from the filter holes and cannot be quantified correctly. In order to know the exact amount of precipitation or solid solution in Ti and V steel, the above analytical method is necessary.

  The amount of C not precipitated as carbide was determined by the solid solution Ti amount, solid solution V amount and extraction residue method determined by the above method.

  Included in the test piece by calculating the value of (Ti-solid solution Ti amount-(S x 48/32)-(N x 48/14)) using the solid solution Ti amount obtained by the above method The amount of Ti precipitated as carbide in Ti was determined. In addition to Ti precipitated as fine Ti-V composite carbides (carbides containing Ti and V), Ti precipitated as Ti carbides that precipitate during rolling and become ferrite nucleation sites included. In the above formula, Ti, S, and N are the Ti content, S content, and N content (all by mass%) of the test piece. Moreover, the amount of V precipitated as a carbide | carbonized_material among V contained in a test piece was calculated | required by calculating the value of (V-solid solution V amount) using the amount of solid solution V calculated | required by said method. In the above formula, V is the V content (mass%) of the test piece.

  Furthermore, using a 10% acetylacetone-1% tetramethylammonium chloride-methanol solution as an electrolyte, the test piece was electrolyzed, and then the extraction residue collected on a 0.2 mm pore size filter was subjected to chemical analysis to become a carbide. The amount of Fe was determined.

  From these analysis results, the amount of C (mass%) not precipitated as carbide was calculated by the following formula. In the following formula, C is the C content (% by mass) of the test piece.

C amount not precipitated as carbide = C − ((Ti amount precipitated as carbide / 48) + (V amount precipitated as carbide / 51)) × 12− (Fe precipitation amount × 12 / (56 × 3))
As is clear from the above formula, the amount of C not precipitated as carbide is the amount of C in a form other than Ti-V composite carbide (Ti, V) C, Ti carbide and cementite Fe ₃ C. is there.
(Iii) Tensile test From the obtained hot-rolled steel sheet, a JIS No. 5 tensile test piece (JIS Z 2201) with the direction perpendicular to the rolling direction is taken, and a tensile test based on the provisions of JIS Z 2241 (2011) is performed. Tensile strength (TS) and total elongation (El) were measured.
(Iv) Hole expansion test A test piece (size: 130 mm x 130 mm) was taken from the obtained hot-rolled steel sheet, and a hole with an initial diameter d ₀ : 10 mm was punched into the test piece (clearance: test piece) 12.5% of the plate thickness). Using these test pieces, a hole expansion test was performed. That is, insert a conical punch with an apex angle of 60 ° from the punch side into the hole, expand the hole, and measure the diameter d (mm) of the hole when the crack penetrates the plate thickness of the steel plate (test piece) Then, the hole expansion ratio λ (%) was calculated by the following formula.
Hole expansion rate λ (%) = {(d−d ₀ ) / d ₀ } × 100

(V) Impact test
In accordance with JIS Z 2242 (2005), a Charpy impact test was performed on a 2.5 mm sub-size V-notch specimen, and the temperature at which the brittle fracture surface ratio was 50% was determined as the ductile brittle transition temperature vTrs (° C). .

  The obtained results are shown in Table 3.

All of the examples of the present invention have high tensile strength of 980 MPa or more, excellent workability of elongation of 15% or more and hole expansion ratio of 40% or more, and ductile brittle transition temperature vTrs is -20 ° C. or less. Has low temperature toughness. On the other hand, the comparative example which is outside the scope of the present invention does not ensure a predetermined high strength, does not ensure a sufficient hole expansion rate, or does not ensure sufficient toughness.
(Example 2)
Molten steel having the composition shown in Table 4 was melted in a converter and continuously cast into a slab (steel material) having a thickness of 250 mm. These slabs (steel materials) were heated to the heating temperatures shown in Table 5 and subjected to rough rolling and finish rolling to finish rolling at the finish rolling finish temperature shown in Table 5, and then the cooling rates shown in Table 5 ( Water cooling was performed at an average cooling rate), and the coil was wound around the coil at the winding temperature shown in Table 5 to obtain a hot-rolled steel sheet having a thickness of 2.6 mm. In addition, about some coils, coil water cooling which immersed in a water tank within 30 minutes after winding was implemented. The immersion time was 2 hours or longer. The other coils were allowed to cool without performing coil water cooling.

  About a part of obtained hot-rolled steel sheet, after pickling and removing the surface scale, the annealing process and the hot dip galvanizing process were performed on the conditions shown in Table 5 using the continuous annealing equipment. Note that the hot dip galvanizing treatment is performed by immersing the steel sheet after the continuous annealing treatment in a hot dip galvanizing bath (0.13% Al-Zn) at 460 ° C, and attaching a hot dip galvanized layer to the surface. ). The cooling after the hot dip galvanizing treatment was gas cooling or water cooling so that the average cooling rate shown in Table 5 was obtained. Further, some of the plated steel sheets were subjected to alloying treatment of a hot dip galvanized layer to obtain alloyed hot dip galvanized steel sheets. The alloying temperature was 520 ° C. The cooling after the alloying treatment was gas cooling or water cooling so that the average cooling rate shown in Table 5 was obtained.

  From the obtained hot-rolled steel sheet, a test piece was collected in the same manner as in Example 1 and subjected to structure observation, chemical analysis, tensile test, hole expansion test, impact test by the same method as in Example 1, Area ratio, average grain size of ferrite phase, shape and size of carbide containing Ti and V, solid solution Ti amount, C amount not precipitated as carbide, tensile strength, total elongation, hole expansion rate (stretch flange) ) And ductile brittle transition temperature.

  The amount of C not precipitated as carbide is determined by chemical analysis in the same manner as in Example 1. The amount of solid solution Mo, the amount of solid solution Nb, the amount of solid solution W is the same as the amount of solid solution Ti and the amount of solid solution V. After obtaining, using the obtained solid solution Mo amount, solid solution Nb amount, solid solution W amount, (Mo-solid solution Mo amount), (Nb-solid solution Nb amount), (W-solid solution W amount) By calculating the value of Mo, Nb, and W, the amount of Mo, Nb, and W precipitated as carbides was determined. In addition to the Ti and V contained in the test piece, the amount of Ti and V precipitated as carbides was added. The amount of C not precipitated as carbide was determined.

  The results obtained are shown in Table 6.

  All of the examples of the present invention have high tensile strength of 980 MPa or more, excellent workability of elongation of 15% or more and hole expansion ratio of 40% or more, and ductile brittle transition temperature vTrs is -20 ° C. or less. Has low temperature toughness. On the other hand, a comparative example that is out of the scope of the present invention does not ensure a predetermined high strength, does not ensure sufficient hole expansibility, or does not ensure sufficient toughness.

Claims (4)

The steel material is heated, subjected to hot rolling consisting of rough rolling and finish rolling, and after finishing rolling is cooled, wound, and made into a hot-rolled steel sheet, the steel material in mass%,
C: 0.07% or more and 0.12% or less, Si: 0.2% or less,
Mn: 0.1% to 1.2%, P: 0.025% or less,
S: 0.005% or less, Al: 0.1% or less,
N: 0.008% or less, Ti: 0.14% or more and 0.20% or less,
V: 0.05% or more and 0.15% or less, C, Ti, V, S and N are contained so as to satisfy the following formulas (1) and (2), and the balance is composed of Fe and inevitable impurities,
The heating temperature of the heating is 1150 ° C or higher and 1350 ° C or lower,
The finish rolling finish temperature of the finish rolling is 880 ° C. or higher,
The winding temperature of the winding is 550 ° C. or more and 700 ° C. or less, the hot-rolled coil after winding is water-cooled ,
The hot-rolled steel sheet has the above-mentioned composition, and solute Ti is 0.04% or less by mass%, and C not precipitated as carbide is 0.022% or more by mass%,
The area ratio of the ferrite phase with respect to the entire structure is 95% or more, the average crystal grain size of the ferrite phase is 5.0 μm or less, and is a carbide containing Ti and V, which was observed from the [001] direction of the ferrite phase. When the shape has a structure in which carbide having a rectangular shape with an average length of 8 nm or less and an average width of 1 nm or less is precipitated,
A method for producing a high-strength hot-rolled steel sheet excellent in workability and toughness, characterized by being a hot-rolled steel sheet having a tensile strength of 980 MPa or more .
Record
((Ti / 48) − (N / 14) − (S / 32) + (Mo / 96) + (Nb / 93) + (W / 184)) / (V / 51)> 1.0… (1)
(C / 12) / ((Ti / 48)-(N / 14)-(S / 32) + (V / 51) + (Mo / 96) + (Nb / 93) + (W / 184))> 1.0… (2)
Here, C, Ti, V, S, N, Nb, Mo, W: Content of each element (mass%) In place of the winding and hot-rolled steel sheet, the winding temperature of the winding is 400 ° C. or higher and 700 ° C. or lower when forming a plating layer on the surface of the hot-rolled steel sheet after winding and hot-rolling steel sheet. and then, cooled with water the hot rolled coil after the winding, subjected to continuous annealing and plating, after the plating treatment, the average cooling rate and characterized by cooling at 20 ° C. / s or higher to a temperature of 0.99 ° C. or less The manufacturing method of the high intensity | strength hot-rolled steel plate of Claim 1 . Following the plating process, subjected to alloying treatment of the plating layer, after alloying treatment, an average cooling rate: according to claim 2, wherein the cooling to 20 ° C. / s or higher at 0.99 ° C. below the temperature Manufacturing method of high strength hot rolled steel sheet. In addition to the above composition, one or more selected from Mo: 0.02% or more and 0.30% or less, Nb: 0.02% or more and 0.10% or less, and W: 0.02% or more and 0.10% or less. The method for producing a high-strength hot-rolled steel sheet according to any one of claims 1 to 3 , wherein the composition is contained.