JP5887903B2 - High strength hot-rolled steel sheet excellent in weldability and method for producing the same - Google Patents

Description

  The present invention relates to a high-strength hot-rolled steel sheet having high tensile strength (TS): 850 MPa or more and excellent weldability, and a method for producing the same, useful for the use of automobile members.

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 efficiency of automobiles, the use of thin high-strength hot-rolled steel sheets as raw materials for automobile parts is increasing because the weight reduction by thinning the members used is the most effective. On the other hand, many automobile parts made of steel plates are assembled by welding parts formed by pressing or burring. Therefore, steel plates for automobile parts are required to have excellent weldability in addition to high strength.

  However, an ultra-high strength steel sheet having a tensile strength of 850 MPa or more usually has a problem that weldability is remarkably deteriorated because it contains a large amount of alloy elements necessary for increasing the strength. Therefore, in order to apply a high-strength hot-rolled steel sheet to automobile parts and the like, it is essential to develop a high-strength hot-rolled steel sheet having weldability, and various techniques have been proposed so far.

  For example, in Patent Document 1, the steel plate composition is C: 0.01 to 0.1%, Si: 0.01 to 2%, Mn: 0.05 to 3%, P ≦ 0.1%, S ≦ 0.03%, Al: 0.005 to 1% in mass%. , N: 0.0005 to 0.005%, Ti: 0.05 to 0.5%, 0 <C- (12 / 48Ti-12 / 14N-12 / 32S) ≦ 0.05%, Mo + Cr ≧ 0.2%, and Cr ≦ 0.5 %, Mo ≤ 0.5%, C, S, N, Ti is contained, the balance is composed of Fe and inevitable impurities, and the steel sheet structure is composed of ferrite or ferrite and bainite. Has been. And according to the technique proposed in the cited document 1, it is said that Cr and Mo added to the steel sheet can suppress the softening of the heat-affected zone of the steel plate by clustering or precipitating with elements such as C during the steel plate welding. ing.

  Moreover, in patent document 2, the composition of the steel plate used as the board | substrate of a hot dip galvanization type steel plate is weight%, C: 0.01-0.1%, Si <= 0.3%, Mn: 0.2-2.0%, P <= 0.04%, S <= 0.02%, Al ≦ 0.1%, N ≦ 0.006%, Ti: 0.03 to 0.2%, including one or more of Mo ≦ 0.5% and W ≦ 1.0%, the balance being substantially Fe, weight %, A composition satisfying 4.8C + 4.2Si + 0.4Mn + 2Ti ≦ 2.5, the structure of the steel sheet is substantially a ferrite single phase, and in a range satisfying (Mo + W) / (Ti + Mo + W) ≧ 0.2 by atomic ratio, Ti and In addition, a technique has been proposed in which a precipitate containing less than 10 nm, preferably 5 nm or less, including one or more of Mo and W is dispersed. According to the technique proposed in Patent Document 2, when a ferrite single-phase structure having a low dislocation density is reinforced with fine precipitates, a high-tensile hot-rolled steel sheet excellent in strength-elongation balance and strength-stretch flangeability balance is obtained. It is supposed to be done. Further, according to the technique proposed in Patent Document 2, the fine precipitate is a carbide containing Ti and one or more of W and Mo, and Ti, Mo, and W in the carbide are further expressed in atomic ratio ( By setting Mo + W) / (Ti + Mo + W) ≧ 0.2, carbide that is difficult to coarsen with heating can be obtained, and the strength of the steel is maintained at a high temperature in the initial temperature range of spot welding. It can be reduced.

JP 2004-218077 A JP 2003-321736 A

  However, in the technique proposed in Patent Document 1, when a high strength steel sheet having a tensile strength of 850 MPa or more is to be obtained, the steel sheet structure is made of a ferrite and bainite, and a solid solution strengthening element is used in the steel sheet. It is necessary to add a large amount of certain Mn and Si. And when a steel sheet containing a low temperature transformation phase (bainite) is welded in this way, in a region where the cooling rate after heating is moderate in the weld heat affected zone that has undergone a thermal history of heating to cooling by welding heat, the low temperature transformation phase is It is tempered and softens. Further, when the content of C, Si, and Mn, which are solid solution strengthening elements, is large, the amount of softening due to the tempering is remarkably increased. Therefore, with the technique proposed in Patent Document 1, it is not possible to obtain a high-strength hot-rolled steel sheet that has a high tensile strength of 850 MPa or more and excellent weldability.

  On the other hand, according to the technique proposed in Patent Document 2, even when the steel sheet structure is substantially a ferrite single phase, a high strength with a tensile strength of 850 MPa or more by precipitating fine carbides and clusters. It can be a hot-rolled steel sheet. However, as shown in the examples, when the steel sheet strength of 850 MPa or more is to be secured by the technique proposed in Patent Document 2, the precipitated carbides and the like are refined to about 2 to 3 nm. It is necessary to let Since carbides and clusters with a size of less than 5 nm are very sensitive to the temperature and holding temperature at the weld heat affected zone, they become coarse due to the welding heat history and are stably welded. It is extremely difficult to suppress softening of the heat affected zone. That is, even if it is a high-strength hot-rolled steel sheet that uses fine carbide without using a low-temperature transformation phase, if the component design of the steel sheet is a component design that does not consider the thermal stability of carbide, it will undergo a welding heat history. Thus, there was a problem that the carbide in the heat affected zone was coarsened and the weld heat affected zone was softened.

  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 excellent weldability, that is, tensile strength capable of suppressing the softening phenomenon of the weld heat-affected zone: 850 MPa or more. And

In order to solve the above-mentioned problems, the present inventors have intensively studied various factors affecting the strength fluctuation and weldability of hot-rolled steel sheets, particularly the strength fluctuation of the weld heat-affected zone formed when welding hot-rolled steel sheets. did. As a result, the following findings (1) to (7) were obtained.
(1) In reducing the softening of the heat affected zone, it is effective to make the structure of the hot rolled steel sheet a ferrite single phase that is not affected by tempering.
(2) The content of Mn, a hardenability element, needs to be less than 0.5% in order to avoid the formation of a low-temperature transformation phase and obtain a hot-rolled steel sheet having a substantially single-phase ferrite structure.
(3) To make the tensile strength 850 MPa or more for a hot-rolled steel sheet that is substantially a ferrite single-phase structure and the content of Mn, a solid solution strengthening element, is reduced to less than 0.5%. It is necessary to adopt precipitation strengthening by Ti carbide, V carbide, and Ti and V composite carbide, and to deposit a lot of these fine carbides in the ferrite grains that are the matrix of the steel sheet.
(4) When adopting precipitation strengthening with carbides, in order to obtain a high-strength steel sheet with excellent weldability, it is effective to use V with a large amount of dissolution after making maximum use of Ti with a small amount of dissolution in steel. It is important to add them together.

(5) Although fine carbides precipitated in ferrite grains are extremely effective in improving the strength of the steel sheet, fine carbides present in the weld heat affected zone melt and coarsen through the welding heat history. It becomes a cause of softening zone generation in the weld heat affected zone.
(6) It is effective to contain a predetermined amount of B in the hot-rolled steel sheet in order to suppress the softening region in the weld heat affected zone, that is, to suppress dissolution and coarsening of carbides.
(7) For hot-rolled steel sheets whose Mn content is less than 0.5% and whose matrix is substantially a ferrite single-phase structure, carbide is precipitated in the ferrite grains to increase the tensile strength to 850 MPa or more and welding heat In order to suppress the dissolution and coarsening of the carbide in the affected part, it is effective that the average particle diameter of the carbide is 5 nm or more and 20 nm or less.

  The mechanism by which the dissolution and coarsening of carbides in the weld heat affected zone is suppressed by the inclusion of B is not always clear, but B is a segregated element that segregates at the carbide and matrix interface, resulting in coarse carbides. It is presumed that the interfacial energy, which is the driving force of, decreases, and the thermal stability of the carbide improves.

The present invention has been completed based on the above findings, and the gist thereof is as follows.
[1] By mass%
C: 0.055% or more and 0.11% or less, Si: 0.4% or less,
Mn: less than 0.5%, P: 0.02% or less,
S: 0.01% or less, Al: 0.1% or less,
N: 0.01% or less, B: 0.0005% or more and 0.005% or less,
A composition containing Ti: 0.05% or more and 0.25% or less, V: 0.01% or more and 0.5% or less, so that C, Ti, and V satisfy the following (1), the balance being Fe and inevitable impurities, and ferrite The area ratio of the phase is 95% or more, the carbide average grain size in the ferrite phase grains is 5 nm or more and 20 nm or less, the tensile strength is 850 MPa or more, welding High strength hot rolled steel sheet with excellent properties.
Record
0.8 ≦ ([C] / 12) / ([Ti] /48+0.7× [V] / 51) ≦ 1.5 ・ ・ ・ (1)
([C], [Ti], [V]: Content of each element (% by mass))

[2] In the above [1], in addition to the above composition, Nb: 0.01% or more and 0.2% or less by mass% is contained so as to satisfy the following formula (2) instead of the formula (1). A high-strength hot-rolled steel sheet with excellent weldability.
Record
0.8 ≦ ([C] / 12) / ([Ti] /48+0.7× [V] / 51 + [Nb] / 93) ≦ 1.5 (2)
([C], [Ti], [V], [Nb]: Content of each element (% by mass))

[3] In the above [1] or [2], in addition to the above composition, any one of Cr: 0.1% or less, Ni: 0.05% or less, W: 0.05% or less, Mo: 0.05% or less in mass% A high-strength hot-rolled steel sheet excellent in weldability, characterized by containing more than seeds.

[4] In any one of the above [1] to [3], in addition to the composition, Ca, Zr, Co, Hf, Sn, Sb, As, Mg, Zn, Cu, Pb, and REM in mass%. A high-strength hot-rolled steel sheet excellent in weldability, characterized by containing one or more of any one or more in total.

[5] A high-strength hot-rolled steel sheet excellent in weldability, characterized in that in any one of the above [1] to [4], the steel sheet surface has a plating layer.

[6] By mass%,
C: 0.055% or more and 0.11% or less, Si: 0.4% or less,
Mn: less than 0.5%, P: 0.02% or less,
S: 0.01% or less, Al: 0.1% or less,
N: 0.01% or less, B: 0.0005% or more and 0.005% or less,
Steel with a composition that contains Ti: 0.05% or more and 0.25% or less, V: 0.01% or more and 0.5% or less, so that C, Ti, and V satisfy the following (1), and the balance is Fe and inevitable impurities Is heated to a temperature of 1100 ° C or higher and 1350 ° C or lower, and is subjected to hot rolling at a finish rolling temperature of 820 ° C or higher, and forced cooling is started within 3 seconds after the hot rolling is completed. After cooling to a temperature range of 700 ° C to 800 ° C at the cooling rate, stop cooling, let cool for 2 to 25 s, cool again at a cooling rate of 20 ° C / s, and 550 to 700 ° C A method for producing a high-strength hot-rolled steel sheet excellent in weldability, characterized by winding in a coil shape at a winding temperature of.
Record
0.8 ≦ ([C] / 12) / ([Ti] /48+0.7× [V] / 51) ≦ 1.5 ・ ・ ・ (1)
([C], [Ti], [V]: Content of each element (% by mass))

[7] In the above [6], in addition to the above composition, Nb: 0.01% or more and 0.2% or less in mass% is contained so as to satisfy the following formula (2) instead of the formula (1). A method for producing a high-strength hot-rolled steel sheet having excellent weldability.
Record
0.8 ≦ ([C] / 12) / ([Ti] /48+0.7× [V] / 51 + [Nb] / 93) ≦ 1.5 (2)
([C], [Ti], [V], [Nb]: Content of each element (% by mass))

[8] In the above [6] or [7], in addition to the above composition, any one of Cr: 0.1% or less, Ni: 0.05% or less, W: 0.05% or less, Mo: 0.05% or less in mass% A method for producing a high-strength hot-rolled steel sheet excellent in weldability, characterized by containing at least a seed.

[9] In any one of the above [6] to [8], in addition to the composition, Ca, Zr, Co, Hf, Sn, Sb, As, Mg, Zn, Cu, Pb, and REM in mass%. A method for producing a high-strength hot-rolled steel sheet excellent in weldability, comprising one or more of any one or more in total.

[10] The method for producing a high-strength hot-rolled steel sheet excellent in weldability, characterized in that in any of the above [6] to [9], a plating treatment is further performed to form a plating layer on the surface.

  According to the present invention, a low-temperature transformation phase that is greatly affected by tempering is not utilized, and a steel sheet structure in which fine carbides are substantially precipitated in a ferrite single-phase matrix is formed. Reduce the softening of the weld heat-affected zone by adding a certain amount of B and improve the thermal stability of the carbide, and obtain a high-strength hot-rolled steel sheet with a tensile strength of 850 MPa or more with excellent weldability Can do. Therefore, according to the present invention, joining by welding which becomes difficult as the strength of the hot-rolled steel sheet becomes higher is possible, and the effects of simplifying the production of automobile members and reducing the weight of the automobile body are remarkable.

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 and the carbide average particle size in the ferrite phase crystal grains is 5 nm or more and 20 nm or less.

Area ratio of ferrite phase: 95% or more The low-temperature transformation phase causes a decrease in the strength of the weld heat-affected zone. In the vicinity of the weld zone, the weld heat affected zone is formed by passing through a thermal history of rapid heating → cooling during welding, but when a steel sheet containing a low temperature transformation phase is welded, the cooling rate after rapid heating of the weld heat affected zone is This is because the low-temperature transformation phase in the gentle region is tempered and softened. For the above reasons, in the present invention, it is preferable that the matrix of the hot-rolled steel sheet is a ferrite single phase for the purpose of suppressing the softening of the weld heat affected zone. Moreover, it is preferable to set it as a ferrite single phase also from a viewpoint of ensuring the workability of a steel plate. However, even if the matrix of the hot-rolled steel sheet is not completely a ferrite single phase, if the area ratio of the ferrite phase is substantially 95% or more, the above-mentioned softening of the heat affected zone is effective. Can be suppressed. Therefore, the area ratio of the ferrite phase is 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. When these structures are present in a large amount in the matrix, they are softened at the weld heat affected zone as described above. 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.

As described above, in the hot-rolled steel sheet of the present invention, in order to reduce the content of Mn, which is a quenching element and a solid solution strengthening element, for the purpose of suppressing the low-temperature transformation phase, The steel sheet strength cannot be improved by melt 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. In the present invention, examples of the carbide that finely precipitates carbide in the ferrite phase grains include Ti carbide, V carbide, and Ti and V composite carbide, or those containing Nb, Mo, and W in these carbides. 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: 5 nm or more and 20 nm or less The hot-rolled steel sheet of the present invention is strengthened by finely dispersing the above carbides in the ferrite phase crystal grains as a matrix. Here, when the carbides are coarsened, the number of carbides that hinder the movement of dislocations is reduced. Therefore, the finer the carbides, the better for increasing the strength of the steel sheet. In order to obtain a high-strength hot-rolled steel sheet having a tensile strength of 850 MPa or more, the average particle diameter of carbides needs to be 20 nm or less. Preferably it is 15 nm or less. On the other hand, extremely fine carbides easily dissolve when exposed to high temperatures during welding. In particular, a carbide having a particle size of less than 5 nm has a small effect of improving thermal stability by B, which will be described later, and therefore the lower limit of the average particle size of the carbide is set to 5 nm.

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% (mass%) unless there is particular notice.
C: 0.055% or more and 0.11% or less
C combines with Ti, V, or even Nb, 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 steel sheet having a tensile strength of 850 MPa or more, the C content needs to be at least 0.055%. On the other hand, if the C content exceeds 0.11%, the amount of C that does not precipitate as carbide increases, resulting in a low temperature transformation phase, temper softening in the weld heat affected zone, and deteriorates weldability. Therefore, the C content is 0.055% or more and 0.11% or less. Preferably they are 0.06% or more and 0.10% or less.

Si: 0.4% or less
Si is a ferrite-forming element, and promotes two-phase separation of austenite and ferrite at high temperatures when the steel sheet heated to the austenite region during cooling is cooled. When transformation is performed at such a high temperature (two-phase separation), the diffusion rate of the elements in the steel sheet is high, so the elements that produce a low-temperature transformation phase such as C and Mn are concentrated in the untransformed austenite. A low temperature transformation phase is easily generated in the weld heat affected zone. From such a viewpoint, the Si content is set to 0.4% or less for the purpose of suppressing the generation of a low-temperature transformation phase in the weld zone (welding heat affected zone). Preferably it is 0.2% 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. As described above, in the present invention, it is essential that the steel sheet contains a predetermined amount of B for the purpose of enhancing the thermal stability of the carbide precipitated in the ferrite grains, but B is an element having a very high hardenability improving effect. It is also an element that promotes the formation of a low temperature transformation phase. Therefore, in the present invention in which the steel sheet structure is substantially a ferrite single phase structure, it is necessary to reduce the content of Mn, which is a hardenability element, and the content is set to less than 0.5%. Preferably it is 0.4% or less. However, if the Mn content is extremely reduced, the ferrite transformation point after hot rolling rises, and there is a concern that carbides will become excessively coarse, so the Mn content may be 0.1% or more. Preferably, it is more preferably 0.2% or more.

P: 0.02% or less
P is a harmful element that significantly deteriorates the weldability because it segregates at the grain boundaries and promotes grain boundary embrittlement in the weld. Therefore, in the present invention, it is desirable to reduce it as much as possible, but since the content up to 0.02% is acceptable, the P content is set to 0.02% or less. Preferably it is 0.015% or less.

S: 0.01% or less
S is a harmful element that generates inclusions in the steel, causes microvoids in the weld, and causes a decrease in ductility. Therefore, in the present invention, it is preferable to reduce the S content as much as possible, but since it is acceptable up to 0.01%, the S content is set to 0.01% or less. Preferably it is 0.005% or less.

Al: 0.1% or less
Al is an element that acts as a deoxidizer. In order to obtain such an effect, it is desirable to contain 0.02% or more. However, if the Al content exceeds 0.1%, coarse inclusions such as Al ₂ O ₃ are formed, which is a cause of strength reduction in the weld. Therefore, the Al content is 0.1% or less. Preferably it is 0.08% or less.

N: 0.01% or less
N is a harmful element that forms a nitride serving as a fracture base in the weld heat affected zone and lowers the weld joint strength. Therefore, in the present invention, it is preferable to reduce the N content as much as possible, but up to 0.01% is acceptable, so 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 desired hot-rolled steel sheet strength (tensile strength: 850 MPa or more) is ensured by precipitating fine carbides in the ferrite crystal grains. It becomes coarse and tends to cause a decrease in strength. In order to suppress dissolution and coarsening of carbides in the weld heat affected zone, it is effective to contain B. When the B content is less than 0.0005%, the above effects are not sufficiently exhibited. On the other hand, since B is an element having an effect of improving hardenability, when the B content exceeds 0.005%, a low temperature transformation phase is likely to be generated. Therefore, the B content is 0.0005% or more and 0.005% or less. Preferably it is 0.0008% or more and 0.002% or less.

Ti: 0.05% or more and 0.25% or less
Ti is an element that combines with C to form fine carbides and contributes to increasing the strength of the steel sheet. In order to obtain a hot rolled steel sheet having a tensile strength of 850 MPa or more, the Ti content needs to be 0.05% or more. On the other hand, if the Ti content exceeds 0.25%, the carbides become coarse and the weld joint strength decreases. Therefore, the Ti content is 0.05% or more and 0.25% or less. Preferably they are 0.05% or more and 0.20% or less.
In addition, from the viewpoint that the steel plate strength is proportional to the precipitation amount of carbide, it is preferable to precipitate solute Ti by precipitating 80% or more of the Ti content as carbide.

V: 0.01% to 0.5%
V, like Ti, is an element that forms a carbide with C and contributes to increasing the strength of the steel sheet. Further, V is effective for increasing the strength of the steel sheet because it combines with Ti to form a fine composite carbide. In order to ensure the desired hot-rolled steel sheet strength (tensile strength: 850 MPa or more), the V content needs to be 0.01% or more. On the other hand, since V has a high solubility in steel, when the V content is excessive, it cannot be completely precipitated and remains in the steel as a solid solution. When the V content exceeds 0.5%, during the welding of the steel sheet, the solid solution V forms coarse carbides in the portion that has received the welding heat history, leading to a decrease in the strength of the weld heat affected zone. Therefore, the V content is 0.01% or more and 0.5% or less. Preferably they are 0.01% or more and 0.40% or less.

  In the present invention, when adopting precipitation strengthening by carbide, in order to obtain a high-strength steel sheet excellent in weldability, the maximum amount of Ti dissolved in the steel is used, and the amount of dissolution is large. It is preferable to add V effectively in combination. Ti and V remaining as a solid solution content cause deterioration of weldability. For this reason, it is preferable to promote precipitation of Ti and V carbides and to reduce solid solution Ti and V as much as possible. Since V easily precipitates with fine Ti carbides as nuclei, in the present invention, it is important to add Ti and V in combination for the purpose of promoting precipitation.

The hot-rolled steel sheet of the present invention contains C, Ti, and V in the above-described range and satisfying the expression (1).
0.8 ≦ ([C] / 12) / ([Ti] /48+0.7× [V] / 51) ≦ 1.5 ・ ・ ・ (1)
([C], [Ti], [V]: Content of each element (% by mass))
The above formula (1) is a requirement that must be satisfied in order that the structure (matrix) of the hot-rolled steel sheet is substantially a ferrite single-phase structure and the tensile strength of the hot-rolled steel sheet is 850 MPa or more. It is an important indicator.

  The hardness difference in the weld heat affected zone (softening phenomenon in the weld heat affected zone described above) has a strong correlation with the amount of solute C present in the steel. Formation of low-temperature transformation phase that causes softening of weld heat affected zone by precipitating almost all of C contained in steel as fine carbides that contribute to strengthening by combining with Ti and V carbide forming elements Can be avoided. Here, if the C content (atomic%) is excessive with respect to the total content (atomic%) of Ti and V, which are carbide forming elements, the amount of solid solution C that does not precipitate as carbide increases. And generation | occurrence | production of a low temperature transformation phase cannot be suppressed, a softening zone arises in a welding heat affected zone, and weldability falls. On the other hand, when the C content (atomic%) with respect to the total content of Ti and V (atomic%) decreases, the amount of C combined with Ti and V becomes insufficient, and the desired hot-rolled steel sheet strength (tensile strength: 850 MPa or more) Therefore, it becomes impossible to secure the amount of carbide precipitation necessary for the purpose.

Therefore, as a result of intensive studies by the present inventors, when ([C] / 12) / ([Ti] /48+0.7× [V] / 51) is less than 0.8, the carbide-forming element is sufficiently precipitated as carbide. However, it was found that the tensile strength of 850 MPa or more could not be obtained, whereas if it exceeded 1.5, the weld heat-affected zone softened due to excess of solid solution C that did not precipitate as fine carbides.
Therefore, in the present invention, ([C] / 12) / ([Ti] /48+0.7× [V] / 51) is set to 0.8 or more and 1.5 or less. Preferably it is 0.9 or more and 1.4 or less.

The above is the basic composition of the present invention. In addition to the above basic composition, Nb: 0.01% or more and 0.2% or less is further contained so as to satisfy the following formula (2) instead of the above formula (1). May be.
0.8 ≦ ([C] / 12) / ([Ti] /48+0.7× [V] / 51 + [Nb] / 93) ≦ 1.5 (2)
([C], [Ti], [V], [Nb]: Content of each element (% by mass))

Nb: 0.01% or more and 0.2% 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. It is an element that contributes to increasing the strength of hot-rolled steel sheets by forming fine carbides. In order to obtain such an effect, the Nb content is preferably set to 0.01% or more. On the other hand, when the Nb content exceeds 0.2%, when manufacturing a hot-rolled steel sheet, coarse Nb carbides are not completely dissolved during heating of the steel material before hot rolling, and the resulting hot-rolled steel sheet is finally obtained. Coarse Nb carbide may remain. And this coarse Nb carbide causes the strength of the steel sheet to decrease, and it becomes difficult to make the tensile strength of the steel sheet 850 MPa or more. Therefore, the Nb content is preferably 0.01% or more and 0.2% or less. Further, it is more preferably 0.01 or more and 0.15 or less.

0.8 ≦ ([C] / 12) / ([Ti] /48+0.7× [V] / 51 + [Nb] / 93) ≦ 1.5 (2)
([C], [Ti], [V], [Nb]: Content of each element (% by mass))
As described above, when the C content (atomic%) is excessive with respect to the total content of carbide-forming elements (atomic%), solid solution C that does not precipitate as carbide increases, resulting in the formation of a low-temperature transformation phase. As a result, a softened zone occurs in the weld heat affected zone, and weldability decreases. On the other hand, when the C content (atomic%) with respect to the total content of carbide forming elements (atomic%) decreases, the amount of C combined with the carbide forming elements becomes insufficient, and the desired hot-rolled steel sheet strength (tensile strength: 850 MPa or more) Therefore, it becomes impossible to secure the amount of carbide precipitation necessary for the purpose. Therefore, when Nb is contained, the formula (2) is satisfied instead of the formula (1).

  When ([C] / 12) / ([Ti] /48+0.7× [V] / 51 + [Nb] / 93) is less than 0.8, the carbide-forming elements do not precipitate sufficiently as carbides and the tensile strength is 850 MPa or more. On the other hand, if it exceeds 1.5, a softening phenomenon of the weld heat affected zone occurs due to an excess of solid solution C that does not precipitate as fine carbides. Therefore, in the present invention, ([C] / 12) / ([Ti] /48+0.7× [V] / 51 + [Nb] / 93) is set to 0.8 to 1.5. Preferably it is 0.9 or more and 1.4 or less.

Furthermore, in order to obtain a hot-rolled steel sheet with a tensile strength of 850 MPa or more, it is necessary to sufficiently precipitate carbide. Therefore, Ti, V or further Nb which is a carbide forming element is contained so as to satisfy the formula (3). It is desirable to make it.
[Ti] /48+0.7× [V] / 51 + [Nb] /93≧0.0031… (3)
([Ti], [V], [Nb]: Content of each element (mass%))
Here, when calculating the value of the left side of equation (3), when the steel sheet does not contain Nb, [Nb] is assumed to be zero.

  In the above formulas (1) to (3), a coefficient of V of 0.7 means that 70% of the added V is precipitated and 30% remains in solid solution. Since V has a smaller driving force for precipitation as compared with Ti, the proportion of V remaining unavoidably as a solid solution is about 30%.

In addition to the above basic composition, it may further contain one or more of Cr: 0.1% or less, Ni: 0.05% or less, W: 0.05% or less, Mo: 0.05% or less.
Cr: 0.1% or less
Cr constitutes a carbide and is an effective element for further increasing the weld joint strength and steel plate strength. However, if the Cr content exceeds 0.1%, the grain boundaries become sensitized and the weldability decreases. Therefore, the Cr content is preferably 0.1% or less. More preferably, it is 0.05% or less.

Ni: 0.05% or less
Ni has the effect of improving the toughness of the weld. However, if the Ni content exceeds 0.05%, a low-temperature transformation phase is generated, so that softening occurs in the heat affected zone and weldability deteriorates. Therefore, the Ni content is preferably 0.05% or less.

W: 0.05% or less
W is an element effective in improving weldability because it forms a carbide with good thermal stability, that is, a carbide that is difficult to melt and coarsen even when subjected to welding heat history. However, if the W content exceeds 0.05%, a low temperature transformation phase is generated, so softening occurs in the weld heat affected zone, and weldability is deteriorated. Therefore, the W content is preferably 0.05% or less.

Mo: 0.05% or less
Mo, like W, forms a carbide with good thermal stability, and is therefore an effective element for improving weldability. However, if the Mo content exceeds 0.05%, a low temperature transformation phase is generated, so softening occurs in the weld heat affected zone, and weldability is reduced. Therefore, the Mo content is preferably 0.05% or less.

  In addition to the basic composition described above, any one or more of Ca, Zr, Co, Hf, Sn, Sb, As, Mg, Zn, Cu, Pb, and REM may be contained in total of 1% or less. . In addition, from the viewpoint of the toughness of the weld zone, the content of these elements is preferably 0.5% or less in total. Components other than the above are Fe and inevitable impurities.

Tensile strength: 850MPa or higher For steel body weight reduction, higher steel plate strength is desirable. In addition, if the tensile strength of the steel sheet with a small softening amount in the weld heat affected zone is 850 MPa or more, the fatigue resistance and tensile strength of the welded portion are dramatically improved in addition to the effect of reducing the weight of the automobile body, which is more desirable. . In response to these demands, the present invention has been designed to make the most of particle dispersion strengthening that finely disperses and strengthens carbides as described above, and measures are taken to suppress the softening phenomenon in the weld heat affected zone. Yes.

  Specifically, in a hot-rolled steel sheet, sufficient carbide to ensure a tensile strength of 850 MPa or more by containing a carbide forming element Ti, V or further Nb at a predetermined ratio with respect to C. Precipitate fine. In addition, the matrix of the hot-rolled steel sheet has a ferrite single-phase structure substantially, the carbide average particle diameter is 5 nm or more and 20 nm or less, and further contains a predetermined amount of B having a carbide coarsening suppressing effect in the weld heat affected zone. The softening phenomenon is suppressed. As a result, a hot-rolled steel sheet having a tensile strength of 850 MPa or more that is excellent in weldability (that is, has a small amount of strength reduction in the weld heat affected zone) can be obtained.

  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 heating temperature in the plating treatment can be 700 ° C. or less, even if the hot-rolled steel sheet of the present invention is subjected to the plating treatment, the above-described effects of the present invention are not impaired. 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 cited as suitable examples. is there.

Next, the manufacturing method of the hot rolled steel sheet of the present invention will be described.
In the present invention, a steel material (steel slab) having the above-described composition is heated, hot-rolled, cooled after completion of finish rolling, and wound into a hot-rolled steel sheet. At this time, the heating temperature of the heating is set to 1100 ° C. or more and 1350 ° C. or less, the finishing rolling temperature of the hot rolling is set to 820 ° C. or more, and forced cooling is started within 3 seconds after completion of hot rolling (after finishing rolling). After cooling to a temperature range of 700 ° C or higher and 800 ° C or lower at a cooling rate of 20 ° C / s or higher, stop cooling, let cool for 2 seconds or longer and 25 seconds or lower, and cool again at a cooling rate of 20 ° C / s or higher And coiled at a winding temperature of 550 ° C. or higher and 700 ° C. or lower.

  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 to 1350 ° C, or if the carbide in the steel material is dissolved, the steel material is heated. You may carry out direct rolling without doing. The rough rolling conditions are not particularly limited.

Finishing rolling temperature: 820 ° C or more When the finishing rolling temperature is lower than 820 ° C, ferrite transformation starts during finish rolling, resulting in a structure in which ferrite grains are extended, and the ferrite grains have partially grown. Due to the grain structure, the steel sheet strength is significantly reduced. Therefore, the finish rolling temperature needs to be 820 ° C. or higher. Preferably it is 850 degreeC or more. The upper limit of the finish rolling temperature is not particularly required, but is naturally determined by the heating temperature before hot rolling, the rolling plate speed, the steel plate thickness, and the like. Therefore, the upper limit of the finish rolling temperature is substantially 980 ° C.

Time to start forced cooling after finishing rolling: Within 3s When steel plate after finishing rolling is kept at high temperature for a long time, carbides are strain-induced precipitation, and the carbides precipitate and grow unevenly and within the coil surface Cause variations in strength. Since such a strength variation causes a failure to obtain a stable weld joint strength, in the present invention, for the purpose of suppressing strain-induced precipitation, it is necessary to start forced cooling immediately after the end of hot rolling, and finish rolling. After completion, start cooling within at least 3s. Preferably it is within 2 s.

Average cooling rate: 20 ° C / s or more As mentioned above, the longer the time that the steel sheet is maintained at the high temperature after finish rolling, the easier the coarsening of carbides due to strain-induced precipitation proceeds and the strength variation as described above. Cause. 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 30 ° C./s or more. However, if the cooling rate after finishing rolling is excessively increased, the cooling stop temperature in the next step is less than 700 ° C, and there is a concern that carbides will be excessively refined. preferable.

Cooling stop temperature: 700 ° C. or higher and 800 ° C. or lower As described above, in the present invention, the average particle size of carbides precipitated in the matrix of the finally obtained hot rolled steel sheet needs to be 5 nm or more and 20 nm or less. Here, when the steel sheet after finish rolling is forcibly cooled to above 700 ° C. at an average cooling rate of 20 ° C./s or more, the average particle size of the precipitated carbides becomes extremely small, and the heat finally obtained It becomes difficult to set the average particle size of carbides precipitated in the matrix of the rolled steel sheet to 5 nm or more. On the other hand, if the above-mentioned forced cooling is stopped in a temperature range exceeding 800 ° C., the average particle size of the precipitated carbide becomes too large, and the average particle size of the carbide precipitated in the matrix of the hot-rolled steel sheet finally obtained Is less than 20 nm. Therefore, the forced cooling stop temperature is set to 700 ° C. or higher and 800 ° C. or lower. Preferably they are 700 degreeC or more and 750 degrees C or less.

Cooling time: 2 s or more and 25 s or less In the present invention, the forced cooling is stopped in the temperature range of 700 ° C. or more and 800 ° C. or less as described above, and then it is allowed to cool, thereby having a predetermined average particle size (5 nm or more and 20 nm or less) Promotes precipitation of carbides. Here, if the cooling time is less than 2 s, the carbide is not sufficiently precipitated, while if the cooling time exceeds 25 s, the carbide becomes excessively coarse. Therefore, the cooling time is 2 to 25 s. Preferably, it is 2 s or more and 15 s or less. Note that the temperature of the steel sheet during the cooling is preferably maintained at 700 ° C. or higher and 800 ° C. or lower.

Average cooling rate of recooling: 20 ° C./s or more In the present invention, it is necessary to maintain the state of the carbide obtained by the above-mentioned cooling. Here, when the average cooling rate after cooling is less than 20 ° C./s, the carbides obtained by cooling are coarsened, and the average particles of carbides precipitated in the matrix of the finally obtained hot-rolled steel sheet It becomes difficult to make the diameter 20 nm or less. Therefore, the average cooling rate after the cooling is set to 20 ° C./s or more. Preferably it is 30 ° C./s or more. However, if the average cooling rate after cooling is more than 150 ° C./s, it may be difficult to control to a desired coiling temperature described later.

Winding temperature: 550 ° C. or higher and 700 ° C. or lower If the winding temperature is lower than 550 ° C., a low temperature transformation phase is generated and softening occurs in the heat affected zone. On the other hand, when the coiling temperature exceeds 700 ° C., the carbides are coarsened, and the hot-rolled steel sheet strength and the welded joint strength are reduced. Therefore, the coiling temperature is set to 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, in the present invention, the matrix is substantially a ferrite single phase, and has a structure in which carbides having an average particle diameter of 5 nm or more and 20 nm or less are sufficiently precipitated in the ferrite crystal grains, and has a tensile strength. A high-strength hot-rolled steel sheet having a weld strength of 850 MPa or more and excellent weldability can be obtained. The hot-rolled steel sheet after being rolled up after hot rolling does not change its properties even if the scale is attached to the surface or the scale is removed by pickling. 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. Note that the hot-rolled steel sheet of the present invention does not change the state of carbides in a temperature range of about 700 ° C. or lower, and the material fluctuation is small even when a heat treatment of about 700 ° C. or lower is performed. Therefore, when forming a plating layer, the continuous plating line which made the annealing temperature 700 degrees C or less can be passed. 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 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.

A 250 mm thick slab (steel material) having the composition shown in Table 1 was subjected to hot rolling under the hot rolling conditions shown in Table 2 to obtain a hot rolled steel sheet having a thickness of 2.0 mm. In addition, the cooling rate described in Table 2 is an average cooling rate from the finish rolling temperature to the cooling start temperature. Moreover, the recooling rate described in Table 2 is an average cooling rate from the end-of-cooling temperature to the winding temperature.
Furthermore, a part of the obtained hot-rolled steel sheet (steel plates No. 8, 10, 13) was passed through a continuous hot dip galvanizing line with an annealing temperature of 700 ° C, and then a 460 ° C plating bath (plating) Composition: Zn-0.13 mass% Al), and then alloyed at 520 ° C. to obtain a plating material (GA material). Plating was formed on both the front and back surfaces of the steel sheet, and the amount of plating was 45 g / m ² per side.

  Samples were taken from the hot-rolled steel sheet (hot-rolled steel sheet and GA material) obtained as described above, and the structure observation and tensile test were performed. The area ratio of the ferrite phase, the type and area ratio of the structure other than the ferrite phase, the ferrite phase The average particle diameter, yield strength, tensile strength, and elongation of the carbide precipitated in the crystal grains were determined. Also, weld samples are prepared by welding the hot-rolled steel sheets obtained as described above, specimens are taken from the hot-rolled steel sheets after welding, hardness tests are performed, and the hardness of the base material and the weld heat-affected zone Asked. Furthermore, a test piece was collected from the hot-rolled steel sheet after welding and subjected to a tensile test to obtain a tensile strength and observe the fracture position of the test piece. The test method is 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 in which corrosion marks are not observed in the grains and the grain boundaries are observed with a smooth curve. 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 relative 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 particle size of the carbides in the ferrite phase grains was measured using a thin film method from the center of the thickness of the obtained hot-rolled steel sheet and observed with a transmission electron microscope (magnification: 120,000 times). It calculated | required by the average of the precipitate particle diameter obtained about 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 ₀ ), and total elongation (El) were determined. The crosshead speed in the tensile test was 10 mm / min.

(Iii) Hardness test of weld sample Arc welding was performed using the obtained hot-rolled steel sheet to prepare a weld sample. The arc welding was carbon dioxide arc welding using a 1.2 mm diameter MGS-80 wire. The welding conditions for the arc welding were butt welding with a welding speed of 80 cm / min, a welding current of 220 A, a welding voltage of 25 V, and a plate gap of 1 mm. After welding, a cross section of the bead part was cut out, and a Vickers hardness test with a test load of 1 kgf was performed at 0.5 mm intervals on the central part of the plate thickness of the cross section. By the Vickers hardness test, the hardness HV ₀ of the base metal and the minimum hardness HV ₁ at the weld heat affected zone were determined, and the difference between the two (HV ₀ −HV ₁ ) was determined.

(Iv) Tensile test of weld sample A JIS No. 13 B tensile test was prepared from the weld sample of (iii) above, and a tensile test was performed under the same conditions as in (ii) above to evaluate weldability. Preparation of the JIS No. 13 B tensile test was performed under the condition that the weld bead portion of the welded sample crossed in the test piece width direction at the center of the distance between the test piece scores.
Evaluation of weldability is based on the difference between the tensile strength TS ₁ of the weld sample and the tensile strength TS ₀ measured in (ii) above (the tensile strength of the specimen before welding) TS ₀ −TS ₁ is 20 MPa or less. In addition, a material having a fracture position of the tensile test piece after the tensile test as the base material was defined as good weldability (◯). On the other hand, the difference TS ₀ −TS ₁ between the tensile strength TS ₁ of the welded sample and the tensile strength TS ₀ (tensile strength of the test piece before welding) measured in (ii) above exceeds 20 MPa, or tensile The test piece where the fracture position of the tensile test piece after the test was the weld heat affected zone was defined as poor weldability (x).
The obtained results are shown in Table 3.

As shown in Table 3, according to the invention example satisfying the conditions of the present invention, a hot-rolled steel sheet having a high strength of 850 MPa or more in tensile strength and excellent weldability is obtained. On the other hand, the comparative example which does not satisfy the conditions of this invention has not obtained desired steel plate strength, or was inferior to weldability. Further, in the comparative example in which the weldability was poor, the difference (HV ₀ −HV ₁ ) between the hardness HV ₀ of the base metal and the minimum hardness HV ₁ in the welding heat-affected zone was large, and the welding was accordingly performed. The amount of decrease in tensile strength due to (TS ₀ -TS ₁ ) also tends to increase. The breaking position of the tensile test at the level where the softening region occurred was the heat-affected zone. This is because tempering of the low temperature transformation phase or coarsening of the carbide occurred in the softening region.

Claims (10)

% By mass
C: 0.055% or more and 0.11% or less, Si: 0.4% or less,
Mn: less than 0.5%, P: 0.02% or less,
S: 0.01% or less, Al: 0.1% or less,
N: 0.01% or less, B: 0.0005% or more and 0.005% or less,
A composition containing Ti: 0.05% or more and 0.25% or less, V: 0.01% or more and 0.5% or less, so that C, Ti, and V satisfy the following (1), the balance being Fe and inevitable impurities, and ferrite The area ratio of the phase is 95% or more, the carbide average grain size in the ferrite phase grains is 5 nm or more and 20 nm or less, the tensile strength is 850 MPa or more, welding High strength hot rolled steel sheet with excellent properties.
Record
0.8 ≦ ([C] / 12) / ([Ti] /48+0.7× [V] / 51) ≦ 1.5 ・ ・ ・ (1)
([C], [Ti], [V]: Content of each element (% by mass)) In addition to the above composition, Nb: 0.01% or more and 0.2% or less by mass% is contained so as to satisfy the following formula (2) instead of the above formula (1): High strength hot-rolled steel sheet with excellent weldability as described.
Record
0.8 ≦ ([C] / 12) / ([Ti] /48+0.7× [V] / 51 + [Nb] / 93) ≦ 1.5 (2)
([C], [Ti], [V], [Nb]: Content of each element (% by mass))   In addition to the above composition, the composition further contains at least one of Cr: 0.1% or less, Ni: 0.05% or less, W: 0.05% or less, Mo: 0.05% or less in mass%. A high-strength hot-rolled steel sheet having excellent weldability according to 1 or 2.   In addition to the above composition, the composition further contains one or more of Ca, Zr, Co, Hf, Sn, Sb, As, Mg, Zn, Cu, Pb, and REM in mass% in total of 1% or less. A high-strength hot-rolled steel sheet excellent in weldability according to any one of claims 1 to 3.   The high-strength hot-rolled steel sheet having excellent weldability according to any one of claims 1 to 4, wherein the steel sheet surface has a plating layer. % By mass
C: 0.055% or more and 0.11% or less, Si: 0.4% or less,
Mn: less than 0.5%, P: 0.02% or less,
S: 0.01% or less, Al: 0.1% or less,
N: 0.01% or less, B: 0.0005% or more and 0.005% or less,
Steel with a composition that contains Ti: 0.05% or more and 0.25% or less, V: 0.01% or more and 0.5% or less, so that C, Ti, and V satisfy the following (1), and the balance is Fe and inevitable impurities Is heated to a temperature of 1100 ° C or higher and 1350 ° C or lower, and is subjected to hot rolling at a finish rolling temperature of 820 ° C or higher, and forced cooling is started within 3 seconds after the hot rolling is completed. After cooling to a temperature range of 700 ° C to 800 ° C at the cooling rate, stop cooling, let cool for 2 to 25 s, cool again at a cooling rate of 20 ° C / s, and 550 to 700 ° C Ri taken up by the coiling temperature in a coil shape, and the area ratio of the ferrite phase is 95% or more, have a tissue carbide average grain size of the crystal grains of the ferrite phase is 5nm or more 20nm or less, the tensile strength Saga characterized hot-rolled steel sheet and to Rukoto not less than 850 MPa, the method of producing a high strength hot-rolled steel sheet with excellent weldability.
Record
0.8 ≦ ([C] / 12) / ([Ti] /48+0.7× [V] / 51) ≦ 1.5 ・ ・ ・ (1)
([C], [Ti], [V]: Content of each element (% by mass)) In addition to the above composition, Nb: 0.01% or more and 0.2% or less by mass% is contained so as to satisfy the following formula (2) instead of the above formula (1): The manufacturing method of the high strength hot-rolled steel plate excellent in the weldability of description.
Record
0.8 ≦ ([C] / 12) / ([Ti] /48+0.7× [V] / 51 + [Nb] / 93) ≦ 1.5 (2)
([C], [Ti], [V], [Nb]: Content of each element (% by mass))   In addition to the above composition, the composition further contains at least one of Cr: 0.1% or less, Ni: 0.05% or less, W: 0.05% or less, Mo: 0.05% or less in mass%. A method for producing a high-strength hot-rolled steel sheet having excellent weldability according to 6 or 7.   In addition to the above composition, the composition further contains one or more of Ca, Zr, Co, Hf, Sn, Sb, As, Mg, Zn, Cu, Pb, and REM in mass% in total of 1% or less. A method for producing a high-strength hot-rolled steel sheet having excellent weldability according to any one of claims 6 to 8. Furthermore, and forming a plating layer on the surface subjected to plating, the method of producing a high strength hot-rolled steel sheet having excellent weldability according to claim 6 which have Shi 9.