KR101997382B1 - Steel sheet and manufacturing method - Google Patents

Abstract

Wherein the metal structure of the steel sheet has a ratio of the number of carbides on the ferrite grain bound to the number of carbides in the ferrite grains of more than 1 and a ferrite grain size of not less than 5 탆 and not more than 50 탆 And the Vickers hardness of the steel sheet is 100 HV or more and 170 HV or less.

Description

Steel sheet and manufacturing method

The present invention relates to a steel sheet and a manufacturing method thereof.

Automotive parts such as gears and clutches are manufactured through processing steps such as punching, forging, and pressing. In order to improve the product quality, stabilize it, and reduce the manufacturing cost in the processing step, it is required to improve the workability of the carbon steel sheet as the material. Further, since these components are used at high strength after? Tempering, excellent quenching is required.

Many proposals have been made so far in order to secure the workability and uniformity of the carbon steel sheet.

Patent Document 1 discloses a steel sheet which comprises 0.20 to 0.45% of C, 0.40 to 1.50% of Mn, 0.03% or less of P, 0.02% or less of S, 0.010% or more of P + S, 0.01 to 0.80% Ti: 0.005 to 0.050%, B: 0.0003 to 0.0050%, the balance being Fe and unavoidable impurities, and further containing 0.05% or less of Sn and 0.05% or less of Te, Of 0.005% or more and comprising a mixed structure of ferrite and pearlite or a mixed structure of ferrite and cementite, which has excellent workability, quenching and toughness after heat treatment.

Patent Document 2 discloses a steel sheet comprising 0.2 to 0.7% of C, 2% or less of Si, 2% or less of Mn, 0.03% or less of P, 0.03% or less of S, (Ar3 transformation point -20 占 폚) or higher to a steel containing 0.01% or less of iron and inevitable impurities as the remainder and then subjected to hot rolling at a cooling rate exceeding 120 占 폚 / Cooling at 620 占 폚 or lower and then winding at a coiling temperature of 600 占 폚 or lower to control the structure to have a bainite phase with a volume ratio exceeding 20%. After pickling, the annealing temperature is 640 占 폚 or higher and the Ac1 transformation point or lower Annealing is carried out to obtain a spheroidized structure, and a method for producing a high-kink high-carbon hot-rolled steel sheet is disclosed.

Japanese Patent No. 4319940 Japanese Patent No. 3879459

However, in the high carbon steel sheet described in Patent Document 1, pearlite having high hardness is also used in the material structure, and the workability is not necessarily excellent. Patent Document 2 does not describe a specific tissue form having excellent processability.

SUMMARY OF THE INVENTION It is an object of the present invention to provide a steel sheet which is improved in moldability and abrasion resistance and is particularly suitable for obtaining parts such as gears and clutches by means of a thick plate forming method and a method of manufacturing the same.

In order to solve the above problems and obtain a steel sheet suitable for a material of a drivetrain part or the like, it is necessary to increase the grain size of the ferrite and to make the carbide (mainly cementite) It is understood that the pearlite structure is reduced. This is for the following reasons.

The ferrite phase has low hardness and high ductility. Therefore, it is possible to increase the material formability by increasing the grain size of a structure mainly composed of ferrite.

Carbide is an indispensable structure for driveline parts because it can impart excellent abrasion resistance and electric fatigue characteristics while appropriately dispersing the carbides in the metal structure while maintaining the material formability. Further, the carbide in the steel sheet is a solid particle which interferes with slippage, and by allowing the carbide to exist in the ferrite grain boundaries, propagation of slip exceeding grain boundaries can be prevented, formation of shear zones can be suppressed, Thereby improving the formability of the steel sheet.

However, the cementite is a hard and brittle structure, and when present in the form of pearlite, which is a layered structure with ferrite, the cementite is hard and brittle and must be present in spherical form. Considering the occurrence of cracking at the time of cold forging or forging, the grain size thereof must be within an appropriate range.

However, a manufacturing method for realizing the above-described structure has not been disclosed so far. Therefore, the inventors of the present invention have made intensive studies on a manufacturing method for realizing the above-described structure.

As a result, the metal structure of the steel sheet after being rolled after hot rolling is rolled at a relatively low temperature (400 to 550 DEG C) in order to obtain a bainite structure in which cementite is dispersed in fine pearlite or fine ferrite having small lamellar spacing. By winding at a relatively low temperature, cementite dispersed in ferrite is also easily spheroidized. Subsequently, as the first-stage annealing, the cementite is partly spheroidized by annealing at a temperature just below the Ac1 point. Subsequently, as a second stage annealing, a part of the ferrite particles is left while annealing in a temperature between the Ac1 point and the Ac3 point (a so-called two-phase region of ferrite and austenite), and a part of the ferrite is austenite transformed. The ferrite grains are gently cooled and then the ferrite grains are grown while ferrite grains are grown to ferrite the austenite with the ferrite grains as nuclei to obtain cementite in the grain boundaries while obtaining a large ferrite phase.

That is, it is difficult to realize a method of manufacturing a steel sheet that simultaneously satisfies both formability and moldability, even if the hot rolling condition or the annealing condition is studied singly, and it can be realized by achieving optimization by a so-called continuous process such as hot rolling and annealing I got it.

The present invention has been made based on the above knowledge, and its gist of the invention is as follows.

(1) A ferritic stainless steel comprising, by mass%, 0.10 to 0.40% of C, 0.01 to 0.30% of Si, 1.00 to 2.00% of Mn, 0.020% or less of P, 0.010% or less of S, 0.001 to 0.10% Mo: not more than 0.10%, Nb: not more than 0.10%, V: not more than 0.10%, Cu: not more than 0.10%, W: not more than 0.10%, Ta: not more than 0.10% 0.10% or less of Ni, not more than 0.050% of Sn, not more than 0.050% of Sb, not more than 0.050% of As, not more than 0.050% of Mg, not more than 0.050% of Ca, not more than 0.050% : 0.050% or less, Ce: 0.050% or less, and the balance being Fe and inevitable impurities, wherein the metal structure of the steel sheet has a ratio of the number of carbides of the ferrite grains to the number of carbides in the ferrite grains , A ferrite grain size of 5 탆 or more and 50 탆 or less and an area ratio of pearlite of 6% or less, and the Vickers hardness of the steel sheet is 100 HV or more and 170 HV or less.

(2) The steel sheet according to the above (1), which contains at least one of Ti and Ti in an amount of not more than 0.10% and not more than 0.010%.

(3) A method of manufacturing the steel sheet according to the above (1) or (2), wherein the steel strip having the component composition of (1) or (2) is hot rolled in a temperature range of 750 ° C. to 850 ° C. Rolling the steel sheet to form a hot-rolled steel sheet, winding the hot-rolled steel sheet at a temperature of 400 ° C to 550 ° C, pickling the wound hot-rolled steel sheet, The annealing is performed in a second stage in which the hot-rolled steel sheet is maintained in the temperature range of 725 DEG C to 790 DEG C for not less than 3 hours and not more than 50 hours. After the annealing, Wherein the hot-rolled steel sheet is cooled to 650 占 폚 at a cooling rate of not less than 1 占 폚 / hour and not more than 30 占 폚 / hour.

According to the present invention, it is possible to provide a steel sheet excellent in moldability and abrasion resistance, particularly suitable for obtaining parts such as gears and clutches, by means of heavy plate casting, and a method of manufacturing the same.

Hereinafter, the present invention will be described in detail. First, the reasons for limiting the composition of the inventive steel sheet will be described. Below. % &Quot; of the component means "% by mass ".

[C: 0.10 to 0.40%]

C is an element effective for forming a carbide in a steel and strengthening steel and making ferrite particles finer. It is necessary to suppress the coarsening of the ferrite grain size in order to suppress crepe formation in the cold working and secure the surface aesthetics of the cold worked part. When the content is less than 0.10%, the volume percentage of the carbide is insufficient, The coarsening of the carbides in the steel can not be suppressed, so C is set to 0.10% or more. Preferably at least 0.12.

On the other hand, if it exceeds 0.40%, the volume ratio of the carbide increases, and when the load is instantaneously loaded, a large amount of cracks are generated as a starting point of the fracture, and the impact resistance characteristic is lowered. Preferably, it is 0.38% or less.

[Si: 0.01 to 0.30%]

Si is an element which acts as a deoxidizing agent and also influences the shape of the carbide. In order to obtain a deoxidizing effect, the Si content should be 0.01% or more. It is preferably at least 0.05%.

On the other hand, if it exceeds 0.30%, the ductility of ferrite lowers, cracking easily occurs in cold working, and cold workability is lowered, so Si is set to 0.30% or less. And preferably 0.28% or less.

[Mn: 1.00 to 2.00%]

Mn is an element contributing to improvement of strength by increasing quenching. If it is less than 1.00%, it becomes difficult to secure the strength after the casting and the residual carbide after casting, so that the Mn content is 1.00% or more. It is preferably 1.09% or more.

On the other hand, if the Mn content exceeds 2.00%, Mn segregation becomes an extreme band shape and the workability is remarkably lowered, so Mn is 2.00% or less. And preferably 1.91% or less.

[Al: 0.001 to 0.10%]

Al acts as a deoxidizing agent for the steel and stabilizes the ferrite. If it is less than 0.001%, the effect of addition can not be obtained sufficiently, so that Al is 0.001% or more. It is preferably 0.004% or more.

On the other hand, if it exceeds 0.10%, a large amount of inclusions are produced and the cold workability is deteriorated. Therefore, the content of Al is 0.10% or less. And preferably 0.08% or less.

The following elements are impurities, and it is necessary to control them to a certain amount or less.

[P: 0.0001 to 0.020%]

P is an element segregating at the ferrite grain boundaries and inhibiting the formation of grain boundary carbides. However, when the P is reduced to less than 0.0001% in the refining process, the polishing cost greatly increases, so that P is set to 0.0001% or more. It is preferably 0.0013% or more.

On the other hand, if it exceeds 0.020%, the ratio of the number of intergranular carbides decreases and the cold workability decreases, so that P is 0.020% or less. Preferably 0.018% or less.

[S: 0.0001 to 0.010%]

S is an impurity element forming a nonmetallic inclusion such as MnS. Since the nonmetallic inclusion is a starting point of occurrence of cracking in cold working, S is preferably as small as possible, but when the S is reduced to less than 0.0001%, the refining cost remarkably increases, so S is 0.0001% or more. It is preferably at least 0.0012%.

On the other hand, when it exceeds 0.010%, the cold workability decreases, so S is set to 0.010% or less. It is preferably 0.007% or less.

[N: 0.0001 to 0.010%]

N is an element which induces embrittlement of ferrite by the incorporation of a large amount, and the smaller N is preferable. The content of N may be 0, but if it is reduced to less than 0.0001%, the refining cost greatly increases, so the actual lower limit is 0.0001 to 0.0006%. On the other hand, if it exceeds 0.010%, ferrite becomes brittle and cold workability decreases, so N is 0.010% or less. It is preferably 0.007% or less.

[O: 0.0001 to 0.020%]

O is an element which forms a coarse oxide in the steel due to the inclusion of a large amount, and is preferably small. The content of O may be zero, but if it is reduced to less than 0.0001%, the refining cost greatly increases, so that the practical lower limit is 0.0001 to 0.0011%. On the other hand, if it exceeds 0.020%, a coarse oxide is generated in the steel and becomes a starting point of cracking in cold working, so that O is 0.020% or less. It is preferably 0.017% or less.

[Sn: 0.001 to 0.050%]

Sn is an element incorporated from a steel raw material (scrap). It is segregated in the grain boundary, and the ratio of the number of intergranular carbides is lowered. The content of Sn may be 0, but if it is reduced to less than 0.001%, the refining cost drastically increases, so that the practical lower limit is 0.001 to 0.002% or more. On the other hand, if it exceeds 0.050%, the ferrite becomes brittle and the cold workability decreases, so that the content of Sn is 0.050% or less. Preferably 0.040% or less.

[Sb: 0.001 to 0.050%]

Sb, like Sn, is an element incorporated from a steel raw material (scrap). It is segregated in the grain boundary, and the ratio of the number of intergranular carbides is lowered. The content of Sb may be zero, but if it is reduced to less than 0.001%, the refining cost drastically increases, so that the practical lower limit is 0.001 to 0.002% or more. On the other hand, if it exceeds 0.050%, the ferrite becomes brittle and the cold workability deteriorates, so that Sb is 0.050% or less. Preferably 0.040% or less.

[As: 0.001 to 0.050%]

As, like Sn and Sb, is an element incorporated from a steel raw material (scrap). It is segregated in the grain boundary, and the ratio of the number of intergranular carbides is lowered. The content of As may be zero, but if it is reduced to less than 0.001%, the refining cost drastically increases, so that the practical lower limit is 0.001 to 0.002% or more. On the other hand, if it exceeds 0.050%, the ratio of the number of intergranular carbides decreases, and the cold workability decreases, so that the content of As is 0.050% or less. Preferably 0.040% or less.

The steel sheet of the present invention contains the above element as a basic component, but may also contain the following elements for the purpose of improving the cold-rolling of the steel sheet. The following elements are not essential for obtaining the effect of the present invention, and thus the content may be zero.

[Cr: 0.50% or less]

Cr is an element that contributes to the improvement of strength by increasing quenching and is an element which is concentrated in the carbide and forms a stable carbide even in the austenite phase. In order to obtain the effect of addition, Cr is preferably 0.001% or more. More preferably, it is 0.007% or more. On the other hand, if it is more than 0.50%, the carbide is stabilized and the dissolution of the carbide is delayed at the time of casting, so that the required quenching strength may not be attained. Therefore, the content of Cr should be 0.50% or less. It is preferably 0.45% or less.

[Mo: 0.10% or less]

Mo, like Mn, is an element effective for controlling the shape of carbides. In order to obtain the effect of addition, Mo is preferably 0.001% or more. More preferably, it is 0.010% or more. On the other hand, if it exceeds 0.10%, the in-plane anisotropy of the r value deteriorates and the cold workability decreases, so Mo is set to 0.10% or less. And preferably 0.08% or less.

[Nb: 0.10% or less]

Nb is an element effective for controlling the shape of a carbide and is an element contributing to improvement of toughness by making the structure finer. In order to obtain the effect of addition, Nb is preferably 0.001% or more. More preferably, it is 0.002% or more. On the other hand, if it exceeds 0.10%, a large number of fine Nb carbides are precipitated, the strength excessively increases, the ratio of the number of intergranular carbides decreases, and the cold workability decreases, so that Nb is 0.10% or less. And preferably 0.08% or less.

[V: 0.10% or less]

V, like Nb, is an element effective for controlling the shape of a carbide, and is an element contributing to improvement of toughness by making the structure finer. In order to obtain the effect of addition, V is preferably 0.001% or more. More preferably, it is 0.004% or more. On the other hand, if it exceeds 0.10%, a large number of fine V carbides are precipitated, the strength excessively increases, the ratio of the number of intergranular carbides decreases, and the cold workability decreases, so V is 0.10% or less. And preferably 0.08% or less.

[Cu: 0.10% or less]

Cu is an element that segregates in the grain boundaries of ferrite and forms fine precipitates and contributes to improvement of strength. In order to obtain the effect of addition, Cu is preferably 0.001% or more. More preferably, it is 0.005% or more. On the other hand, if it exceeds 0.10%, the brittle brittleness is generated and the productivity in hot rolling is lowered, so that the content of Cu is 0.10% or less. And preferably 0.08% or less.

[W: 0.10% or less]

Like W, Nb and V, W is an element effective for controlling the shape of carbide. In order to obtain the effect of addition, W is preferably 0.001% or more. More preferably, it is 0.003% or more. On the other hand, if it exceeds 0.10%, a large number of fine W carbides are precipitated, the strength excessively increases, the ratio of the number of intergranular carbides decreases, and the cold workability decreases, so W is 0.10% or less. And preferably 0.08% or less.

[Ta: 0.10% or less]

Like Ta, Nb, V, and W, Ta is an element effective for controlling the shape of carbide. In order to obtain the effect of addition, the content of Ta is preferably 0.001% or more. More preferably, it is 0.005% or more. On the other hand, if it exceeds 0.10%, a large number of fine W carbides are precipitated, the strength excessively increases, the ratio of the number of intergranular carbides decreases, and the cold workability decreases, so that Ta is 0.10% or less. And preferably 0.08% or less.

[Ni: 0.10% or less]

Ni is an element effective for improving the toughness of a component. In order to obtain the effect of addition, Ni is preferably 0.001% or more. More preferably, it is 0.004% or more. On the other hand, if it exceeds 0.10%, the ratio of the number of intergranular carbides decreases and the cold workability decreases, so that the Ni content is 0.10% or less. And preferably 0.08% or less.

[Mg: 0.050% or less]

Mg is an element capable of controlling the form of sulfide by the addition of a trace amount. In order to obtain the effect of addition, Mg is preferably 0.0001% or more. More preferably, it is 0.0008% or more. On the other hand, if it exceeds 0.050%, the ferrite becomes brittle and the cold workability decreases, so that the Mg content is 0.050% or less. Preferably 0.040% or less.

[Ca: 0.050% or less]

Ca, like Mg, is an element capable of controlling the form of sulfide by the addition of a trace amount. In order to obtain the effect of addition, Ca is preferably 0.001% or more. More preferably, it is 0.003% or more. On the other hand, if it exceeds 0.050%, a coarse Ca oxide is generated and becomes a starting point of occurrence of cracking in cold working, so Ca should be 0.050% or less. Preferably 0.040% or less.

[Y: 0.050% or less]

Y, like Mg and Ca, is an element capable of controlling the form of sulfide by the addition of a trace amount. In order to obtain the effect of addition, it is preferable that Y is 0.001% or more. More preferably, it is 0.003% or more. On the other hand, if it exceeds 0.050%, a coarse Y oxide is produced, which is a starting point of cracking in cold working, so that Y is 0.050% or less. Preferably 0.035% or less.

[Zr: 0.050% or less]

Zr, like Mg, Ca, and Y, is an element capable of controlling the form of sulfide by the addition of a trace amount. In order to obtain the effect of addition, Zr is preferably 0.001% or more. More preferably, it is 0.004% or more. On the other hand, when it exceeds 0.050%, a coarse Zr oxide is generated and becomes a starting point of cracking in cold working, so that Zr is 0.050% or less. It is preferably 0.045% or less.

[La: 0.050% or less]

La is an element effective for controlling the shape of the sulfide by the addition of a trace amount, and is an element which is segregated in grain boundaries and causes a decrease in the ratio of the number of grain boundaries. In order to obtain the effect of addition, La is preferably 0.001% or more. More preferably, it is 0.004% or more. On the other hand, if it exceeds 0.050%, the ratio of the number of intergranular carbides decreases and the cold workability decreases. Therefore, La should be 0.050% or less. It is preferably 0.045% or less.

[Ce: 0.050% or less]

As with La, Ce is an element capable of controlling the form of sulfide by the addition of a trace amount, and is also an element which is segregated in grain boundaries and causes a decrease in the ratio of the number of intergranular carbides. In order to obtain the effect of addition, Ce is preferably 0.001% or more. More preferably, it is 0.004% or more. On the other hand, if it exceeds 0.050%, the ratio of the number of intergranular carbides decreases and the cold workability decreases. Therefore, Ce is made 0.050% or less. And preferably 0.046% or less.

The balance of the composition of the inventive steel sheet is Fe and inevitable impurities.

Instead of the above-mentioned part of Fe, one or two of Ti and B may be contained.

[Ti: 0.10% or less]

Ti is an element effective for controlling the shape of a carbide and is an element contributing to improvement of toughness by making the structure finer. In order to obtain the effect of addition, Ti is preferably 0.001% or more. More preferably, it is 0.005% or more. On the other hand, if it exceeds 0.10%, a coarse Ti oxide is generated and becomes a starting point of cracking in cold working, so that Ti is made 0.10% or less. And preferably 0.08% or less.

[B: 0.0001 to 0.010%]

B is an element which contributes to improvement of toughness by improving the uniformity of the structure during the heat treatment of the component. In order to obtain the effect of addition, B is preferably 0.0001% or more. More preferably, it is 0.0006% or more. On the other hand, if it exceeds 0.010%, a coarse B oxide is generated and becomes a starting point of cracking in cold working, so B is set to 0.010% or less. It is preferably 0.009% or less.

Next, the steel sheet structure of the present invention will be described.

The structure of the steel sheet of the present invention is a structure composed substantially of ferrite and carbide. In addition to cementite (Fe ₃ C), which is a compound of iron and carbon, the carbide may be a compound obtained by substituting an Fe element in cementite with an alloy element such as Mn or Cr or an alloy carbide (M ₂₃ C ₆ , M ₆ C, MC [M: metal element added as Fe and other alloys]).

When the steel sheet is formed into a predetermined shape, a shear band is formed in the macrostructure of the steel sheet, and slip deformation is concentrated in the vicinity of the shear band. The slip deformation is accompanied by the proliferation of dislocations, and a region having a high dislocation density is formed in the vicinity of the shearing stage. Along with the increase of the deformation amount given to the steel sheet, the slip deformation is promoted, and the dislocation density is increased.

In the cold forging, steel processing of more than 1 is carried out in a considerable variation. As a result, in the conventional steel sheet, generation of voids and / or cracks accompanying the increase of the dislocation density can not be prevented, and it is difficult to improve the cold forging in the conventional steel sheet. In order to solve this problem, it is effective to suppress the formation of the shearing stage at the time of molding.

From the viewpoint of the microstructure, the formation of the shearing zone is understood as a phenomenon in which the slip caused by one of the crystal grains propagates continuously to the adjacent crystal grains over the grain boundaries. Therefore, in order to suppress the formation of the shearing zone, it is necessary to prevent the propagation of slip exceeding the grain boundaries.

Carbides in the steel sheet are solid particles which interfere with slippage, and by causing carbides to exist in the ferrite grain boundaries, propagation of slip exceeding grain boundaries can be prevented, formation of a shearing zone can be suppressed, and cold sectioning can be improved. At the same time, the formability of the steel sheet is improved.

The formability of the steel sheet is largely due to the accumulation of deformation (accumulation of dislocations) in the crystal grain, and when the propagation of the deformation to the adjacent crystal grains is inhibited in the crystal grain system, the deformation amount in the crystal grain increases. As a result, the work hardening rate is increased and the moldability is improved.

Based on the theory and the principle, it is considered that the cold workability is strongly influenced by the coverage of the carbide of the ferrite grain boundaries, and therefore it is necessary to measure the coverage rate with high precision.

In order to measure the coating rate of the carbide in the ferrite grain boundaries in the three-dimensional space, serial scanning SEM observation in which sample cutting and observation with FIB is repeated in a scanning electron microscope, or three-dimensional EBSP observation It becomes necessary to have a large measurement time, and accumulation of technical know-how becomes indispensable. We have clarified this by concluding that the general analytical method is not suitable.

As a result of searching for a simple and highly accurate evaluation index, it is possible to evaluate the cold workability by using the ratio of the number of carbides on the ferrite grain boundaries to the number of carbides in the ferrite grains as an index, The inventors have found that when the ratio of the number of carbides in the ferrite grain bound to the number of ferrite grains exceeds 1, the cold workability remarkably improves.

In addition, since buckling, bending, and folding of the steel sheet occurring during cold working are all caused by localization of deformation accompanying formation of the shear band, the presence of the carbide in the ferrite grain boundaries alleviates localization of shear band formation and deformation , Buckling, bending, and folding can be effectively suppressed.

If the spheroidization ratio of the carbide on the grain boundaries is less than 80%, the deformation concentrates locally on the rod-like or plate-like carbide, and voids and / or cracks are likely to be generated, so that the spheroidization ratio of the carbides on the grain boundaries is preferably 80% More preferably, it is 90% or more.

If the average particle diameter of the carbide is less than 0.1 mu m, the hardness of the steel sheet is remarkably increased and the workability is lowered. Therefore, the average particle diameter of the carbide is preferably 0.1 mu m or more. More preferably, it is 0.17 탆 or more. On the other hand, when the average particle diameter of the carbide exceeds 2.0 占 퐉, the coarse carbide is used as a starting point in the cold working to cause cracking, and the cold workability is lowered, so that the average particle diameter of the carbide is preferably 2.0 占 퐉 or less. More preferably 1.95 占 퐉 or less.

The carbide is observed with a scanning electron microscope. Prior to observation, the sample for tissue observation was polished by wet grinding by emery paper and diamond abrasive grains having an average particle size of 1 mu m, and the observation surface was mirror-finished, and then the structure was etched with a 3% nitric acid- I will. The magnification of the observation is selected at a magnification of 3,000 times at which the ferrite and the carbide can be discriminated. At a selected magnification, eight fields of 30 占 퐉 x 40 占 퐉 are randomly photographed in the 1/4 sheet thickness layer.

With respect to the obtained tissue image, the area of each carbide contained in the area is measured in detail by image analysis softwareware represented by Win ROOF (Win ROOF). (= 2 x? (Area / 3.14)) is determined from the area of each carbide, and the average value is taken as the carbide particle diameter. The spheroidization rate of carbide is obtained by calculating the ratio although the ratio of the maximum length to the maximum length in the direction perpendicular to the carbide is approximated to an ellipse having the same area and the same moment of inertia as the equivalent area.

Further, in order to suppress the influence of the measurement error due to noise, the carbide having an area of 0.01 탆 ² or less was excluded from the object of evaluation. The number of carbides present on the ferrite grain boundary phase was counted and the number of carbides on the grain boundary was subtracted from the total number of carbide numbers to obtain the number of carbides in the ferrite grains. Based on the measured number, the ratio of the number of carbides on the ferrite grain boundaries to the carbides in the ferrite grains was determined.

In the structure after annealing the cold-rolled steel sheet, the cold workability can be improved by setting the ferrite grain size to 5.0 탆 or more. If the ferrite grain size is less than 5 占 퐉, the hardness increases and cracks and cracks tend to occur during cold working, so that the ferrite grain size is set to 5 占 퐉 or more. Preferably 7 mu m or more.

On the other hand, if it exceeds 50 탆, the number of carbides on the grain boundaries which suppress slip propagation decreases, and the cold workability decreases. Therefore, the ferrite grain size is set to 50 탆 or less. Preferably not more than 37 mu m.

The ferrite particle diameter is measured by observing the structure of the observation surface with an optical microscope or a scanning electron microscope after polishing the observation surface of the sample with a mirror surface and etching with a 3% nitric acid-alcohol solution in the above-described polishing method, By using the line segment method.

Cementite, which is a carbide of iron, is a hard and brittle structure. When present in the form of pearlite which is a layered structure with ferrite, the steel becomes hard and brittle. Therefore, the pearlite needs to be minimized as much as possible. In the steel sheet of the present invention, the area ratio is set to 6% or less.

Since pearlite has a peculiar lamellar structure, it can be classified by SEM and optical microscope observation. By calculating the area of the lamellar structure in an arbitrary section, the area ratio of pearlite can be obtained.

Further, by setting the Vickers hardness of the steel sheet to 100 HV or more and 170 HV or less, the cold workability can be improved. If the Vickers hardness is less than 100 HV, buckling tends to occur during cold working, so the Vickers hardness should be 100 HV or more. Preferably 110 HV or more.

On the other hand, when the Vickers hardness exceeds 170 HV, ductility is lowered and internal cracking occurs easily during cold working, so that the Vickers hardness is made 170 HV or less. And preferably 168 HV or less.

Next, the production method of the present invention will be described.

The manufacturing method of the present invention is based on the basic idea of controlling the steel sheet by controlling the hot rolling condition and the annealing condition by using the above-mentioned steel sheet having the component composition.

First, a hot rolled steel sheet obtained by continuously casting molten steel having the required composition is provided. The cast piece after continuous casting may be provided to the hot rolling directly, or may be provided to the hot rolling after being once cooled after being heated.

When the billet is once cooled and then heated to be supplied to hot rolling, the heating temperature is preferably 1000 占 폚 to 1250 占 폚, and the heating time is preferably 0.5 hour to 3 hours. When continuous cast steel is provided for direct hot rolling, the temperature of the steel strip to be provided for hot rolling is preferably 1000 ° C or more and 1250 ° C or less.

When the steel strip temperature or the steel strip heating temperature exceeds 1250 占 폚 or the steel strip heating time exceeds 3 hours, decarburization from the surface layer of the steel strip becomes remarkable, and during heating before carburization, the austenite grains in the surface layer of the steel strip become abnormal And the impact resistance is lowered. Therefore, the steel strip temperature or the steel strip heating temperature is preferably 1250 DEG C or less, and the heating time is preferably 3 hours or less. More preferably, it is 1,200 DEG C or less and 2.5 hours or less.

If the steel strip temperature or the steel strip heating temperature is less than 1000 占 폚 or the heating time is less than 0.5 hour, the micro-segregation or macroscopic segregation produced by the casting is not solved and the alloy element such as Si or Mn is locally thickened And the impact resistance is lowered. For this reason, it is preferable that the steel strip temperature or the steel strip heating temperature is 1000 ° C or more, and the heating time is 0.5 hours or more. More preferably 1050 DEG C or more, and 1 hour or more.

Finishing rolling in hot rolling is completed in a temperature range of 750 DEG C or more and 850 DEG C or less. When the finishing rolling temperature is lower than 750 캜, the deformation resistance of the steel sheet increases, the rolling load remarkably increases, the amount of wear of the roll increases, and the productivity is lowered. Further, the recrystallization required for improving the plastic anisotropy does not sufficiently proceed Therefore, the finishing rolling temperature should be 750 ° C or higher. Is preferably 770 DEG C or higher in terms of promoting recrystallization.

If the finish rolling temperature exceeds 850 ° C, a thick scale is generated in the run out table (ROT), scratches are formed on the surface of the steel sheet due to this scale, impact load is applied after cold forging and carburizing heat tempering Cracks tend to occur starting from scratches, so that the impact resistance of the steel sheet is lowered. For this reason, the finishing rolling temperature is set to 850 DEG C or less. Preferably 830 DEG C or less.

When the hot-rolled steel sheet after the finish rolling is cooled in the ROT, the cooling rate is preferably 10 占 폚 / sec or more and 100 占 폚 / sec or less. If the cooling rate is less than 10 ° C / second, a thick scale is generated during cooling, the occurrence of scratches due to the scaling can not be suppressed, and the impact resistance is lowered. Therefore, the cooling rate is preferably 10 ° C / second or more. More preferably not less than 20 ° C / second.

When the steel sheet is cooled at a cooling rate exceeding 100 deg. C / second from the surface layer to the inside of the steel sheet, the outermost surface layer is excessively cooled, and a low temperature transformation structure such as bainite or martensite is generated. When the hot-rolled steel coils cooled at 100 ° C to room temperature after coiling are discharged, microcracks are generated in the low-temperature transformed structure. It is difficult to remove this fine crack by pickling and cold rolling.

Further, when an impact load is applied to the steel sheet after cold forging and carburizing? Tempering, the crack progresses starting from the micro crack, so that the impact resistance is lowered. For this reason, in order to suppress the occurrence of low-temperature transformation textures such as bainite and martensite at the outermost surface of the steel sheet, the cooling rate is preferably 100 ° C / sec or less. More preferably not higher than 90 占 폚 / sec.

The cooling rate is set so that the cooling rate of the hot rolled steel sheet from the cooling facility of each sump section at the time of cooling on the ROT from the time when the hot-rolled steel sheet after passing through the waterless section to the target temperature of water- Indicates the cooling capability to be received and does not indicate the average cooling rate from the start of the cycle to the temperature of the winding by the winder.

The coiling temperature is set to 400 ° C or more and 550 ° C or less. This is a temperature lower than the general coiling temperature, and is a condition that is not normally performed when the content of C is particularly high. By winding the hot-rolled steel sheet produced under the above-described conditions within this temperature range, the steel sheet can be made into a bainite structure in which carbides are dispersed in fine ferrite.

If the coiling temperature is less than 400 캜, austenite which is not transformed before the winding is transformed into hard martensite, cracks are generated in the surface layer of the hot-rolled steel sheet when the hot-rolled steel coils are discharged, and the impact resistance is lowered.

Further, at the time of recrystallization from austenite to ferrite, since the recrystallization driving force is small, orientation of the recrystallized ferrite particles is strongly influenced by the orientation of the austenite grains, making it difficult to randomize the aggregate structure. Therefore, the coiling temperature is set to 400 캜 or higher. Preferably 430 ° C or higher.

When the coiling temperature exceeds 550 캜, pearlite having a large lamellar spacing is produced, and a thick needle-shaped carbide with high thermal stability is produced. The carbide of this bed remains after the two-stage annealing. During the cold forging of a steel sheet, cracks are generated from the carbide of the acicular phase as a starting point.

Further, at the time of recrystallization of ferrite from austenite, on the contrary, the recrystallization driving force becomes excessively large, and even in this case, recrystallized ferrite particles strongly dependent on the orientation of the austenite grains are formed, and the aggregate structure is not randomized. Therefore, the coiling temperature is set to 550 占 폚 or less. Preferably 520 DEG C or less.

After the hot-rolled steel coils are discharged and pickled, annealing in two steps (two-step annealing) is carried out in two temperature zones. By performing the two-stage annealing on the hot-rolled steel sheet, the stability of the carbide is controlled and the generation of carbide in the ferrite grain boundary is promoted.

If cold rolling is performed on the steel sheet after pickling before the annealing treatment, the ferrite grains become finer, so that the steel sheet is hardly softened. Therefore, in the present invention, it is not preferable to carry out the cold rolling before the annealing, and it is preferable to carry out the annealing treatment after the pickling and without cold rolling.

The first-stage annealing is performed in a temperature range of 650 to 720 deg. C, preferably at least A _c1 point. By this annealing, the carbide is coarsened, partially spheroidized, and the alloying element is concentrated in the carbide to increase the thermal stability of the carbide.

In the first-stage annealing, the heating rate up to the annealing temperature (hereinafter referred to as " first-stage heating rate ") is 3 ° C / hour or more and 150 ° C / hour or less. If the heating rate in the first stage is less than 3 ° C / hour, it takes time to raise the temperature and the productivity is lowered. Therefore, the heating rate in the first stage is 3 ° C / hour or more. Preferably 10 ° C / hour or more.

On the other hand, if the heating rate in the first stage exceeds 150 ° C / hour, the temperature difference between the outer peripheral portion and the inner portion of the hot-rolled steel coil increases, and friction scratches or baking caused by the difference in thermal expansion occur, thereby forming irregularities on the surface of the steel sheet. During cold forging or other molding, the unevenness becomes a starting point, cracks are formed, the cold step composition is lowered, and the impact resistance after molding and carburizing step tempering is lowered. Therefore, the first heating rate is 150 ° C / do. Preferably 130 ° C / hour or less.

The annealing temperature in the first-stage annealing (hereinafter referred to as " first-stage annealing temperature ") is 650 ° C or higher and 720 ° C or lower. If the annealing temperature at the first stage is less than 650 캜, the stabilization of the carbide is not sufficient, and it becomes difficult to retain the carbide in the austenite at the second stage annealing. For this reason, the annealing temperature in the first stage is set to 650 ° C or higher. Preferably 670 DEG C or more.

On the other hand, if the annealing temperature at the first stage exceeds 720 deg. C, austenite is generated before the stability of the carbide rises, and the control of the above-described structure change becomes difficult. Therefore, the annealing temperature at the first stage is set to 720 deg. Preferably 700 DEG C or less.

The annealing time in the first-stage annealing (hereinafter referred to as " first-stage annealing time ") is from 3 hours to 60 hours. If the annealing time in the first stage is less than 3 hours, the stabilization of the carbides is not sufficient, and it becomes difficult to retain the carbides in the austenite during the second stage annealing. For this reason, the annealing time in the first stage is 3 hours or more. Preferably 5 hours or more.

On the other hand, if the annealing time in the first stage exceeds 60 hours, further stabilization of the carbides can not be anticipated and the productivity is lowered, so that the annealing time in the first stage is 60 hours or less. Preferably 55 hours or less.

Thereafter, the temperature was raised to a temperature range of 725 to below 790 ℃, preferably from A _c1 point than A ₃ point, to produce an austenite in the tissue. At this time, the carbides in the fine ferrite grains dissolve in the austenite, but the carbides that are coarsened by the first-stage annealing remain in the austenite.

When cooling is performed without performing the second-stage annealing, the ferrite grain size does not become large, and an ideal structure can not be obtained.

The heating rate to the annealing temperature of the second-stage annealing (hereinafter referred to as " second-stage heating rate ") is from 1 ° C / hr to 80 ° C / hr. During the second-stage annealing, austenite is generated from the ferrite grain boundaries to grow. At this time, by slowing the heating rate to the annealing temperature, nucleation of austenite can be suppressed, and it is possible to increase the grain boundary coverage of carbide in a structure formed by gradual cooling after annealing.

Therefore, it is preferable that the heating rate in the second stage is slower. However, if the heating rate is lower than 1 占 폚 / hour, time is required for raising the temperature and productivity is lowered. Preferably 10 ° C / hour or more.

If the heating rate in the second stage exceeds 80 DEG C / hour, the temperature difference between the outer peripheral portion and the inner portion of the hot-rolled steel coil increases, and friction scratches or baking caused by a large thermal expansion difference due to transformation occur, . During cold forging, cracks are generated starting from the unevenness, resulting in a decrease in cold forging and formability, and a decrease in impact resistance after carburizing and tempering. Therefore, the second stage heating rate is set to 80 캜 / hour or less. Preferably 70 DEG C / hour or less.

The annealing temperature (hereinafter referred to as " second-stage annealing temperature ") in the second-stage annealing is 725 ° C or higher and 790 ° C or lower. If the annealing temperature in the second stage is less than 725 DEG C, the amount of austenite to be produced is reduced, and the number of carbides in the ferrite grain boundaries is reduced and the ferrite grain size is reduced after cooling in the second stage after annealing. For this reason, the annealing temperature in the second stage is set to 725 DEG C or higher. Preferably 735 DEG C or more.

On the other hand, when the annealing temperature in the second stage exceeds 790 DEG C, it is difficult to keep the carbide in the austenite, and it becomes difficult to control the structure change, so the annealing temperature in the second stage is set to 790 DEG C or less. Preferably 770 占 폚 or less.

The annealing time (second-stage annealing time) in the second-stage annealing is set to be not less than 3 hours and less than 50 hours. If the annealing time in the second stage is less than 3 hours, the amount of austenite to be produced is small and the dissolution of carbides in the ferrite grains does not progress sufficiently, so that it is difficult to increase the number of carbides on the ferrite grain boundaries and the ferrite grain size becomes small . For this reason, the annealing time in the second stage is set to 3 hours or more. Preferably 5 hours or more.

On the other hand, if the annealing time in the second stage exceeds 50 hours, it becomes difficult to keep the carbide in the austenite and the manufacturing cost also increases, so that the annealing time in the second stage is less than 50 hours. Preferably 40 hours or less.

After the two-stage annealing, the steel sheet is cooled to 650 DEG C at a cooling rate of 1 DEG C / hour or more and 30 DEG C / hour or less.

By slowly cooling the austenite produced in the second stage of annealing by gradual cooling, carbon atoms are adsorbed on the carbide remaining in the austenite while being transformed into ferrite, so that the carbides and austenite cover the ferrite grain boundaries, A structure in which many carbides exist in the grain boundary can be obtained.

For this purpose, the cooling rate is preferably slower, but if it is less than 1 占 폚 / hour, the time required for cooling is increased and the productivity is lowered. Therefore, the cooling rate is 1 占 폚 / hour or more. Preferably 10 ° C / hour or more.

On the other hand, if the cooling rate exceeds 30 DEG C / hour, the austenite is transformed into pearlite, the hardness of the steel sheet increases, the cold step composition decreases, and the impact resistance after carburizing heat tempering decreases. 30 ° C / hour or less. Preferably not higher than 20 ° C / hour.

Further, the steel sheet cooled to 650 deg. C is cooled to room temperature. The cooling rate at this time is not limited.

The atmosphere in the two-stage annealing is not particularly limited to a specific atmosphere. For example, the atmosphere may be at least 95% nitrogen atmosphere, at least 95% hydrogen atmosphere, or atmospheric atmosphere.

INDUSTRIAL APPLICABILITY As described above, according to the manufacturing method of controlling the hot-rolled condition and the annealing condition of the present invention and controlling the texture of the steel sheet, it is possible to provide a steel sheet excellent in moldability during cold forging combined with drawing, It is possible to manufacture a steel sheet excellent in ignitability required for improvement of impact resistance after quenching.

Example

Next, the embodiment will be described, but the level of the embodiment is an example of a condition adopted to confirm the feasibility and effect of the present invention, and the present invention is not limited to this one conditional example. The present invention can adopt various conditions as long as the object of the present invention is achieved without departing from the gist of the present invention.

The evaluation of the cold workability was carried out by taking a JIS No. 5 tensile test specimen from an annealing material having a thickness of 3 mm and performing a tensile test to evaluate the overall elongation in the direction of 0 占 from the rolling direction and in the 90 占 direction in both directions, % Or more, and the difference in total elongation | DELTA EL | in each direction is not more than 4%, the cold workability is superior.

The evaluation of the quenching was carried out by grinding an annealing material having a thickness of 3 mm at a plate thickness of 1.5 mm and holding it at 880 캜 for 10 minutes in a vacuum atmosphere and cooling at a rate of 30 캜 / If it is more than 60%, it is said to be superior.

(Example 1)

The continuous casting pieces (ingot) having the constituent compositions shown in Table 1 were heated at 1240 占 폚 for 1.8 hours and then subjected to hot rolling. After finishing the hot rolling at the temperatures shown in Table 2, And a hot-rolled coil having a thickness of 3.0 mm was produced. The hot-rolled coil was picked up and the hot-rolled coil was charged into the box-shaped annealing furnace. The atmosphere was controlled to 95% hydrogen-5% nitrogen, heated from room temperature to 705 캜, held at 705 캜 for 36 hours, And then heated to 760 DEG C and held at 760 DEG C for 10 hours.

Thereafter, the sample was cooled to 650 ° C at a cooling rate of 10 ° C / hour, and then cooled to room temperature to obtain a sample for evaluation of characteristics. The texture of the sample was measured by the above-described method.

Table 2 shows the results of measuring or evaluating the Vickers hardness of the produced samples, the ratio of the number of carbides on the ferrite grain boundaries to the number of carbides in the ferrite grains, the pearlite area ratio, the cold workability and the shattering.

As shown in Table 2, the inventive steels B-1, E-1, F-1, H-1, J-1, K-1, L-1, M-1, N-1, The ratio of the number of carbides in the ferrite grain boundaries to the number of carbides in the ferrite grains is not less than 1, Is more than 1, Vickers hardness is not more than 170 HV, and is excellent in cold workability and quenching.

On the other hand, the comparative steel G-1 had a high C content and a poor cold workability. The comparative steel O-1 has a high amount of Mo and Cr and a high stability of carbide, so that the carbide is not dissolved at the time of casting, the amount of austenite produced is small, and quenching is poor.

The comparative steels Q-1 and AD-1 have a high amount of Si and Al and a high A3 point, so the amount of austenite produced is small and the quenching is poor. In Comparative Example U-1, the amount of S was high, coarse MnS was produced in the steel, and the cold workability was low. In Comparative Example AA-1, the amount of Mn is low and the quenching is inferior.

In Comparative Example I-1, the finishing temperature of hot rolling was low and the productivity was lowered. In Comparative Example D-1, the finishing temperature of hot rolling was high, and scale scratches were formed on the surface of the steel sheet. In Comparative Examples C-1 and S-1, the coiling temperature of hot-rolled steel was low and the low-temperature transformed structures such as bainite and martensite were increased and became brittle.

In Comparative Examples A-1 and V-1, coarse carbide having high coiling temperature of hot-rolled, high pearlite with thick lamellar spacing in hot-rolled structure and high thermal stability was produced, and even after this two- It remained in the steel sheet, and the cold workability was lowered.

(Example 2)

In order to investigate the influence of the annealing conditions, the steel strips having the composition shown in Table 1 were heated at 1240 占 폚 for 1.8 hours and then subjected to hot rolling. After finishing the hot rolling at 820 占 폚, Cooled to 520 deg. C, rolled at 510 deg. C to obtain a hot-rolled coil having a thickness of 3.0 mm, and subjected to box annealing in a two-step step type under the annealing conditions shown in Table 3, .

Table 3 shows the results of measuring or evaluating the carbide diameter, the ferrite grain size, the Vickers hardness, the ratio of the number of carbides on the ferrite grain boundaries to the number of carbides in the ferrite grains, the pearlite area ratio, the cold workability, .

2, C-2, E-2, F-2, H-2, I-2, J-2, K-2, M-2, N-2 and R 2, S-2, V-2, Z-2 and AC-2 all have a ratio of the number of carbides on the ferrite grains to the number of carbides in the ferrite grains exceeding 1, a Vickers hardness of 170 HV or less , Excellent in cold workability and quenching.

On the other hand, the comparative steel G-1 had a high C content and a poor cold workability. The comparative steel O-1 had a high amount of Mo and Cr, and the cold workability deteriorated. Further, since the stability of the carbide is high, the carbide is not dissolved at the time of casting, the austenite formation amount is small, and the quenching property is inferior.

The comparative steel Q-1 had a high amount of Si and a high hardness of ferrite, so that workability deteriorated. Also, since the A3 point is high, the austenite production amount is small and the quenching is inferior during casting. The comparative steel AD-1 has a high amount of Al and a high A3 point, so a small amount of austenite is produced at the time of casting, and quenching is inferior. In the comparative steel U-1, the amount of S was high, coarse MnS was generated in the steel, and cold workability deteriorated. The comparative steel AA-1 has a low Mn content and a low quenching characteristic.

The comparative steel T-2 had a low holding temperature at the first stage annealing of the two-step step type box annealing, The coarsening treatment of the carbide at a temperature lower than the temperature is insufficient and the thermal stability of the carbide is insufficient so that the carbide remaining at the second stage annealing is decreased and the pearlite transformation can not be suppressed in the post- The cold workability was lowered.

The comparative steel A-2 has a high holding temperature at the first stage annealing of the two-step step type box annealing, austenite is generated during annealing, the stability of the carbide can not be enhanced, Carbide was reduced, pearlite transformation was not able to be suppressed in the post-cooling structure, and the cold workability was lowered.

The comparative steel L-2 had a short holding time at the first-stage annealing in the two-step step type box annealing, insufficient coarsening treatment of the carbide at the Ac1 temperature or lower, and insufficient thermal stability of the carbide, The remaining carbides at the second stage of annealing were decreased and the pearlite transformation could not be suppressed in the post-cooling structure, and the cold workability was lowered.

The comparative steel W-2 had a longer holding time at the first-stage annealing at the time of the two-step annealing, and the productivity was lowered. In Comparative Strength X-2, the holding temperature at the second stage annealing at the time of the two-step annealing is low, the amount of austenite is small and the ratio of the number of carbides in the grain boundary can not be increased, .

The comparative steel AB-2 had a higher holding temperature at the second stage annealing of the two-step step type box annealing, and the carbide remained so accelerated, so that the residual carbide was reduced and the pearlite transformation And the cold forging workability was deteriorated.

The comparative steel P-2 had a low holding temperature at the second stage of annealing in the two-step step type box annealing, a small amount of austenite was produced, and the ratio of the number of carbides in the ferrite grain boundary could not be increased, The workability was deteriorated. The comparative steel Y-2 had a longer holding time at the second stage annealing of the two-step step type box annealing and accelerated the dissolution of the carbide, so that the remaining carbide was reduced and the pearlite transformation And the cold forging workability was deteriorated.

In Comparative Strength D-2, the cooling rate from 650 ° C to the end of the second stage annealing of the two-step step type box annealing was large, pearlite transformation occurred during cooling, and cold workability deteriorated.

INDUSTRIAL APPLICABILITY As described above, according to the present invention, it is possible to produce and provide a steel sheet excellent in moldability and abrasion resistance. INDUSTRIAL APPLICABILITY The steel sheet of the present invention is a steel sheet suitable as a material for automobile parts, cutlery, and other mechanical parts, which are manufactured through processing steps such as punching, bending, and pressing.

Claims (3)

In terms of% by mass,
C: 0.10 to 0.40%,
Si: 0.01 to 0.30%
Mn: 1.09 to 2.00%
P: 0.020% or less,
S: 0.010% or less,
Al: 0.001 to 0.10%
N: 0.010% or less,
O: 0.020% or less,
0.50% or less of Cr,
Mo: 0.10% or less,
Nb: 0.10% or less,
V: 0.10% or less,
Cu: not more than 0.10%
W: 0.10% or less,
Ta: 0.10% or less,
Ni: 0.10% or less,
Sn: 0.050% or less,
Sb: 0.050% or less,
As: 0.050% or less,
Mg: 0.050% or less,
Ca: 0.050% or less,
Y: 0.050% or less,
Zr: 0.050% or less,
La: 0.050% or less,
Ce: not more than 0.050%
And the remainder being Fe and inevitable impurities,
The metal structure of the steel sheet
The ratio of the number of carbides on the ferrite grain bound to the number of carbides in the ferrite grains is more than 1,
The spheroidization ratio of the carbide on the grain boundary is 80% or more,
The average particle diameter of the carbide is 0.1 占 퐉 or more, 2.0 占 퐉 or less,
A ferrite grain size of 5 탆 or more and 50 탆 or less and
If the area ratio of pearlite is less than 6%
Lt; / RTI >
Wherein the steel sheet has a Vickers hardness of 100 HV or more and 170 HV or less,
The overall elongation in the 0 占 direction from the rolling direction and in the 90 占 direction from the rolling direction is 35% or more in both directions
. The method according to claim 1,
Ti: 0.10% or less and
B: 0.010% or less,
Or a mixture of two or more of them. A manufacturing method for manufacturing the steel sheet according to any one of claims 1 to 3,
A hot-rolled steel sheet is obtained by subjecting a steel sheet having the composition described in claim 1 or 2 to hot rolling to finish finish rolling in a temperature range of 750 ° C to 850 ° C,
The hot-rolled steel sheet is rolled at a temperature of 400 ° C or more and 550 ° C or less,
The picked hot-rolled steel sheet was pickled,
The first hot rolled steel sheet is subjected to the first stage annealing in which the hot rolled steel sheet is held in a temperature range of 650 DEG C or higher and 720 DEG C or lower for 3 hours or more and 60 hours or less,
Stage annealing in which the hot-rolled steel sheet is held in a temperature range of 725 DEG C to 790 DEG C for 3 hours to 50 hours,
The hot-rolled steel sheet after annealing is cooled to 650 占 폚 at a cooling rate of 1 占 폚 / hour or more and 30 占 폚 / hour or less
Wherein the steel sheet is produced by a method comprising the steps of: