KR20170138508A - Steel sheet and manufacturing method thereof - Google Patents

Abstract

A steel sheet having a predetermined composition and having an average grain size of 0.4 to 2.0 탆, an area ratio of pearlite of 6% or less, an average grain size of the carbides in the ferrite grains Wherein the ratio of the number of carbides on the ferrite grain bound to the number of carbides is greater than 1 and the Vickers hardness is from 100 HV to 180 HV.

Description

Steel sheet and manufacturing method thereof

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

The steel sheet containing 0.1 to 0.4% by carbon of carbon is subjected to press forming, hole expanding molding, bending molding, drawing molding, thickness increase and thickness reduction molding, and cold forging in combination thereof. Gears, clutches, and the like. The amount of deformation accumulated in the material increases in the cold forging compared to the conventional hot forging or the like and causes cracking of the material or buckling at the time of forming, thereby causing deterioration of the part characteristics.

Particularly, after carrying out carburization and tempering in a molding material in order to obtain wear resistance, residual stress is generated by heat treatment, so that cracks are generated and progress from the cracks and buckling portions. It is required to acquire an impact resistance characteristic to prevent brittle fracture against a load with a large load instantaneously due to the engagement of the gear at the time of starting. Therefore, the above steel sheet has excellent cold forging It is required to secure the impact resistance characteristics after carburizing and tempering.

Heretofore, many proposals have been made on techniques for improving the cold forging of a steel sheet and the impact resistance characteristics after carburizing (see, for example, Patent Documents 1 to 5).

For example, Patent Document 1 discloses a steel for machine structural use in which the toughness is improved by suppressing the coarsening of crystal grains in the carburizing heat treatment. The steel contains 0.10 to 0.30% of C, 0.05 to 2.0% of Si, 0.05 to 2.0% of Si, 0.10 to 0.50% P, 0.030% or less of S, 0.030% or less of S, 1.80 to 3.00% of Cr, 0.005 to 0.050% of Al, 0.02 to 0.10% of Nb and 0.0300% or less of N, And a ferrite-pearlite structure before cold working, and an average value of ferrite grain sizes of 15 mu m or more.

Patent Document 2 discloses a steel which is excellent in cold workability and carburizing property and contains 0.15 to 0.40% of C, 1.00% or less of Si, 0.40% or less of Mn, not more than 0.02% of sol. : 0.005 to 0.050%, the balance being Fe and inevitable impurities, and having a structure mainly composed of a ferrite phase and a graphite phase.

Patent Document 3 discloses a steel material for a carburizing bevel gear having excellent impact strength, a tough carburizing bevel gear, and a manufacturing method thereof.

Patent Document 4 discloses that cold forging is performed after spheroidizing annealing to suppress the coarsening of crystal grains even after carburizing with excellent workability for parts manufactured in the carburizing and quenching step and excellent impact resistance, A steel for a carburizing part is disclosed.

Patent Document 5 discloses a cold tool steel for carburizing a plasma which contains 0.40 to 0.80% of C, 0.05 to 1.50% of Si, 0.05 to 1.50% of Mn and 1.8 to 6.0% of V, 0.10 to 2.50% : 0.1 to 2.0% and Mo: 3.0% or less, and the balance of Fe and unavoidable impurities.

Japanese Patent Application Laid-Open No. 2013-040376 Japanese Patent Application Laid-Open No. 06-116679 Japanese Patent Application Laid-Open No. 09-201644 Japanese Patent Application Laid-Open No. 2006-213951 Japanese Patent Application Laid-Open No. 10-158780

The structure of the steel for machine structural use of Patent Document 1 is a structure of ferrite + pearlite and has a large hardness in comparison with the structure of ferrite + cementite, so that it is not possible to suppress the malfunction of the die during cold forging, It can not be said to be a mechanical structural steel having excellent composition.

In the steel of Patent Document 2, the graphitization treatment of cementite requires annealing at a high temperature, so that it is not possible to suppress the decrease of the yield and the increase of the production cost.

The manufacturing method of Patent Document 3 is required to further perform hot forging after cold forging and carburizing, and hot forging is essential, so that it is not a manufacturing method which leads to a low cost.

It is not clear whether or not the steel for carburizing parts in Patent Document 4 can exhibit the same effect in the cold forging to which a large deformation is imparted and the specific method of controlling the structure and the structure of the steel are unclear. It is not a steel which exhibits excellent workability even in the case of forging in which a large deformation is imparted by cold, such as forging.

Patent Document 5 does not disclose any knowledge and technique regarding the optimum components and the morphology of the steel for improving the formability of the steel, particularly the cold syngas.

The present invention provides a steel sheet having excellent impact resistance characteristics after cold forging and carburizing surface tempering and particularly suitable for obtaining parts such as high cycle gears by plate forming and a method of manufacturing the steel sheet To be a problem.

In order to solve the above problems and obtain a steel sheet suitable for a material such as a drivetrain part or the like, it is necessary to increase the grain size of ferrite in the steel sheet containing C necessary for enhancing the shading and to increase the carbide (mainly cementite) It can be understood that it is necessary to reduce the pearlite structure. 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.

By appropriately dispersing the carbides in the metal structure, it is an indispensable structure for the drivetrain parts, because it can impart excellent abrasion resistance and electric fatigue characteristics while maintaining the material formability. The presence of carbide in the ferrite grain boundaries prevents the propagation of slip exceeding the grain boundaries to suppress the formation of the shearing zone and improves the cold-rolled composition. At the same time, Thereby improving the formability of the steel sheet.

However, the cementite is a hard and brittle structure, and if present in the form of pearlite, which is a layered structure with ferrite, the steel is hard and embrittled, so it needs to be present in spherical form. Considering the occurrence of cracking at the time of cold forging or forging, the grain size needs to be in an appropriate range.

However, a manufacturing method for realizing the above-described structure has not been disclosed so far. Therefore, the present inventors 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 the hot rolling is rolled at a relatively low temperature (400 캜 to 550 캜) 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, a part of the ferrite grains are partially austenitized by annealing at a temperature between the Ac1 point and Ac3 point (so-called two-phase region of ferrite and austenite) as the second stage annealing. The ferrite grains are slowly cooled and left behind, and ferrite is transformed into ferrite grains while the austenite is used as nuclei to obtain cementite in the grain boundaries while obtaining a large ferrite phase, whereby the above-described structure can be realized.

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 devised singly, and it can be realized by achieving optimization in a so- did.

In order to improve the drawability during cold forging, it is necessary to reduce the plastic anisotropy and it is important to adjust the hot rolling conditions for this improvement.

The present invention is based on these findings, and its main points are as follows.

(1) A ferritic stainless steel comprising (1) a ferritic stainless steel having a composition of 0.10 to 0.40% by mass, 0.01 to 0.30% by Si, 0.30 to 1.00% by Mn, 0.001 to 0.10% by mass, 0.50 to 2.00% Pb: not more than 0.050%, Sn: not more than 0.050%, Sb: not more than 0.050%, P: not more than 0.020%, S: not more than 0.010%, N: not more than 0.020%, O: not more than 0.020% 0.10% or less of Cu, 0.10% or less of W, 0.10% or less of Ta, 0.10% or less of Ni, 0.050% or less of Mg, Carbon steel sheet is a low carbon steel plate containing not more than 0.050% of Y, not more than 0.050% of Y, not more than 0.050% of Zr, not more than 0.050% of La and not more than 0.050% of Ce and the balance Fe and impurities, Of 0.4 to 2.0 탆, a pearlite area ratio of 6% or less and a ratio of the number of carbides of the ferrite grain boundaries to the number of carbides in the ferrite grains exceeds 1, and the Vickers hardness of the low carbon steel sheet is 100 HV to 180 HV Steel sheet according to claim.

(2) A method of manufacturing the steel sheet according to (1), wherein the steel strip having the component composition of (1) is subjected to hot rolling to finish hot rolling at a temperature range of 650 ° C to 950 ° C to obtain a hot- , Hot rolling the hot rolled steel sheet at a temperature of not less than 400 ° C and not more than 600 ° C, pickling the hot rolled steel sheet, picking the pickled hot rolled steel sheet at a heating rate of not less than 30 ° C / The annealing is performed at a temperature of 725 DEG C or more and 790 DEG C or more at a heating rate of 1 DEG C / hour or more and 80 DEG C / hour or less, And annealing is performed for 3 hours to 50 hours, and 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 100 占 폚 / hour or less Characterized by a cold stage And a method for producing the steel sheet.

INDUSTRIAL APPLICABILITY According to the present invention, it is possible to provide a steel sheet excellent in impact resistance characteristics after cold forging and carburizing surface tempering, and particularly suitable for obtaining parts such as high cycle gears by plate forming, and a method for producing the same.

Fig. 1 is a diagram schematically showing an outline of a cold forging test and a form of a crack introduced in cold forging. (a) shows the test piece cut from the hot-rolled steel sheet, (b) shows the shape of the test piece after cold forging, and (c) shows the cross-sectional shape of the test piece after cold forging.
2 is a diagram schematically showing an outline of a drop test in which the impact resistance characteristic of a sample subjected to carburizing surface tempering is evaluated.
3 is a graph showing the relationship between the ratio of the number of grain boundaries to the number of carbides in the grain, the crack length of the cold forging test piece, and the impact resistance characteristic after carburizing surface tempering.
4 is a diagram showing the relationship between the ratio of the number of intergranular carbides to the number of carbides in the particle, the crack length of the cold forging test piece, and the impact resistance characteristic after carburizing surface tempering.

Hereinafter, the present invention will be described in detail. First, the reason for limiting the composition of the inventive steel sheet will be described. Here, "% " with respect to the component composition means " mass% ".

[C: 0.10 to 0.40%]

C is an element effective for forming a carbide in a steel, strengthening the steel and making the ferrite particles finer. In order to suppress the generation of crepe during cold working and to secure the surface aesthetics of the cold forged part, it is necessary to suppress coarsening of ferrite grain size. When the content is less than 0.10%, the volume fraction of carbide is insufficient. The coarsening can not be suppressed, so C should be 0.10% or more. It is preferably at least 0.11%.

On the other hand, if it exceeds 0.40%, the volume ratio of carbide increases, and when a load is instantaneously loaded, a large amount of cracks are generated as a starting point of fracture, resulting in a decrease in the impact resistance characteristic. do. 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 reduce the number of carbides in the ferrite grains to obtain the deoxidizing effect and to increase the number of carbides on the ferrite grain boundaries, an austenite phase is generated during the annealing by a two-step step annealing, It is necessary to slowly cool and accelerate the generation of carbides into the ferrite grain boundary.

If the Si content exceeds 0.30%, the ductility of the ferrite decreases, cracks tend to occur at the time of cold forging, and the impact resistance characteristics after cold forging and carburizing finish tempering deteriorate. And preferably 0.28% or less.

Si is preferably as small as possible, but reduction to less than 0.01% leads to a significant increase in refining cost, so Si is set to 0.01% or more. It is preferably 0.02% or more.

[Mn: 0.30 to 1.00%]

Mn is an element that controls the shape of the carbide in a two-step step annealing. If it is less than 0.30%, it is difficult to generate carbide on the ferrite grain boundary in the gradual cooling after the second-stage annealing, so that Mn is set to 0.30% or more. It is preferably 0.33% or more.

On the other hand, if it exceeds 1.00%, the toughness after carburizing quenching is reduced, so that the Mn content is 1.00% or less. Preferably, it is 0.96% or less.

[Al: 0.001 to 0.10%]

Al acts as a deoxidizer for steel to stabilize ferrite. If it is less than 0.001%, the effect of addition is not sufficiently obtained, so that the content of Al is 0.001% or more. It is preferably 0.004% or more.

On the other hand, if it exceeds 0.10%, the ratio of the number of carbides on the grain boundary is decreased and the crack length during cold forging is increased, so that the Al content is 0.10% or less. It is preferably 0.09% or less.

[Cr: 0.50 to 2.00%]

Cr and Mo are elements improving toughness. Cr is an element effective for stabilizing carbides at the time of heat treatment. If it is less than 0.50%, it is difficult to retain the carbide at the time of carburizing, the coarsening of the austenite grain size in the surface layer is caused, and the impact resistance characteristic is lowered. It is preferably not less than 0.52%.

On the other hand, if it exceeds 2.00%, the amount of Cr to be concentrated in the carbide increases, and a large amount of fine carbide remains in the austenite phase produced by the two-step step annealing. Therefore, The ratio of the increase in hardness to the number of intergranular carbides decreases, and the cold-rolled steel composition decreases, so that the Cr content is made to be 2.00% or less. Preferably 1.94% or less.

[Mo: 0.001 to 1.00%]

Mo is an effective element for controlling the morphology of carbides. If it is less than 0.001%, the effect of addition is not sufficiently obtained, so that Mo is set to 0.001% or more. And preferably 0.017% or more.

On the other hand, if it exceeds 1.00%, Mo is concentrated in the carbide, and stable carbide is increased in the austenite phase. Therefore, carbide is present also in the grain after the slow cooling to cause increase in hardness and decrease in the number ratio of intergranular carbides , And the composition of the cold step is lowered, so that the Mo content is 1.00% or less. It is preferably 0.94% or less.

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

[P: 0.020% or less]

P is segregated at the ferrite grain boundaries and is an element inhibiting the generation of grain boundary carbides. The smaller the number, the better. The content of P may be zero. However, in order to obtain a high purity of less than 0.0001% in the refining process, a long time is required for refining and a substantial increase in production cost is caused, so that the practical lower limit is 0.0001 to 0.0013%.

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

[S: 0.010% or less]

S is an impurity element forming a non-metallic inclusion such as MnS. Since nonmetallic inclusions are a starting point of crack generation in cold forging, the smaller the S, the better. The content of S may be zero, but if the S is reduced to less than 0.0001%, the refining cost greatly increases, so that the practical lower limit is 0.0001 to 0.0012%.

On the other hand, if it exceeds 0.010%, the crack length during cold forging is increased, so S is set to 0.010% or less. It is preferably 0.009% or less.

[N: 0.020% or less]

N is an element that segregates into a ferrite grain boundary and inhibits the generation of carbide on grain boundaries. The smaller the number, the better. The content of N may be zero, but if it is reduced to less than 0.0001%, the refining cost greatly increases, so that the actual lower limit is 0.0001 to 0.0006%.

On the other hand, if it exceeds 0.020%, the ratio of the number of carbides on the ferrite grain boundaries to the number of carbides in the ferrite grains becomes less than 1, and the cold hardening is reduced even when the two-phase region annealing and slow cooling are performed. Or less. It is preferably 0.017% or less.

[O: 0.0001 to 0.020%]

O is an element that forms oxides in the steel. Since the oxide present in the ferrite particles is a site for producing carbide, it is preferable that the oxide is small. The content of O may be zero. However, if O is reduced to less than 0.0001%, the refining cost is significantly increased, so that the actual lower limit is 0.0001 to 0.0006%.

On the other hand, if it exceeds 0.020%, the ratio of the number of carbides on the ferrite grain boundaries to the number of carbides in the ferrite grains becomes less than 1, and the cold step composition is lowered, so that O is 0.020% or less. It is preferably 0.017% or less.

[Ti: 0.010% or less]

Ti is an element that is important for controlling the shape of the carbide and is an element that promotes the formation of carbide in the ferrite grains by the inclusion of a large amount. The content of Ti may be zero, but if it is reduced to less than 0.0001%, the polishing cost greatly increases, so that the practical lower limit is 0.0001 to 0.0006%.

On the other hand, if it is more than 0.010%, the ratio of the number of carbides on the ferrite grain boundaries to the number of carbides in the ferrite grains becomes less than 1, and the cold step composition is lowered. It is preferably 0.007% or less.

[B: 0.0005% or less]

B is an element effective for controlling slippage of potential in cold forging. Since the activity of the slip system is limited by the inclusion of a large amount, it is preferable that B is small. The content of B may be zero. Attention must be paid to the detection of less than 0.0001% B and, depending on the analytical equipment, the detection limit is below the detection limit.

On the other hand, if it exceeds 0.0005%, the cross slip of the potential at the shearing stage formed by the cold forging is suppressed, the deformation concentrates locally and cracks are generated, so that B is 0.0005% or less. It is preferably 0.0005% or less.

[Sn: 0.050% or less]

Sn is an element incorporated from a steel raw material (scrap), and the smaller the number, the better. The content of Sn 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%.

On the other hand, if it exceeds 0.050%, the ferrite becomes brittle and the cold step composition decreases, so that the content of Sn is 0.050% or less. Preferably 0.048% or less.

[Sb: 0.050% or less]

Sb, like Sn, is an element incorporated from steel raw material (scrap). Sb is segregated at the grain boundaries and decreases the ratio of the number of intergranular carbides. The content of Sb may be zero, but if it is reduced to less than 0.001%, the refining cost is significantly increased, so that the practical lower limit is 0.001 to 0.002%.

On the other hand, if it exceeds 0.050%, the cold-rolled composition is lowered, so that Sb is 0.050% or less. Preferably 0.048% or less.

[As: 0.050% or less]

As, like Sn and Sb, is an element incorporated from a steel raw material (scrap). As is segregated at grain boundaries and decreases the number ratio of intergranular carbides. The content of As may be zero, but if it is reduced to less than 0.001%, the refining cost greatly increases, so that the practical lower limit is 0.001 to 0.002%.

On the other hand, if it exceeds 0.050%, the ratio of the number of intergranular carbides decreases and the cold step composition decreases, so that the content of As is 0.050% or less. It is preferably 0.045% or less.

The steel sheet of the present invention may contain the following elements for the purpose of improving the cold-rolling and other properties, while the above elements are used as basic elements. The following elements are not essential for attaining the effect of the present invention, and therefore the content may be zero.

[Nb: 0.10% or less]

Nb is an element effective for controlling the morphology of carbides and is an element contributing to improvement of toughness by making the structure finer. If it is less than 0.001%, the effect of addition is not sufficiently obtained, and therefore 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 number ratio of intergranular carbides decreases, and the cold step composition is lowered, so that Nb is 0.10% or less. It is preferably 0.09% 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. If it is less than 0.001%, the effect of addition is not sufficiently obtained, and therefore 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 step composition decreases, so V is 0.10% or less. It is preferably 0.09% or less.

[Cu: 0.10% or less]

Cu forms fine precipitates and contributes to improvement of strength. If it is less than 0.001%, the effect of improving the strength can not be sufficiently obtained. Therefore, the content of Cu is preferably 0.001% or more. More preferably, it is 0.008% or more.

On the other hand, if it exceeds 0.10%, the hot brittleness is manifested during hot rolling, and the productivity is lowered, so that the content of Cu is 0.10% or less. It is preferably 0.09% or less.

[W: 0.10% or less]

W, like Nb and V, is an element effective for controlling the shape of carbide. If it is less than 0.001%, the effect of addition is not sufficiently obtained, and therefore it is preferable that W is 0.001% or more. More preferably, it is 0.003% or more.

On the other hand, when the content exceeds 0.10%, a large number of fine W carbides are precipitated, the strength excessively increases, the number ratio of intergranular carbides decreases, and the cold hardening decreases, so W is 0.10% or less. And preferably 0.08% or less.

[Ta: 0.10% or less]

Ta, like Nb, V, and W, is an element effective for controlling the shape of carbide. If it is less than 0.001%, the effect of addition is not sufficiently obtained, and therefore it is preferable that Ta is 0.001% or more. It is preferably at least 0.007%.

On the other hand, if it exceeds 0.10%, a large number of fine W carbides are precipitated, the strength excessively increases, the number ratio of intergranular carbides decreases, and the cold step composition is lowered. It is preferably 0.09% or less.

[Ni: 0.10% or less]

Ni is an effective element for improving the impact resistance of a component. If it is less than 0.001%, the effect of addition is not sufficiently obtained, and therefore it is preferable that the content of Ni is 0.001% or more. More preferably, it is 0.002% or more.

On the other hand, if it exceeds 0.10%, the number ratio of intergranular carbides decreases and the cold-rolled steel composition decreases, so that the content of Ni is 0.10% or less. It is preferably 0.09% or less.

[Mg: 0.050% or less]

Mg is an element capable of controlling the form of sulfide by adding a trace amount of Mg. If it is less than 0.0001%, the effect of addition is not sufficiently obtained, and therefore, 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 step composition is lowered, so that the Mg content is 0.050% or less. Preferably 0.049% 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. If it is less than 0.001%, the effect of addition is not sufficiently obtained, and Ca is preferably 0.001% or more. More preferably, it is 0.003% or more.

On the other hand, if it exceeds 0.050%, coarse Ca oxides are produced and Ca becomes a starting point of cracking at the time of cold forging, so Ca should be 0.050% or less. Preferably 0.04% 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. If it is less than 0.001%, the effect of addition is not sufficiently obtained. Therefore, Y is preferably 0.001% or more. More preferably, it is 0.003% or more.

On the other hand, if it is more than 0.050%, coarse Y oxide is produced and Y becomes 0.050% or less since it becomes a starting point of cracking in cold forging. It is preferably 0.031% or less.

[Zr: 0.050% or less]

Like Mg, Ca and Y, Zr is an element capable of controlling the form of sulfide by the addition of a trace amount. If it is less than 0.001%, the effect of addition is not sufficiently obtained, and therefore Zr is preferably 0.001% or more. More preferably, it is 0.004% or more.

On the other hand, if it exceeds 0.050%, a coarse Zr oxide is produced and Zr becomes 0.050% or less since it becomes a starting point of crack generation in cold forging. It is preferably 0.045% or less.

[La: 0.050% or less]

La is an element effective for controlling the shape of a sulfide by the addition of a trace amount and is an element segregating in grain boundaries and lowering the ratio of the number of intergranular carbides. If it is less than 0.001%, the shape control effect can not be sufficiently obtained. Therefore, La is preferably 0.001% or more. More preferably, it is 0.003% or more.

On the other hand, if it is more than 0.050%, the number ratio of intergranular carbides is lowered and the cold step composition is lowered, so that La is set to 0.050% or less. Preferably 0.047% or less.

[Ce: 0.050% or less]

Like La, Ce is an element capable of controlling the form of sulfide by the addition of a trace amount, and is an element segregating in grain boundaries and lowering the ratio of the number of intergranular carbides. If the content is less than 0.001%, the shape control effect can not be sufficiently obtained. Therefore, Ce is preferably 0.001% or more. More preferably, it is 0.003% or more.

On the other hand, if it exceeds 0.050%, the number ratio of intergranular carbides is lowered and the cold-rolled steel composition is lowered, so that Ce is 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.

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 consisting essentially of ferrite and carbide. In addition to cementite (Fe ₃ C), which is a compound of iron and carbon, carbide is a compound in which Fe atoms in cementite are substituted with Mn, Cr, alloy carbide (M ₂₃ C ₆ , M ₆ C, And other metal elements).

When the steel sheet is formed into a predetermined part shape, a shear band is formed in the macrostructure of the steel sheet, and a 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. With an increase in the deformation amount given to the steel sheet, the slip deformation is promoted, and the dislocation density is increased.

For cold forging, more than 1 steel is machined in equivalent variants. As a result, in the conventional steel sheet, generation of voids and / or cracks accompanying an increase in dislocation density can not be prevented, and it is difficult to improve the cold-rolled steel sheet.

In order to solve this difficult problem, it is effective to suppress the formation of the shear band at the time of molding. From the viewpoint of the microstructure, the formation of the shear band can be understood as a phenomenon in which the slip generated in any one of the particles continuously propagates to the adjacent particles over the grain boundaries. Therefore, in order to suppress the formation of the shearing zone, it is necessary to prevent the propagation of the slip over the grain boundaries.

The carbide in the steel sheet is a solid particle which interferes with slippage, and by forming the carbide in the ferrite grain boundary, it is possible to suppress the formation of the shearing zone and to improve the cold hardening.

In order to obtain such an effect, it is necessary to disperse the carbide to a proper size in the metal structure. Therefore, the average particle diameter of the carbide is set to 0.4 μm or more and 2.0 μm or less. If the particle diameter of the carbide is less than 0.4 mu m, the hardness of the steel sheet is remarkably increased and the cold step composition is lowered. More preferably not less than 0.6 mu m.

On the other hand, when the average particle diameter of the carbide exceeds 2.0 占 퐉, the carbide becomes a starting point of the crack during cold forming. More preferably 1.95 占 퐉 or less.

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 is 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 unique lamellar structure, it can be strictly distinguished by SEM and optical microscopic observation. By calculating the area of the lamellar structure in an arbitrary section, the area ratio of pearlite can be obtained.

Based on the theory and the principle, it is considered that the cold-rolled steel is strongly influenced by the coverage of the carbide of the ferrite grain boundaries, and measurement of the high precision is required. However, the covering ratio of the carbide to the ferrite grain boundaries in the three- , Serial scanning SEM observation or 3-dimensional EBSP observation in which sample cutting and observation are repeatedly performed by FIB in a scanning electron microscope is required, and a large amount of measuring time is required, and accumulation of technical know-how This becomes indispensable.

The present inventors have made this clearer and searched for evaluation indexes that are simpler and more accurate as a result of the search. As a result, it is possible to evaluate the cold-rolled steel composition 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 present inventors have found that when the ratio of the number of carbides on the ferrite grain boundaries to the number of carbides in the grains exceeds 1, the cold step composition remarkably improves.

In addition, buckling, bending, and folding of the steel sheet occurring during cold forging are all caused by localization of the deformation accompanying formation of the shear band. Likewise, localization of the formation and deformation of the shear band is made possible by the presence of the carbide in the ferrite grain boundaries When it is relaxed, occurrence of buckling, bending, and folding can be suppressed.

The observation of the carbide is performed by a scanning electron microscope. Prior to observation, the sample for tissue observation was polished by wet grinding with emery paper and diamond abrasive grains having an average particle size of 1 mu m, and the observation surface was finished with a mirror surface, and then the tissue was treated with a saturated picric acid- Etched.

The magnification of observation is set to 3000 times, and the field of view of 30 占 퐉 x 40 占 퐉 at the 1/4 sheet thickness is randomly photographed. With respect to the obtained tissue image, the area of each carbide contained in the area is measured in detail by image analysis software represented by Win ROOF (manufactured by Mitani Shoji Co., Ltd.). (= 2 x? (Area / 3.14)] is determined from the area of each carbide, and the average value is determined as the carbide particle diameter.

Further, in order to suppress the influence of the measurement error due to the noise, the carbide having an area of 0.01 탆 ² or less is excluded from the evaluation object.

The number of carbides present on the ferrite grain boundary phase is counted and the number of carbides on the grain boundaries is subtracted from the total number of carbides to obtain the number of carbides in the ferrite grains. The ratio of the number of carbides in the grain boundaries to the carbides in the ferrite grains is calculated on the basis of the measured number.

By setting the ferrite grain size to 3.0 占 퐉 or more and 50.0 占 퐉 or less as the structure after annealing, it is possible to improve the cold hardening. If the ferrite grain size is less than 3 占 퐉, the hardness increases and cracks and cracks easily occur during cold forging. Therefore, the ferrite grain size is preferably 3.0 占 퐉 or more. More preferably at least 7.5 mu m.

On the other hand, when the ferrite grain size exceeds 50.0 占 퐉, the number of carbides on the grain boundaries for suppressing slip propagation decreases and the cold-rolled steel composition is lowered, so that the ferrite grain size is preferably 50.0 占 퐉 or less. More preferably, it is 37.9 占 퐉 or less.

The ferrite particle size was measured in the above-described procedure. The observation surface of the sample for tissue observation was polished to a mirror surface, and the structure of the observation surface etched with a 3% nitric acid-alcohol solution was observed with an optical microscope or a scanning electron microscope, A line segment method is used to measure one image.

By setting the Vickers hardness of the steel sheet to 100 HV or more and 180 HV or less, it is possible to improve the impact resistance characteristics after cold step forming and carburizing surface tempering. If the Vickers hardness is less than 100 HV, buckling tends to occur during cold forging, and the buckling portion is bent and folded to reduce the impact resistance characteristics. Therefore, the Vickers hardness should be 100 HV or more. Preferably 110 HV or more.

On the other hand, when the Vickers hardness exceeds 180 HV, the ductility is lowered, internal cracks tend to occur during cold forging, and the impact resistance characteristic deteriorates, so that the Vickers hardness is made 180 HV or lower. And preferably 170 HV or less.

Next, a method of evaluating the cold hardening will be described.

Fig. 1 schematically shows the outline of the cold forging test and the type of crack introduced in cold forging. Fig. 1 (a) shows a test piece cut from a hot-rolled steel sheet, Fig. 1 (b) shows a test piece after cold forging, Fig. 1 (c) / RTI >

As shown in Fig. 1, a disc test piece 1 having a diameter of 70 mm was cut out from a hot-rolled steel sheet having a thickness of 5.2 mm (see Fig. 1 (a)), (Not shown). Subsequently, the vertical wall portion of the cup-shaped test material was subjected to thickness increase molding (cold forging) at a thickness increase ratio of 1.54 (= 8 mm / 5.2 mm) using a one-shot forming press machine of Moriteko Co., Thereby producing a cup-shaped test material 2 having a thickness of 8 mm (see Fig. 1 (b)).

The cup-shaped test material 2 subjected to the thickness increasing molding is cut so that the cross section of the diameter portion appears in the FANUC wire-cut electric discharge machine (see Fig. 1 (c)). The ratio of the maximum length L of the cracks existing in the vertical wall portion to the thickness of the vertical wall portion after the increase in thickness (= maximum length of the crack L / Thickness of the vertical wall 8 mm) is measured. By this measurement value, the cold forging is evaluated.

Further, even if the initial plate thickness is other than 5.2 mm, if the diameter of the circular test piece to be cut is adjusted so that the height of the end wall after the increase in thickness becomes 30 mm, and molding is performed at the same thickness increase ratio of 1.54, And the evaluation result can be reproduced. Therefore, the hot-rolled steel sheet to which the present invention is applied is not limited to the hot-rolled steel sheet having the sheet thickness of 5.2 mm. INDUSTRIAL APPLICABILITY The present invention can improve the impact resistance characteristics after cold forging and carburizing surface tempering even in a hot-rolled steel sheet having a general thickness (2 to 15 mm).

Next, the manufacturing method of the present invention will be described. The technical idea of the manufacturing method of the present invention is to uniformly manage the hot rolling condition and the annealing condition at the time of producing the steel sheet from the above-mentioned steel sheet having the component composition and to improve the cold forging property and the impact resistance characteristic after carburizing surface tempering.

The characteristics of the manufacturing method of the present invention will be described.

[Characteristics of hot rolled steel]

And the slab is continuously supplied to hot rolling as in the conventional method or is once heated and then subjected to hot rolling so as to be heated to a temperature of 650 DEG C or higher and 950 DEG C or lower Finish the hot rolling in the area. The hot-rolled steel sheet after the finish rolling is cooled on the ROT and wound at a coiling temperature of 400 DEG C or more and 600 DEG C or less.

[Characteristics of annealing]

After the pickling, the hot-rolled steel sheet is subjected to a two-step annealing process in which the hot-rolled steel sheet is held in two temperature zones. At this time, Annealing is carried out at a heating rate so as to be maintained for at least 3 hours but not more than 60 hours in a temperature range of 650 占 폚 to 720 占 폚.

In the second annealing step, the hot-rolled steel sheet is heated to an annealing temperature at a heating rate of 1 deg. C / hr or more and 80 deg. C / hr or less and maintained in a temperature range of 725 DEG C to 790 DEG C for 3 hours to 50 hours Lt; / RTI >

Then, the hot-rolled steel sheet after annealing is cooled to 650 DEG C at a cooling rate of 1 DEG C / hour or more and 100 DEG C / hour or less, and then cooled to room temperature.

By the cooperation of the hot rolling condition and the annealing condition, it is possible to obtain a low carbon steel sheet having excellent impact resistance characteristics after cold forging and carburizing surface tempering.

Hereinafter, the process conditions of the production method of the present invention will be described in detail.

[Hot Rolling]

Finishing hot rolling temperature: 650 ℃ or more and 950 ℃ or less

Coiling temperature: 400 占 폚 to 600 占 폚

The hot-rolled steel sheet is subjected to hot rolling at a temperature of not lower than 400 占 폚 and not higher than 400 占 폚 in a temperature range of 650 占 폚 to 950 占 폚, Wind at 600 ° C or less.

The slab heating temperature is preferably 1300 DEG C or lower, and the heating time in which the temperature of the slab surface layer is maintained at 1000 DEG C or higher is preferably 7 hours or less.

When the heating temperature exceeds 1300 占 폚 or the heating time exceeds 7 hours, the decarburization of the slab surface layer becomes significant, and the austenite grains in the surface layer grow abnormally during heating before the casting, The heating temperature is preferably 1300 占 폚 or less, and the heating time is preferably 7 hours or less. More preferably, the heating temperature is 1280 占 폚 or less and the heating time is 6 hours or less.

The finished hot rolled steel is finished at a temperature of 650 ° C or more and 950 ° C or less. If the finish hot rolling temperature is less than 650 占 폚, the finish rolling hot-rolled temperature is set to 650 占 폚 or more because the rolling load is remarkably increased from the increase of the deformation resistance of the steel material, the roll wear amount is increased and the productivity is lowered. Preferably 680 DEG C or more.

On the other hand, when the finish hot rolling temperature exceeds 950 占 폚, a thick scale is generated during passage of the ROT (Run Out Table), scratches are generated on the surface of the steel sheet due to the scale, and during cold forging and / When the impact load is applied after the tempering, cracks are generated starting from scratches and the impact resistance characteristic is lowered. Therefore, the finish hot rolling temperature is set to 950 占 폚 or less. Preferably 920 占 폚 or lower.

The cooling rate at the time of cooling the hot-rolled steel sheet on the ROT is preferably 10 ° C / sec to 100 ° C / sec. If the cooling rate is less than 10 占 폚 / sec, generation of a thick scale and generation of scratches caused therefrom can not be suppressed during cooling and the impact resistance characteristic is lowered. Therefore, the cooling rate is preferably 10 占 폚 / sec or more . More preferably not less than 20 ° C / second.

On the other hand, when the hot-rolled 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 do.

It is difficult to remove cracks in the subsequent pickling process and the cold rolling process due to generation of microcracks in the low temperature transformed structure at the time of dispensing the hot-rolled steel sheet at a temperature of 100 ° C to room temperature after the coiling, When an impact load is applied after tempering, the crack progresses starting from cracks, and the impact resistance characteristic is lowered. Therefore, the cooling rate is preferably 100 占 폚 / sec or less. More preferably not higher than 80 占 폚 / sec.

In addition, the cooling rate is set so that, at a point of time when the hot-rolled steel sheet after finishing hot rolling passes through the waterless section and then cooled on the ROT from the time when water is cooled in the main section to the target temperature of winding, 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 being wound by the winder.

The coiling temperature is 400 ° C or higher and 600 ° C or lower. This is a temperature lower than the general coiling temperature. 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 lower than 400 占 폚, the austenite which has been transformed to the unmodified state before the coiling is transformed into hard martensite, and the cohesion occurs in the surface layer when the coagulated hot- Or more. Preferably 430 ° C or higher.

On the other hand, when the coiling temperature exceeds 600 ° C, pearlite having a large lamellar spacing is produced, and a thick needle-shaped carbide having high thermal stability is formed. Even after the annealing in the two-step step type, the carbide of the needle remains. During cold forging, cracks are generated starting from the carbide of this needle bed as a starting point, so that the winding temperature is set to 600 캜 or less. Preferably 570 DEG C or less.

The hot-rolled steel sheet produced under the above conditions is subjected to annealing in a two-step step type that is held in two temperature zones after pickling. By performing the two-step step annealing on the hot-rolled steel sheet, the stability of the carbide is controlled and the formation of the carbide in the ferrite grain boundary is promoted.

First, the technical idea of the two-step step annealing will be described.

The first-stage annealing is performed in a temperature range of Ac1 points or less to coarsen the carbides and to thicken the additional metal elements to increase the thermal stability of the carbides. Thereafter, the temperature is raised to the temperature range of Ac1 or higher, and austenite is produced in the structure, the carbide in the fine ferrite grains is dissolved in the austenite, and the coarse carbide remains in the austenite.

Subsequently, by slow cooling, the austenite is transformed into ferrite and the carbon concentration in the austenite is increased. By progressing slowly cooling, it becomes possible to form a structure in which carbon atoms are adsorbed on the carbide remaining in the austenite, carbides and austenite cover the grain boundaries of the ferrite, and finally, a large amount of carbides are present in the ferrite grain boundaries. As a result, it is apparent that the structure defined in the present invention can not be formed by simple annealing.

Hereinafter, specific annealing conditions will be described.

[1st-stage annealing]

Heating rate to annealing temperature: 30 占 폚 / hour or more 150 占 폚 / hour

Annealing temperature: 650 ° C or more and 720 ° C or less

Holding time at annealing temperature: 3 hours or more and 60 hours or less

The heating rate to the first-stage annealing temperature is 30 占 폚 / hour or more and 150 占 폚 / hour or less. If the heating rate is less than 30 캜 / hour, it takes time to raise the temperature and the productivity is lowered. Therefore, the heating rate should be 30 캜 / hour or more. Preferably 40 DEG C / hour or more.

On the other hand, if the heating rate exceeds 150 DEG C / hour, the temperature difference between the outer peripheral portion and the inner portion of the coil increases, friction rubbing or seizing occurs due to the difference in thermal expansion, and irregularities are formed on the surface of the steel sheet. During cold forging, cracks are generated starting from the unevenness, resulting in a decrease in cold-rolled steel composition and a decrease in impact resistance characteristics after carburizing and tempering. Therefore, the heating rate is set to 150 ° C / hour or less. And preferably not more than 120 ° C / hour.

The annealing temperature (first-stage annealing temperature) in the first-stage annealing is set to be 650 ° C or higher and 720 ° C or lower. If the annealing temperature in the first stage is less than 650 deg. C, the stability of the carbide is insufficient, and it becomes difficult to retain the carbide in the austenite in the second stage annealing, so the annealing temperature in the first stage is set to 650 deg. Preferably 670 DEG C or more.

On the other hand, if the annealing temperature exceeds 720 占 폚, the austenite is produced before the stability of the carbide becomes high, and the above-described structure change can not be controlled, so that the annealing temperature is 720 占 폚 or less. Preferably 700 DEG C or less.

The holding time (first holding time) in the first-stage annealing is 3 hours or longer and 60 hours or shorter. If the holding time at the first stage is less than 3 hours, stabilization of the carbide is not sufficient and it is difficult to keep the carbide in the second-stage annealing. Therefore, the first-stage holding time is set to 3 hours or more. Preferably 10 hours or more.

On the other hand, if the holding time of the first stage exceeds 60 hours, the stability of the carbide of the first layer can not be expected to be improved and the productivity is lowered. Therefore, the holding time of the first stage is set to 60 hours or less. Preferably 50 hours or less.

[2nd-stage annealing]

Heating rate to the annealing temperature: 1 DEG C / hour or more 80 DEG C / hour

Annealing temperature: 725 ℃ or more and 790 ℃ or less

Holding time at annealing temperature: 3 hours or more and 50 hours or less

After completion of the holding in the first-stage annealing, the hot-rolled steel sheet is heated to an annealing temperature at a heating rate of 1 占 폚 / hour or more and 80 占 폚 / hour or less. In the case of cooling without performing the second-stage annealing, the ferrite grain size is not increased and an ideal structure can not be obtained.

In the second-stage annealing, austenite is generated from the ferrite grain boundaries to grow. By slowing the heating rate, nucleation of austenite can be suppressed, and it is possible to increase the grain boundary coverage of carbide in a structure obtained after slow cooling. Therefore, it is preferable that the heating rate in the second-stage annealing is small.

If the heating rate is less than 1 캜 / hour, it takes time to raise the temperature and the productivity is lowered. Therefore, the heating rate is set to 1 캜 / hour or more. Preferably 10 ° C / hour or more.

On the other hand, if the heating rate exceeds 80 DEG C / hour, the temperature difference between the outer peripheral portion and the inner portion of the coil increases, and friction shaking or seizing occurs due to a large thermal expansion difference due to the transformation, and irregularities are formed on the surface of the steel sheet. At the time of cold forging, cracks are generated starting from the unevenness, resulting in a decrease in cold-rolled steel composition and a decrease in impact resistance characteristics after carburizing and tempering, so that the heating rate is 80 ° C / hour or less.

The annealing temperature (second-stage annealing temperature) in the second-stage annealing is 725 ° C or higher and 790 ° C or lower. If the annealing temperature of the second stage is less than 725 DEG C, the amount of austenite to be produced is reduced, and the number of carbides on the ferrite grain boundary decreases and the ferrite grain size becomes smaller after the second stage 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, if the annealing temperature at the second stage exceeds 790 DEG C, it becomes difficult to keep the carbide in the austenite, and it becomes difficult to control by the above-described structure change. Therefore, the annealing temperature at the second stage is set to 790 DEG C or less. Preferably 780 DEG C or less.

The holding time (second holding time) in the second-stage annealing is 1 hour to 50 hours. If the holding time in the second stage is less than 1 hour, the amount of the austenite to be produced is small and the solubility of the carbides in the ferrite particles is insufficient, making it difficult to increase the number ratio of carbides on the ferrite grain boundaries, Therefore, the second-stage holding time is set to one hour or more. Preferably 5 hours or more.

On the other hand, if the holding time in the second stage exceeds 50 hours, it becomes difficult to keep the carbide in the austenite, so the holding time in the second stage is set to 50 hours or less. Preferably 45 hours or less.

[Cooling after annealing]

Cooling stop temperature: 650 ℃

Cooling rate: 1 ° C / hour or more 100 ° C / hour or less

After the holding in the second-stage annealing is completed, the hot-rolled steel sheet after annealing is cooled to 650 DEG C at a cooling rate of 1 DEG C / hour to 100 DEG C / hour or less. When the quenching temperature of the quenching exceeds 650 ° C, the unconverted austenite is transformed into perlite or bainite by the cooling rate exceeding 100 ° C / hour to the room temperature thereafter to increase the hardness, Since the composition is lowered, the cooling stop temperature is set to 650 ° C.

In order to cool the austenite produced in the second stage annealing to transform it into ferrite and adsorb the carbon with the carbide remaining in the austenite, the cooling rate is preferably slow. If the cooling rate 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 100 ° C / hour, the austenite is transformed into pearlite, the hardness of the steel sheet is increased, the cold step composition is lowered and the impact resistance characteristic is lowered after carburizing surface tempering, Is set at 100 DEG C / hour or less. Preferably 90 [deg.] C / hour.

Here, the cooling stop temperature is a temperature to be controlled by the above-mentioned cooling rate. When cooling to 650 DEG C is performed at a cooling rate of 1 DEG C / hour or more and 100 DEG C / hour or less, It does not.

Further, the atmosphere of the annealing is not limited to a specific atmosphere. For example, the atmosphere may be 95% or more of nitrogen atmosphere, 95% or more of hydrogen atmosphere, or air 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 structure of the steel sheet, it is possible to exhibit excellent cold-rolled steel for cold forging combined with drawing and thickness- It is possible to produce a low carbon steel sheet excellent in impact resistance characteristics after carburization and tempering.

Example

Next, the embodiment will be described, but the level of the embodiment is an example of the execution condition adopted to confirm the feasibility and effect of the present invention, and the present invention is not limited to this example condition. 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.

Continuous casting pieces (ingot) having the composition shown in Table 1 were heated at 1240 占 폚 for 1.8 hours and then subjected to hot rolling. The hot rolling was finished at 890 占 폚, cooled to 520 占 폚 at a cooling rate of 45 占 폚 / sec on the ROT, and rolled at 510 占 폚 to produce a hot-rolled coil having a thickness of 5.2 mm.

The hot-rolled coil was pickled, the coil was charged into the box-shaped annealing furnace, the atmosphere was controlled to 95% hydrogen-5% nitrogen, and then the furnace was heated from room temperature to 705 ° C at a heating rate of 100 ° C / The temperature distribution in the coil was made uniform by keeping it for 36 hours. Thereafter, the mixture was heated to 760 占 폚 at a heating rate of 5 占 폚 / hour, maintained at 760 占 폚 for 10 hours, cooled to 650 占 폚 at a cooling rate of 10 占 폚 / hour, A sample for evaluation was prepared.

The texture of the sample was observed by the method described above, and the crack length present in the sample after cold forging was measured by the method described above.

Carburization of the molded sample was carried out by gas carburization. In order to diffuse carbon into the steel through the sample surface layer from the atmospheric gas in the furnace, treatment was carried out at 940 占 폚 for 120 minutes in a furnace where the carbon potential was controlled to 0.5 mass% C, did.

Subsequently, the mixture was heated from room temperature to 840 占 폚, followed by holding for 20 minutes, and the flask was immersed in an oil at 60 占 폚. The sample was kept at 170 占 폚 for 60 minutes and then subjected to air cooling so as to prepare a carburizing quenched tempering sample.

The impact resistance of the carburizing quenched samples was evaluated by a drop weight test. Fig. 2 schematically shows the outline of the drop test in which the impact resistance characteristic of a sample subjected to carburizing treatment is evaluated. The bottom of the cup of the cup-shaped sample 4 subjected to the carburizing quenching was fixed with a jig and a weight of 2 kg of weight was dropped on the cup side (phase width: 50 mm, lower width: 10 mm, height: 110 mm) was allowed to fall freely from an upper portion spaced apart by 4 m, impact of about 80 J was applied to the vertical wall of the sample 4, and the presence or absence of cracks in the sample was observed to evaluate the impact resistance characteristics.

As a result of free fall, samples with no cracks or fractures were rated "OK" with excellent impact resistance, and samples with cracks or fracture were rated "NG" with poor impact resistance.

Table 2 shows the ratio of the carbide diameter, the pearlite area ratio, the ferrite grain size, the Vickers hardness, the ratio of the number of carbides in the ferrite grains to the number of carbides in the ferrite grains, the maximum crack length And the results of the measurement of the impact resistance and the evaluation results.

As shown in Table 2, the inventive strengths A-1, B-1, C-1, D-1, E-1, F-1, G-1, H-1, I- -1 all show a ratio of the number of carbides on the ferrite grain boundaries to the number of carbides in the ferrite grains exceeding 1 and a Vickers hardness of 100 HV or more and 180 HV or less and excellent impact resistance after cold forging and carburizing .

On the other hand, the comparative steel L-1 has a low C content because the amount of C is low and the hardness before cold forging is less than 100 HV. Since the comparative steels M-1, P-1 and Z-1 contain excess P, Al and N and the amount of segregation at the γ / α interface is large during the second stage annealing, Is suppressed.

Comparative steel S-1 contains excess Si and has low ductility due to low ductility of ferrite. Since the comparative steels N-1 and T-1 each contain excess Mo and Cr, the carbides are finely dispersed in the ferrite grains and the hardness exceeds 180 HV. The comparative steel Q-1 contains Mn excessively, so that the impact resistance after the carburization quenching is remarkably low.

The comparative steel O-1 has a small amount of Cr, and since the austenite grains in the surface layer are abnormally coarsened at carburizing, the impact resistance is low. Since the comparative steel R-1 contains excess S, coarse MnS is produced in the steel and the cold steel composition is low. Since comparative steel U-1 contains excess C, coarse carbide is produced inside the steel thickness increase, and coarse carbide remains after carburizing, so that the impact resistance characteristic is low.

The comparative steel V-1 had a low amount of Mn and it was difficult to increase the stability of the carbide, so that the impact resistance characteristics after the cold-forging and carburizing were tempered. Since the comparative steels W-1 and X-1 contain excessive amounts of O and Ti, the oxides and TiC present in the ferrite grains become carbide formation sites in the slow cooling after the two-phase region annealing, Is suppressed, and the cold step composition is low. The comparative steel Y-1 contains B in excess, so the cold-rolled steel is low.

Subsequently, slabs having component compositions of A, B, C, D, E, F, G, H, I, J and K shown in Table 1 were subjected to hot rolling And annealing conditions, a hot-rolled sheet annealing sample having a thickness of 5.2 mm was produced.

Table 4 shows the ratio of the carbide diameter, the pearlite area ratio, the ferrite grain size, the Vickers hardness, the ratio of the number of carbides in the ferrite grain boundaries to the number of carbides in the ferrite grains, and the maximum crack length Ratio and impact resistance, and evaluation results.

The comparative steel E-3 has low finish hot rolling temperature and low productivity due to increased rolling load. The comparative steel D-2 had a high hot finish temperature and a scale trace on the surface of the steel sheet. Therefore, cracks and peeling occurred on the basis of scale marks when they were provided to the abrasion resistance test after pattern tempering, . The comparative steel F-2 has a slow cooling rate in the ROT (Run Out Table), resulting in a decrease in productivity and generation of scales.

The comparative steel C-4 had a cooling rate of 100 DEG C / sec in the ROT, and the outermost surface of the steel sheet was excessively cooled, resulting in fine cracks in the outermost surface layer. The comparative steel C-2 has a low coiling temperature, a large number of low-temperature transformed structures such as bainite and martensite are generated and embrittled, and cracks are frequently generated when the hot-rolled coil is discharged, and productivity is lowered. In addition, the abrasion resistance of the sample taken from the cracked piece is low.

The comparative steel G-2 has a high coiling temperature, a high pearlite having a large lamellar spacing in the hot-rolled structure, a high thermal stability of the coarse carbide in the acicular phase, and even after the annealing in two- So that the workability is low. The comparative steel H-4 has a low productivity due to the slow heating rate in the first stage annealing of the two-step step annealing.

The comparative steel E-3 has a high heating rate in the first-stage annealing, so that the temperature difference between the inside of the coil and the inner and outer peripheral portions of the coil becomes large, friction rubbing and seizing due to the thermal expansion difference occur, When it was provided to the evaluation test of the characteristics, cracking and peeling occurred from the absorbent portion, and the abrasion resistance property was deteriorated.

The comparative steel G-4 had a low holding temperature (annealing temperature) in the first-stage annealing, insufficient coarsening treatment of the carbide at the Ac1 point or less, and insufficient thermal stability of the carbide, , The remaining carbide is decreased, and the pearlite transformation can not be suppressed in the post-annealed structure. Therefore, the machinability is low.

The comparative steel D-4 has a high holding temperature (annealing temperature) in the first-stage annealing, austenite is generated during annealing, and the stability of the carbide can not be increased. Therefore, pearlite is generated after annealing, , And the workability is low. The comparative steel J-4 has a short holding time in the first-stage annealing, can not increase the stability of the carbide, and has low workability.

The comparative steel F-2 has a long holding time in the first-stage annealing and a low productivity, in addition to the generation of shrinkage and low abrasion resistance. The comparative steel B-4 has a low productivity due to a slow heating rate in the second-stage annealing of the two-step step annealing. In the comparative steel A-3, since the heating speed in the second-stage annealing is fast, the temperature difference between the inside and the outer circumferential portion of the coil becomes large, and friction and seizing due to a large difference in thermal expansion due to transformation occur. The abrasion resistance property after polishing is low.

The comparative steel K-2 has a low holding temperature (annealing temperature) in the second-stage annealing, a small amount of austenite to be formed, and a low number of carbides in the ferrite grain boundaries. In comparative steel C-4, since the holding temperature (annealing temperature) in the second-stage annealing was high and the dissolution of the carbide during annealing was accelerated, it was difficult to form the intergranular carbide after gradual cooling and pearlite was generated, Hardness exceeds 180 HV, and the workability is low.

The comparative steel J-3 has a low workability because the holding time in the second-stage annealing is long and the dissolution of carbide is accelerated. The comparative steel D-3 had a slow cooling rate from 650 ° C to the second-stage annealing, a low productivity, and a coarse carbide was formed in the structure after the slow cooling. In the cold forging, And the cold-rolled composition was lowered. The comparative steel I-3 has a low cooling rate because the cooling rate from the second-stage annealing to 650 ° C is fast and the pearlite transformation occurs during cooling to increase the hardness.

Next, in order to investigate the allowable content of other elements, continuous casting pieces (ingot) having the composition shown in Tables 5 and 6 (following Table 5) were heated at 1240 占 폚 for 1.8 hours and then subjected to hot rolling . The hot rolling was finished at 890 占 폚, cooled to 520 占 폚 at a cooling rate of 45 占 폚 / sec on the ROT, and rolled at 510 占 폚 to produce a hot-rolled coil having a thickness of 5.2 mm.

The hot-rolled coil was pickled, the hot-rolled coil was charged into the box-shaped annealing furnace, the atmosphere was controlled to 95% hydrogen-5% nitrogen, and then the furnace was heated from room temperature to 705 ° C at a heating rate of 100 ° C / For 36 hours to uniformize the temperature distribution in the coil and thereafter heated to 760 占 폚 at a heating rate of 5 占 폚 / hour, maintained at 760 占 폚 for 10 hours, cooled to 650 占 폚 at 10 占 폚 / , And then cooled to room temperature to prepare a sample for evaluation of properties.

Further, the texture of the sample was observed by the above-mentioned method, and the crack length existing in the sample after cold forging was measured by the above-mentioned method.

Table 7 shows the ratio of the carbide diameter, the pearlite area ratio, the ferrite grain size, the Vickers hardness, the ratio of the number of carbides in the ferrite grains to the number of carbides in the ferrite grains, the maximum crack length And the results of the measurement of the impact resistance and the evaluation results.

As shown in Table 7, the inventive steels AA-1, AB-1, AC-1, AD-1, AE-1, AF-1, AG-1, AH-1, AI- 1, AL-1, AM-1, AN-1, AO-1, AP-1 and AQ-1 all have a ratio of the number of carbides on the ferrite grains to the number of carbides in the ferrite grains exceeds 1 , Vickers hardness of 100 HV or more and 180 HV or less, and excellent impact resistance after cold forging and carburizing surface tempering.

In comparison, the comparative steels AR-1, AS-1, AW-1, AZ-1, BB-1 and BF-1 contain excess La, As, Cu, Ni, Sb and Ce, The amount of segregation at the? /? Interface increases during the annealing of the grain boundary, so generation of carbide in the grain boundary is suppressed. The comparative steel BG-1 contains excess Si and has low ductility due to low ductility of ferrite.

Since the comparative steels AT-1, AV-1, BA-1, BC-1, BH-1 and BJ-1 contain excess Mo, Nb, Cr, Ta, W and V respectively, Is finely dispersed and the hardness exceeds 180 HV. Since the comparative steel BF-1 contains Mn in excess, the impact resistance after the carburization hardening is remarkably low.

The comparative steels AU-1, AX-1, AY-1 and BE-1 each contain excess Zr, Ca, Mg and Y and coarse oxides or nonmetallic inclusions are generated in the steel, Cracks were generated starting from oxides or coarse nonmetallic inclusions, and the cold step composition was lowered. The comparative steel BD-1 contains excess Sn, ferrite is brittle, and low cold-rolled steel is low. Since comparative steel BK-1 contains excess C, coarse carbides are produced in the steel thickness increase, and coarse carbides remain even after carburizing, and the impact resistance characteristics are deteriorated.

Next, in order to investigate the influence of the manufacturing conditions, the values of AA, AB, AC, AD, AE, AF, AG, AH, AI, AJ, AK, AL, AM, A slab having a component composition was subjected to a hot-rolled sheet annealing sample having a thickness of 5.2 mm under the hot rolling condition and the annealing condition shown in Table 8.

Table 9 shows the ratio of the carbide diameter, the pearlite area ratio, the ferrite grain size, the Vickers hardness, the ratio of the number of carbides in the ferrite grains to the number of carbides in the ferrite grains, the maximum crack length And the results of the measurement of the impact resistance and the evaluation results.

The comparative steel AC-2 has a low hot finish temperature and low productivity. The comparative steel AN-4 had high finish hot rolling temperature, scale marks were formed on the surface of the steel sheet, cracks were generated from the scratches when impact load was applied after cold forging and carburizing corrosion tempering, .

Inventive steel AB-3 has a slow cooling rate in the ROT, resulting in a decrease in productivity and the derivation of scales. Inventive steels AJ-3 and AD-4 had a cooling rate of 100 ° C / sec in the ROT, and the outermost surface of the steel sheet was excessively cooled, resulting in a fine crack at the outermost surface.

The comparative steel AN-3 had a low coiling temperature, a large number of low-temperature transformed textures such as bainite and martensite were generated and became brittle, cracks occurred frequently when the hot-rolled coil was discharged, and the productivity deteriorated. In addition, the impact resistance characteristics after cold forging and carburizing shrinkage tempering in samples taken from the crack pieces were inferior.

The comparative steel AH-3 has a high coiling temperature, a high pearlite with a large lamellar spacing in the hot-rolled structure, a high thermal stability of the acicular carbide in the acicular form, and even after the annealing in two- And therefore, the cold-rolled composition is low.

The comparative steel AF-4 has a low productivity due to the slow heating rate in the first-stage annealing in the two-step step annealing. Since the comparison steel AG-2 has a high heating rate in the first-stage annealing, the temperature difference between the inside and the outer circumferential portion of the coil becomes large, friction rubbing and seizing due to the difference in thermal expansion occur, and cold forging and carburizing refinement tempering The impact resistance characteristics after the test were lowered.

The comparative steel AA-2 had a low holding temperature (annealing temperature) in the first-stage annealing, insufficient coarsening treatment of the carbide at the Ac1 point or less, insufficient thermal stability of the carbide, The remaining carbide was decreased, pearlite transformation was not suppressed in the post-cooling structure, and the cold step composition was lowered.

The comparative steel AM-3 had a high holding temperature (annealing temperature) at the first stage, austenite was generated during annealing, the stability of the carbide could not be enhanced, and the impact resistance characteristics after cold-tempering and carburizing were tempered. The comparative steel AF-2 has a short holding time in the first-stage annealing, can not increase the stability of the carbide, and has a low cold-rolled composition. The comparative steel AO-4 has a long holding time in the first-stage annealing and low productivity.

The comparative steel AP-4 has a low productivity due to the slow heating rate in the second-stage annealing of the two-step step annealing. Since the comparative steel AI-3 has a high heating rate in the second-stage annealing, the temperature difference between the inside and the outer circumferential portion of the coil becomes large, friction rubbing and seizing due to a large thermal expansion difference due to the transformation occur, When an impact load was applied later, cracks were generated from the corresponding scratches, and the impact resistance characteristics deteriorated.

The comparative steel AL-3 had a low holding temperature (annealing temperature) in the second-stage annealing, a small amount of austenite, no increase in the number of carbides in the ferrite grain boundaries, . The comparative steel AD-2 had a high holding temperature (annealing temperature) in the second-stage annealing and promoted dissolution of the carbide during the annealing, making it difficult to generate grain boundary carbide after gradual cooling, The impact resistance characteristic after quenching was deteriorated.

The comparative steel AJ-4 has a low cold-rolled steel composition because the holding time in the second-stage annealing is long and the melting of carbide is promoted. The comparative steel AQ-3 had a slow cooling rate from 650 ° C to the second-stage annealing, a low productivity, and a coarse carbide was formed in the post-cooling structure. During the cold forging, Resulting in a decrease in cold-rolled steel composition. The comparative steel AP-2 had a rapid cooling rate from the second stage annealing to 650 DEG C, pearlite transformation occurred at the time of cooling, and hardness increased, so that the cold step composition was lowered.

Here, FIG. 3 shows the relationship between the ratio of the number of intergranular carbides to the number of carbides in the particle, the crack length of the cold forging test piece, and the impact resistance characteristics after carburizing and tempering.

3, when the number ratio (= number of intergranular carbides / number of carbides in particles) exceeds 1, the ratio of the length of cracks introduced in cold forging can be suppressed, and excellent impact resistance after carburization is tempered Is obtained.

Fig. 4 shows another relationship between the ratio of the number of intergranular carbides to the number of carbides in the particle, the crack length of the cold forging test piece, and the impact resistance characteristic after carburizing surface tempering. 4 is a diagram showing that crack length can be suppressed even in a steel sheet to which an additive element is added.

4, even when the steel sheet is added with an appropriate range of elements, when the number ratio (= number of intergranular carbides / number of carbides in particles) exceeds 1, the ratio of the length of cracks introduced in cold forging is suppressed And excellent impact resistance can be obtained after carburizing quenching.

INDUSTRIAL APPLICABILITY As described above, according to the present invention, it is possible to provide a low carbon steel sheet excellent in impact resistance characteristics after cold forging and carburizing and tempering, and a method for producing the same. INDUSTRIAL APPLICABILITY The steel sheet of the present invention is suitable as a material for obtaining parts such as high cycle gears, for example, by cold forging such as plate forming.

1: Round test specimen
2: cup-shaped test piece
3: Crack
4: Samples
5: Reducing
L: Maximum length of crack

Claims (2)

The composition of matter, in% by mass,
C: 0.10 to 0.40%,
Si: 0.01 to 0.30%
Mn: 0.30 to 1.00%
Al: 0.001 to 0.10%
0.50 to 2.00% Cr,
Mo: 0.001 to 1.00%
P: 0.020% or less,
S: 0.010% or less,
N: 0.020% or less,
O: 0.020% or less,
Ti: 0.010% or less,
B: 0.0005% or less,
Sn: 0.050% or less,
Sb: 0.050% or less,
As: 0.050% 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,
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 and
Ce: 0.050%
And the remainder being Fe and impurities,
Wherein the metal structure of the low-
A carbide particle diameter of 0.4 to 2.0 mu m,
If the pearlite area ratio is 6% or less and
When the ratio of the number of carbides on the ferrite grains to the number of carbides in the ferrite grains exceeds 1
Lt; / RTI >
Wherein the low-carbon steel sheet has a Vickers hardness of 100 HV or more and 180 HV or less
. A manufacturing method for manufacturing the steel sheet according to claim 1,
A hot-rolled steel sheet is obtained by subjecting a steel strip having the composition described in claim 1 to hot rolling to finish hot rolling in a temperature range of 650 占 폚 to 950 占 폚,
The hot-rolled steel sheet is rolled at a temperature of 400 ° C or higher and 600 ° C or lower,
The picked hot-rolled steel sheet is pickled, and the pickled hot-rolled steel sheet is heated at an annealing temperature of 650 ° C to 720 ° C at a heating rate of 30 ° C / hour or more and 150 ° C / hour or less, The first-stage annealing is performed, and subsequently,
The hot-rolled steel sheet is heated at an annealing temperature of not less than 725 DEG C and not more than 790 DEG C at a heating rate of not less than 1 DEG C / hour and not more than 80 DEG C / hour, and is held for not less than 3 hours and not more than 50 hours, The steel sheet is cooled to 650 占 폚 at a cooling rate of 1 占 폚 / hour or more and 100 占 폚 / hour or less
Wherein the steel sheet is produced by a method comprising the steps of:

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