WO2016204288A1

WO2016204288A1 - Steel sheet and manufacturing method

Info

Publication number: WO2016204288A1
Application number: PCT/JP2016/068169
Authority: WO
Inventors: 匹田　和夫; 元仙橋本; 健悟竹田; 高田　健
Original assignee: 新日鐵住金株式会社
Priority date: 2015-06-17
Filing date: 2016-06-17
Publication date: 2016-12-22
Also published as: EP3312299A4; KR101997382B1; CN107735505B; US20180171445A1; TW201708564A; JPWO2016204288A1; CN107735505A; BR112017025756A2; JP6206601B2; EP3312299A1; MX2017015266A; TWI588270B; KR20180004262A

Abstract

A steel sheet, in which moldability and abrasion resistance have been improved, is characterized in that: the steel sheet has a prescribed component composition; the metal texture of the steel sheet satisfies the ratio of the number of ferrite grain boundary carbides to the number of carbides inside the ferrite particles being greater than 1 and the ferrite grain diameter being 5 µm to 50 µm; and the Vickers hardness of the steel sheet is 100 HV to 170 HV.

Description

Steel plate and manufacturing method

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

Automotive parts such as gears and clutches are manufactured through processing steps such as punching, forging, and pressing. In the machining process, improvement of workability of the carbon steel plate as a material is required to improve the product quality, stabilize it, and reduce the manufacturing cost. Moreover, since these parts are used with high strength after quenching and tempering, excellent hardenability is required.

Many proposals have been made to secure the workability and hardenability of carbon steel sheets.

In Patent Document 1, in mass%, C: 0.20 to 0.45%, Mn: 0.40 to 1.50%, P: 0.03% or less, S: 0.02% or less, P + S: Containing 0.010% or more, Cr: 0.01-0.80%, Ti: 0.005-0.050%, B: 0.0003-0.0050%, and the balance consisting of Fe and inevitable impurities Further, Sn: 0.05% or less, Te: 0.05% or less, and the total content of Sn + Te is 0.005% or more, and a mixed structure of ferrite and pearlite, or a mixture of ferrite and cementite A high carbon steel sheet excellent in workability, hardenability, and toughness after heat treatment characterized by comprising a structure is disclosed.

In Patent Document 2, in mass%, C: 0.2 to 0.7%, Si: 2% or less, Mn: 2% or less, P: 0.03% or less, S: 0.03% or less, sol . A steel containing Al: 0.08% or less, N: 0.01% or less, the balance being iron and inevitable impurities, after hot rolling at a finishing temperature (Ar3 transformation point −20 ° C.) or more, After cooling at a cooling rate of 120 ° C./second and a cooling end temperature of 620 ° C. or less, then winding at a coiling temperature of 600 ° C. or less, and controlling to a structure having a bainite phase exceeding 20% in volume ratio, A method for producing a high hardenability high carbon hot-rolled steel sheet characterized by annealing after pickling and annealing at an annealing temperature of 640 ° C. or higher and an Ac1 transformation point or lower to form a spheroidized structure is disclosed.

Japanese Patent No. 4319940 Japanese Patent No. 3879459

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

In view of the current state of the art, the present invention provides a steel plate suitable for improving the formability and wear resistance, particularly for obtaining parts such as gears and clutches by thick plate molding, and a method for producing the same. Objective.

In order to solve the above problems and to obtain a steel sheet suitable for a material such as a drive train component, in the steel sheet containing C necessary for improving the hardenability, the ferrite grain size is increased, and carbide (mainly cementite) is obtained. ) Can be spheroidized with an appropriate particle size to reduce the pearlite structure. This is due to the following reason.

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

By properly dispersing carbide in the metal structure, it can provide excellent wear resistance and rolling fatigue characteristics while maintaining material formability. It is. Also, carbides in the steel sheet are strong particles that prevent slipping, and by allowing carbides to exist at the ferrite grain boundaries, it is possible to prevent the propagation of slips across the crystal grain boundaries and suppress the formation of shear bands. It can improve the cold forgeability and at the same time improve the formability of the steel sheet.

However, cementite is a hard and brittle structure, and if it exists in the state of pearlite, which is a layered structure with ferrite, the steel becomes hard and brittle, so it must be present in a spherical shape. In consideration of cold forgeability and generation of cracks during forging, the particle size needs to be in an appropriate range.

However, a manufacturing method for realizing the above organization has not been disclosed so far. Therefore, the present inventors have intensively studied a manufacturing method for realizing the above structure.

As a result, the steel structure after coiling after hot rolling is made into a bainite structure in which cementite is dispersed in fine pearlite or fine ferrite with a small lamellar spacing, so that it is at a relatively low temperature (400 to 550 ° C). Take up. By winding at a relatively low temperature, cementite dispersed in the ferrite is also easily spheroidized. Subsequently, the cementite is partially spheroidized by annealing at a temperature just below the Ac1 point as the first stage annealing. Next, as the second stage annealing, annealing is performed at a temperature between Ac1 point and Ac3 point (so-called two-phase region of ferrite and austenite), and a part of the ferrite grains is left, and a part thereof is austenite transformed. Thereafter, the ferrite grains left by slow cooling were grown, and austenite was transformed into ferrite by using the ferrite grains as a nucleus, so that cementite was precipitated at the grain boundaries while obtaining a large ferrite phase, and the above structure was realized.

In other words, a steel sheet manufacturing method that satisfies both hardenability and formability is difficult to achieve even if the hot rolling conditions and annealing conditions are devised by a single method. It was found that this can be achieved by achieving optimization.

The present invention has been made on the basis of the above findings, and the gist thereof is as follows.

(1) By mass%, C: 0.10 to 0.40%, Si: 0.01 to 0.30%, Mn: 1.00 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, Cr: 0.50% or less, Mo: 0.10% or less, Nb : 0.10% or less, V: 0.10% or less, Cu: 0.10% or less, W: 0.10% or less, Ta: 0.10% or less, Ni: 0.10% or less, Sn: 0 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 %, La: 0.050% or less, Ce: 0.050% or less, the balance being Fe and unavoidable impurities steel plate, The ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grain to the number of carbides in the ferrite grains is more than 1, the ferrite grain size is 5 μm to 50 μm, and the pearlite area ratio is 6% or less. A steel plate having a hardness of 100 HV or more and 170 HV or less.

(2) The steel sheet according to (1) above, which contains one or two of Ti: 0.10% or less and B: 0.010% or less in place of a part of the Fe.

(3) A production method for producing the steel sheet of (1) or (2), wherein the steel slab having the component composition of (1) or (2) is finish-rolled in a temperature range of 750 ° C. or higher and 850 ° C. or lower. The hot-rolled steel sheet is rolled into a hot-rolled steel sheet at a temperature of 400 ° C. or higher and 550 ° C. or lower, and the wound hot-rolled steel sheet is pickled and the pickled hot-rolled steel sheet is 650 ° C. or higher. A first-stage annealing is performed in a temperature range of 720 ° C. or lower, which is maintained for 3 hours or more and 60 hours or less, and then the hot-rolled steel sheet is held in a temperature range of 725 ° C. or more and 790 ° C. or less for 3 hours or more and 50 hours or less 2 A method for producing a steel sheet, comprising subjecting a stage of annealing and cooling the hot-rolled steel sheet after annealing to 650 ° C. at a cooling rate of 1 ° C./hour or more and 30 ° C./hour or less.

According to the present invention, it is possible to provide a steel plate that is excellent in formability and wear resistance, and that is particularly suitable for obtaining parts such as gears and clutches by thick plate forming and a method for manufacturing the steel plate.

Hereinafter, the present invention will be described in detail. First, the reasons for limiting the component composition of the steel sheet of the present invention will be described. Less than. “%” For a component means “mass%”.

[C: 0.10 to 0.40%]
C is an element that forms carbides in steel and is effective in strengthening steel and refining ferrite grains. In order to suppress the occurrence of satin in cold working and ensure the surface aesthetics of the cold worked parts, it is necessary to suppress the coarsening of the ferrite grain size, but if it is less than 0.10%, the volume fraction of carbides Is insufficient, and coarsening of carbide during box annealing cannot be suppressed, so C is made 0.10% or more. Preferably it is 0.12 or more.

On the other hand, if it exceeds 0.40%, the volume fraction of carbide increases, and when a load is instantaneously applied, a large amount of cracks that become the starting point of fracture are generated, and the impact resistance characteristics are reduced. 0.40% or less. Preferably it is 0.38% or less.

[Si: 0.01-0.30%]
Si is an element that acts as a deoxidizer and affects the morphology of carbides. In order to obtain a deoxidizing effect, Si is made 0.01% or more. Preferably it is 0.05% or more.

On the other hand, if it exceeds 0.30%, the ductility of ferrite is lowered, cracking is likely to occur during cold working, and cold workability is lowered, so Si is made 0.30% or less. Preferably it is 0.28% or less.

[Mn: 1.00 to 2.00%]
Mn is an element that enhances hardenability and contributes to improvement in strength. If it is less than 1.00%, it becomes difficult to ensure the strength after quenching and the residual carbide after quenching, so Mn is made 1.00% or more. Preferably it is 1.09% or more.

On the other hand, if it exceeds 2.00%, Mn segregation becomes an extreme band shape and the workability is remarkably deteriorated, so Mn is made 2.00% or less. Preferably it is 1.91% or less.

[Al: 0.001 to 0.10%]
Al is an element that acts as a deoxidizer for steel and stabilizes ferrite. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so Al is made 0.001% or more. Preferably it is 0.004% or more.

On the other hand, if it exceeds 0.10%, a large amount of inclusions are produced and cold workability is lowered, so Al is made 0.10% or less. Preferably it is 0.08% or less.

The following elements are impurities and must be controlled to a certain amount or less.

[P: 0.0001 to 0.020%]
P is an element that segregates at the ferrite grain boundaries and suppresses the formation of grain boundary carbides. The smaller the amount, the better. However, if P is reduced to less than 0.0001% in the refining process, the refining cost increases significantly. Therefore, P is set to 0.0001% or more. Preferably it is 0.0013% or more.

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

[S: 0.0001 to 0.010%]
S is an impurity element that forms non-metallic inclusions such as MnS. Since non-metallic inclusions are the starting point for cracking during cold working, the smaller the amount of S, the better. However, if S is reduced to less than 0.0001%, the refining cost will be significantly increased, so S will be 0.2. 0001% or more. Preferably it is 0.0012% or more.

On the other hand, if it exceeds 0.010%, cold workability deteriorates, so S is made 0.010% or less. Preferably it is 0.007% or less.

[N: 0.0001 to 0.010%]
N is an element that causes embrittlement of ferrite due to the inclusion of a large amount, and the smaller the amount, the better. The N content may be 0, but if the content is reduced to less than 0.0001%, the refining cost increases significantly, so the practical lower limit is 0.0001 to 0.0006%. On the other hand, if it exceeds 0.010%, ferrite becomes brittle and cold workability deteriorates, so N is made 0.010% or less. Preferably it is 0.007% or less.

[O: 0.0001 to 0.020%]
O is an element that forms a coarse oxide in steel due to its large content, and it is preferable that O be small. The O content may be 0, but if the content is reduced to less than 0.0001%, the refining cost increases significantly, so 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 during cold working, so O is made 0.020% or less. Preferably it is 0.017% or less.

[Sn: 0.001 to 0.050%]
Sn is an element mixed from the steel raw material (scrap). Since it segregates at a grain boundary and causes a decrease in the number ratio of grain boundary carbides, the smaller the number, the better. The Sn content may be 0, but if it is reduced to less than 0.001%, the refining cost will increase significantly, so the practical lower limit is 0.001 to 0.002% or more. On the other hand, if it exceeds 0.050%, ferrite becomes brittle and cold workability deteriorates, so Sn is made 0.050% or less. Preferably it is 0.040% or less.

[Sb: 0.001 to 0.050%]
Sb is an element mixed from steel raw material (scrap) like Sn. Since it segregates at a grain boundary and causes a decrease in the number ratio of grain boundary carbides, the smaller the number, the better. The Sb content may be 0, but if the content is reduced to less than 0.001%, the refining cost increases significantly, so the substantial lower limit is 0.001 to 0.002% or more. On the other hand, if it exceeds 0.050%, ferrite becomes brittle and cold workability deteriorates, so Sb is made 0.050% or less. Preferably it is 0.040% or less.

[As: 0.001 to 0.050%]
As is an element mixed from steel raw material (scrap), as in Sn and Sb. Since it segregates at a grain boundary and causes a decrease in the number ratio of grain boundary carbides, the smaller the number, the better. The content of As may be 0, but if the content is reduced to less than 0.001%, the refining cost increases significantly, so the practical lower limit is 0.001 to 0.002% or more. On the other hand, if it exceeds 0.050%, the number ratio of grain boundary carbides decreases and cold workability deteriorates, so As is made 0.050% or less. Preferably it is 0.040% or less.

The steel sheet of the present invention contains the above elements as basic components, but may further contain the following elements for the purpose of improving the cold forgeability of the steel sheet. The following elements are not essential for obtaining the effects of the present invention, so the content may be zero.

[Cr: 0.50% or less]
Cr is an element that improves hardenability and contributes to the improvement of strength, and is an element that concentrates in carbides and forms stable carbides even in the austenite phase. In order to obtain the additive effect, Cr is preferably 0.001% or more. More preferably, it is 0.007% or more. On the other hand, if it exceeds 0.50%, the carbide is stabilized, the dissolution of the carbide is delayed at the time of quenching, and the required quenching strength may not be achieved, so Cr is 0.50% or less. Preferably it is 0.45% or less.

[Mo: 0.10% or less]
Mo, like Mn, is an element effective for controlling the morphology 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 deteriorates, so Mo is made 0.10% or less. Preferably it is 0.08% or less.

[Nb: 0.10% or less]
Nb is an element that is effective for controlling the morphology of carbides, and is an element that refines the structure and contributes to improved toughness. 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%, many fine Nb carbides precipitate, the strength excessively increases, the number ratio of grain boundary carbides decreases, and cold workability decreases, so Nb is 0 10% or less. Preferably it is 0.08% or less.

[V: 0.10% or less]
V, like Nb, is an element that is effective for controlling the morphology of carbides, and is an element that contributes to refinement of the structure and improvement of toughness. 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%, many fine V carbides precipitate, the strength increases excessively, the number ratio of grain boundary carbides decreases, and cold workability decreases, so V is 0. 10% or less. Preferably it is 0.08% or less.

[Cu: 0.10% or less]
Cu is an element that segregates at the ferrite grain boundaries and contributes to the improvement of strength by forming fine precipitates. 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%, red heat embrittlement occurs and productivity in hot rolling decreases, so Cu is made 0.10% or less. Preferably it is 0.08% or less.

[W: 0.10% or less]
W, like Nb and V, is an element effective for controlling the form 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 precipitate, the strength excessively increases, the number ratio of grain boundary carbides decreases, and the cold workability decreases, so W is 0. 10% or less. Preferably it is 0.08% or less.

[Ta: 0.10% or less]
Ta, as well as Nb, V, and W, is an element effective for controlling the morphology of carbides. In order to obtain the effect of addition, Ta is preferably 0.001% or more. More preferably, it is 0.005% or more. On the other hand, if it exceeds 0.10%, many fine W carbides precipitate, the strength increases excessively, the number ratio of grain boundary carbides decreases, and cold workability decreases, so Ta is 0. 10% or less. Preferably it is 0.08% or less.

[Ni: 0.10% or less]
Ni is an element effective for improving the toughness of parts. In order to obtain the additive effect, 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 number ratio of grain boundary carbides decreases and cold workability decreases, so Ni is made 0.10% or less. Preferably it is 0.08% or less.

[Mg: 0.050% or less]
Mg is an element that can control the form of sulfide by addition of a small 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%, ferrite becomes brittle and cold workability deteriorates, so Mg is made 0.050% or less. Preferably it is 0.040% or less.

[Ca: 0.050% or less]
Ca, like Mg, is an element that can control the form of sulfide with a small amount of addition. 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%, coarse Ca oxide is produced and becomes a starting point of cracking during cold working, so Ca is made 0.050% or less. Preferably it is 0.040% or less.

[Y: 0.050% or less]
Y, like Mg and Ca, is an element that can control the form of sulfide by addition of a trace amount. In order to obtain the effect of addition, Y is preferably 0.001% or more. More preferably, it is 0.003% or more. On the other hand, if it exceeds 0.050%, coarse Y oxide is generated and becomes a starting point of cracking during cold working, so Y is made 0.050% or less. Preferably it is 0.035% or less.

[Zr: 0.050% or less]
Zr, like Mg, Ca, and Y, is an element that can control the form of sulfide by adding a small 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, if it exceeds 0.050%, a coarse Zr oxide is generated and becomes a starting point of cracking during cold working, so Zr is made 0.050% or less. Preferably it is 0.045% or less.

[La: 0.050% or less]
La is an element that is effective for controlling the form of sulfide when added in a small amount, but is also an element that segregates at the grain boundary and causes a decrease in the number ratio of grain boundary carbides. 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 number ratio of grain boundary carbides decreases and cold workability decreases, so La is made 0.050% or less. Preferably it is 0.045% or less.

[Ce: 0.050% or less]
Ce, like La, is an element that can control the form of the sulfide with a small amount of addition, but is also an element that segregates at the grain boundary and causes a decrease in the number ratio of grain boundary 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 number ratio of grain boundary carbides decreases and cold workability decreases, so Ce is made 0.050% or less. Preferably it is 0.046% or less.

The balance of the component composition of the steel sheet of the present invention is Fe and inevitable impurities.

In addition, it may replace with a part of said Fe, and may contain 1 type or 2 types of Ti and B.

[Ti: 0.10% or less]
Ti is an element effective for controlling the form of carbide, and is also an element contributing to improvement of toughness by refining the structure. 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%, coarse Ti oxide is generated and becomes a starting point of cracking during cold working, so Ti is made 0.10% or less. Preferably it is 0.08% or less.

[B: 0.0001 to 0.010%]
B is an element that contributes to improving the toughness by increasing the hardenability during the heat treatment of the parts, making the structure uniform. 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%, coarse B oxide is generated and becomes a starting point of cracking during cold working, so B is made 0.010% or less. Preferably it is 0.009% or less.

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

The structure of the steel sheet of the present invention is substantially a structure composed of ferrite and carbide. In addition to cementite (Fe ₃ C), which is a compound of iron and carbon, carbides include compounds in which Fe atoms in cementite are substituted with alloy elements such as Mn and Cr, and alloy carbides (M ₂₃ C ₆ , M ₆ C MC, etc. [M: Fe and other metal elements added as alloys]).

When a steel sheet is formed into a predetermined shape, a shear band is formed in the macro structure of the steel sheet, and slip deformation is concentrated near the shear band. Slip deformation is accompanied by dislocation growth, and a region having a high dislocation density is formed in the vicinity of the shear band. As the amount of strain applied to the steel sheet increases, slip deformation is promoted and the dislocation density increases.

In cold forging, strong processing exceeding 1 is applied with considerable strain. For this reason, in the conventional steel plate, generation | occurrence | production of the void and / or a crack accompanying the increase in a dislocation density cannot be prevented, and the improvement of cold forgeability was difficult in the conventional steel plate. In order to solve this problem, it is effective to suppress the formation of shear bands during molding.

From the viewpoint of the microstructure, the formation of a shear band is understood as a phenomenon in which a slip generated in one crystal grain overcomes the grain boundary and continuously propagates to adjacent crystal grains. Therefore, in order to suppress the formation of shear bands, it is necessary to prevent the propagation of slip across the grain boundary.

Carbides in the steel sheet are strong particles that prevent slipping, and by allowing the carbides to exist at the ferrite grain boundaries, it is possible to prevent the propagation of slips across the crystal grain boundaries and suppress the formation of shear bands. It becomes possible to improve cold forgeability. At the same time, the formability of the steel sheet is improved.

The formability of a steel sheet is largely due to the accumulation of strain (accumulation of dislocations) in the crystal grains. If the propagation of strain to adjacent crystal grains is prevented at the grain boundaries, the amount of strain in the crystal grains is reduced. Increase. As a result, the work hardening rate is increased and the moldability is improved.

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

In order to measure the carbide coverage at the ferrite grain boundary in a three-dimensional space, serial sectioning SEM observation or three-dimensional EBSP observation in which sample cutting and observation by FIB is repeated in a scanning electron microscope is essential. As a result, enormous measurement time is required and accumulation of technical know-how is essential. The present inventors clarified this and concluded that a general analysis method is not suitable.

For this reason, as a result of searching for a simple and highly accurate evaluation index, it is possible to evaluate cold workability by using the ratio of the number of carbides at the ferrite grain boundary 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 in the ferrite grain boundary to the number of carbides in the ferrite grains exceeds 1, the cold workability is remarkably improved.

In addition, since any buckling, folding, and folding of the steel sheet that occurs during cold working are caused by the localization of strain accompanying the formation of shear bands, the presence of carbides at the ferrite grain boundaries can cause shearing. Band formation and strain localization can be alleviated, and buckling, folding, and folding can be effectively suppressed.

When the spheroidization rate of the carbide on the grain boundary is less than 80%, strain is concentrated locally on the rod-like or plate-like carbide, and voids and / or cracks are likely to occur. The spheroidization rate of the carbide is preferably 80% or more, and more preferably 90% or more.

If the average particle diameter of the carbide is less than 0.1 μm, the hardness of the steel sheet is remarkably increased and the workability is lowered, so the average particle diameter of the carbide is preferably 0.1 μm or more. More preferably, it is 0.17 μm or more. On the other hand, if the average particle diameter of the carbide exceeds 2.0 μm, coarse carbides are generated as a starting point during cold processing, cracking occurs, and cold workability is deteriorated. Therefore, the average particle diameter of the carbide is 2.0 μm or less. preferable. More preferably, it is 1.95 μm or less.

Carbide is observed with a scanning electron microscope. Prior to observation, a sample for tissue observation was wet-polished with emery paper and polished with diamond abrasive grains having an average particle size of 1 μm, and the observation surface was finished to a mirror surface, and the tissue was then washed with a 3% nitric acid-alcohol solution. Etch. The magnification for observation is selected to be a magnification capable of discriminating ferrite and carbides from 3000 times. At the selected magnification, 8 images of a 30 μm × 40 μm field of view in a 1/4 layer thickness are taken at random.

For the obtained tissue image, the area of each carbide contained in the region is measured in detail by image analysis software typified by Mitani Corporation (Win ROOF). The equivalent circle diameter (= 2 × √ (area / 3.14)) is determined from the area of each carbide, and the average value is defined as the carbide particle diameter. Further, the spheroidization rate of the carbide was obtained by calculating the ratio of the carbide that approximates an ellipse having the same area and the same moment of inertia, and the ratio of the maximum length to the maximum length in the perpendicular direction is less than 3. .

In order to suppress the influence of measurement error due to noise, carbides having an area of 0.01 μm ² or less were excluded from the evaluation targets. The number of carbides present on the ferrite grain boundaries was counted, and the number of carbides in the ferrite grains was determined by subtracting the number of carbides on the grain boundaries from the total number of carbides. Based on the measured number, the number ratio of the carbide on the ferrite grain boundary to the carbide in the ferrite grain was determined.

The cold workability can be improved by setting the ferrite grain size to 5.0 μm or more in the structure after annealing the cold-rolled steel sheet. If the ferrite particle size is less than 5 μm, the hardness increases and cracks and cracks are likely to occur during cold working, so the ferrite particle size is set to 5 μm or more. Preferably it is 7 micrometers or more.

On the other hand, if the thickness exceeds 50 μm, the number of carbides on the grain boundaries that suppress the propagation of slip is reduced and the cold workability is lowered, so the ferrite grain size is set to 50 μm or less. Preferably it is 37 micrometers or less.

The ferrite grain size was measured by polishing the sample observation surface to a mirror surface using the polishing method described above, etching with a 3% nitric acid-alcohol solution, and observing the structure of the observation surface with an optical microscope or a scanning electron microscope. The line segment method is applied to the image and measured.

Also, cementite, which is a carbide of iron, has a hard and brittle structure, and if it exists in the state of pearlite, which is a layered structure with ferrite, the steel becomes hard and brittle. Therefore, it is necessary to reduce pearlite as much as possible. In the steel sheet of the present invention, the area ratio is set to 6% or less.

Since perlite has a unique lamellar structure, it can be distinguished by SEM and optical microscope observation. The area ratio of pearlite can be obtained by calculating the region of the lamellar structure in an arbitrary cross section.

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

On the other hand, if the Vickers hardness exceeds 170 HV, the ductility is lowered and internal cracks are likely to occur during cold working, so the Vickers hardness is set to 170 HV or less. Preferably it is 168HV or less.

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

The manufacturing method of the present invention is based on the basic idea that the steel strip having the above-described composition is used to consistently manage the hot rolling conditions and the annealing conditions and to control the structure of the steel sheet.

First, a steel slab in which molten steel having a required composition is continuously cast is subjected to hot rolling. The slab after continuous casting may be directly subjected to hot rolling, or may be subjected to hot rolling after being once cooled and heated.

When the steel slab is once cooled and then heated and subjected to hot rolling, the heating temperature is preferably 1000 ° C. or more and 1250 ° C. or less, and the heating time is preferably 0.5 hours or more and 3 hours or less. When the continuously cast steel slab is directly subjected to hot rolling, the temperature of the steel slab subjected to hot rolling is preferably 1000 ° C. or more and 1250 ° C.

If the slab temperature or slab heating temperature exceeds 1250 ° C, or if the slab heating time exceeds 3 hours, decarburization from the slab surface layer becomes significant, and during heating before carburizing and quenching, austenite grains on the steel sheet surface layer Grows abnormally and impact resistance decreases. For this reason, the slab temperature or the slab heating temperature is preferably 1250 ° C. or less, and the heating time is preferably 3 hours or less. More preferably, it is 1200 degrees C or less and 2.5 hours or less.

If the billet temperature or billet heating temperature is less than 1000 ° C, or if the heating time is less than 0.5 hours, microsegregation and macrosegregation generated by casting will not be eliminated, and Si or A region where an alloy element such as Mn is locally concentrated remains and impact resistance is lowered. For this reason, the steel slab temperature or the steel slab heating temperature is preferably 1000 ° C. or more, and the heating time is preferably 0.5 hours or more. More preferably, it is 1050 ° C. or more and 1 hour or more.

Finish rolling in hot rolling is completed in a temperature range of 750 ° C. or higher and 850 ° C. or lower. When the finish rolling temperature is less than 750 ° C., the deformation resistance of the steel sheet increases, the rolling load increases significantly, the roll wear amount increases, the productivity decreases, and the plastic anisotropy is improved. Since the recrystallization necessary for this does not proceed sufficiently, the finish rolling temperature is set to 750 ° C. or higher. In terms of promoting recrystallization, the temperature is preferably 770 ° C. or higher.

When the finish rolling temperature exceeds 850 ° C., a thick scale is generated in the run-out table (ROT) through the plate, resulting in wrinkles on the surface of the steel plate, after cold forging and carburizing and tempering. When an impact load is applied, cracks are likely to occur starting from wrinkles, so that the impact resistance of the steel sheet is reduced. For this reason, finish rolling temperature shall be 850 degrees C or less. Preferably it is 830 degrees C or less.

When the hot-rolled steel sheet after finish rolling is cooled by ROT, the cooling rate is preferably 10 ° C./second or more and 100 ° C./second or less. When the cooling rate is less than 10 ° C./second, a thick scale is formed during the cooling, and the generation of wrinkles due to the scale cannot be suppressed, and the impact resistance is reduced. Therefore, the cooling rate is preferably 10 ° C./second or more. . More preferably, it is 20 ° C./second or more.

When the steel sheet is cooled from the surface layer to the inside at a cooling rate exceeding 100 ° C./second, the outermost layer part is excessively cooled, and low temperature transformation structures such as bainite and martensite are generated. When the hot-rolled steel sheet coil cooled to 100 ° C. to room temperature is taken out after winding, microcracks are generated in the low-temperature transformation structure. It is difficult to remove these micro cracks by pickling and cold rolling.

And, when an impact load is applied to a steel sheet after cold forging and carburizing quenching and tempering, the crack develops starting from a microcrack, so the impact resistance is lowered. For this reason, in order to suppress generation | occurrence | production of low temperature transformation structures, such as a bainite and a martensite, in the outermost layer part of a steel plate, 100 degrees C / sec or less is preferable. More preferably, it is 90 ° C./second or less.

In addition, the cooling rate is determined at each water injection section from the time when the hot-rolled steel sheet after finish rolling passes through the non-water injection section to receive water cooling in the water injection section to the time when it is cooled on the ROT to the winding target temperature. It refers to the cooling capacity received from the cooling equipment, and does not indicate the average cooling rate from the water injection start point to the temperature taken up by the winder.

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

When the coiling temperature is less than 400 ° C., the austenite that has not been transformed before winding is transformed into hard martensite, and when the hot-rolled steel sheet coil is discharged, cracks occur in the surface layer of the hot-rolled steel sheet, resulting in impact resistance. Sexuality decreases.

Furthermore, since the recrystallization driving force is small during recrystallization from austenite to ferrite, the orientation of the recrystallized ferrite grains is strongly influenced by the orientation of the austenite grains, making it difficult to randomize the texture. Therefore, the winding temperature is 400 ° C. or higher. Preferably it is 430 degreeC or more.

When the coiling temperature exceeds 550 ° C., pearlite with large lamella spacing is generated, and thick needle-like carbides with high thermal stability are generated. This acicular carbide remains even after two-stage annealing. At the time of forming such as cold forging of a steel plate, a crack is generated starting from this acicular carbide.

Also, when recrystallizing ferrite from austenite, on the contrary, the recrystallization driving force becomes too large, and even in this case, the recrystallized ferrite grains strongly depend on the orientation of the austenite grains, and the texture is not randomized. Therefore, the coiling temperature is 550 ° C. or less. Preferably it is 520 degrees C or less.

¡Take out the hot-rolled steel sheet coil, pickle it, and then perform two-step annealing (two-step annealing) that keeps it in two temperature ranges. By subjecting the hot-rolled steel sheet to two-stage annealing, the stability of the carbide is controlled and the formation of carbide at the ferrite grain boundary is promoted.

If the steel plate after pickling is cold-rolled before the annealing treatment, the ferrite grains become finer, so that the steel plate becomes difficult to soften. Therefore, in the present invention, it is not preferable to perform cold rolling before annealing, and it is preferable to perform annealing treatment without pickling after pickling.

The first stage annealing is performed in a temperature range of 650 to 720 ° C., preferably the A _c1 point or less. By this annealing, the carbide is coarsened and partially spheroidized, and the alloy elements are concentrated in the carbide, thereby improving 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 30 ° C./hour or more and 150 ° C./hour or less. If the first stage heating rate is less than 30 ° C./hour, it takes time to raise the temperature and the productivity is lowered. Therefore, the first stage heating rate is set to 3 ° C./hour or more. Preferably, it is 10 ° C./hour or more.

On the other hand, if the first stage heating rate exceeds 150 ° C./hour, the temperature difference between the outer peripheral portion and the inside of the hot-rolled steel sheet coil increases, and slag and seizure due to the difference in thermal expansion occurs. Unevenness is formed on the surface. When forming such as cold forging, cracks are generated as a starting point, and cold forgeability is deteriorated, and impact resistance after carburizing and quenching and tempering is reduced. It shall be below ℃ / hour. Preferably it is 130 degrees C / hour or less.

The annealing temperature in the first stage annealing (hereinafter referred to as “first stage annealing temperature”) is 650 ° C. or more and 720 ° C. or less. If the first stage annealing temperature is less than 650 ° C., the carbide is not sufficiently stabilized, and it becomes difficult to leave the carbide in the austenite during the second stage annealing. For this reason, the first stage annealing temperature is set to 650 ° C. or higher. Preferably it is 670 degreeC or more.

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

The annealing time in the first stage annealing (hereinafter referred to as “first stage annealing time”) is 3 hours or more and 60 hours or less. If the first stage annealing time is less than 3 hours, the carbide is not sufficiently stabilized, and it becomes difficult to leave the carbide in the austenite during the second stage annealing. For this reason, the first stage annealing time is set to 3 hours or more. Preferably it is 5 hours or more.

On the other hand, if the first stage annealing time exceeds 60 hours, further stabilization of the carbide cannot be expected, and further, the productivity is lowered. Therefore, the first stage annealing time is set to 60 hours or less. Preferably it is 55 hours or less.

Thereafter, the temperature is raised to 725 to 790 ° C., preferably in the temperature range from A _{c1 to} A ₃ , and austenite is generated in the structure. At this time, the carbides in the fine ferrite grains are dissolved in the austenite, but the carbides coarsened by the first stage annealing remain in the austenite.

When cooled without performing the second stage annealing, the ferrite grain size does not increase and an ideal structure cannot be obtained.

The heating rate of the second stage annealing to the annealing temperature (hereinafter referred to as “second stage heating rate”) is 1 ° C./hour or more and 80 ° C./hour or less. During the second stage annealing, austenite is generated and grows from the ferrite grain boundary. At that time, by slowing the heating rate up to the annealing temperature, it becomes possible to suppress austenite nucleation and increase the grain boundary coverage of the carbide in the structure formed by annealing after annealing.

Therefore, it is preferable that the second stage heating rate is slow. However, if it is less than 1 ° C./hour, it takes time to raise the temperature and the productivity decreases, so the second stage heating rate is 1 ° C./hour or more. And Preferably, it is 10 ° C./hour or more.

When the second stage heating rate exceeds 80 ° C./hour, in the hot-rolled steel sheet coil, the temperature difference between the outer peripheral portion and the inside increases, and scouring and seizure due to a large difference in thermal expansion due to transformation occurs. Unevenness is formed on the surface of the steel plate. At the time of cold forging, cracks are generated starting from this unevenness, cold forgeability and formability are reduced, and impact resistance after carburizing and quenching and tempering is also reduced, so the second stage heating rate is 80 ° C / Less than hours. Preferably it is 70 degrees C / hour or less.

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

On the other hand, if the second stage annealing temperature exceeds 790 ° C., it becomes difficult to leave the carbides in the austenite and it becomes difficult to control the structure change, so the second stage annealing temperature is set to 790 ° C. or less. Preferably it is 770 degrees C or less.

The annealing time in the second stage annealing (second stage annealing time) is 3 hours or more and less than 50 hours. If the second stage annealing time is less than 3 hours, the amount of austenite produced is small, and the dissolution of carbides in the ferrite grains does not proceed sufficiently, making it difficult to increase the number of carbides at the ferrite grain boundaries, In addition, the ferrite grain size is reduced. For this reason, the second stage annealing time is set to 3 hours or more. Preferably it is 5 hours or more.

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

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

By slow cooling, the austenite generated in the second stage annealing is transformed into ferrite, carbon atoms are adsorbed on the carbide remaining in the austenite, and the carbide and austenite cover the ferrite grain boundary, and finally In addition, a structure in which a large number of carbides exist in the ferrite grain boundary can be obtained.

For this purpose, it is preferable that the cooling rate is slow, but if it is less than 1 ° C./hour, the time required for cooling increases and the productivity decreases, so the cooling rate is 1 ° C./hour or more. Preferably, it is 10 ° C./hour or more.

On the other hand, when the cooling rate exceeds 30 ° C./hour, austenite transforms into pearlite, the hardness of the steel sheet increases, cold forgeability decreases, and impact resistance after carburizing and quenching and tempering decreases. Therefore, the cooling rate is set to 30 ° C./hour or less. Preferably it is 20 degrees C / hour or less.

Furthermore, the steel sheet cooled to 650 ° 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, any atmosphere of 95% or more nitrogen atmosphere, 95% or more hydrogen atmosphere, or air atmosphere may be used.

As explained above, according to the manufacturing method that consistently manages the hot rolling conditions and annealing conditions of the present invention and performs the structure control of the steel sheet, the formability during cold forging combined with drawing and thickening is achieved. Further, it is possible to produce a steel sheet that is excellent and further has excellent hardenability necessary for improving impact resistance after carburizing, quenching, and tempering.

Next, although an Example is described, the level of an Example is an example of the conditions employ | adopted in order to confirm the feasibility and effect of this invention, and this invention is limited to this one condition example. is not. 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.

Cold workability is evaluated by taking a JIS No. 5 tensile specimen from an as-annealed material with a thickness of 3 mm and conducting a tensile test to evaluate the total elongation in the 0 ° direction from the rolling direction and the 90 ° direction from the rolling direction. The cold workability is said to be superior when both directions are 35% or more and the total elongation difference | ΔEL | in each direction is 4% or less.

The evaluation of the hardenability was performed by grinding the material as it was annealed with a plate thickness of 3 mm to a plate thickness of 1.5 mm, holding at 880 ° C. for 10 minutes in a vacuum atmosphere, and quenching at a cooling rate of 30 ° C./second, If the martensite fraction is 60% or more, it is said that the hardenability is superior.

Example 1
A continuous cast slab (steel ingot) having the composition shown in Table 1 was heated at 1240 ° C. for 1.8 hours, then subjected to hot rolling, and after finishing hot rolling at 890 ° C., it was wound at 510 ° C. A hot rolled coil having a thickness of 3.0 mm was manufactured. The hot rolled coil is pickled, the hot rolled coil is placed in a box-type annealing furnace, the atmosphere is controlled to 95% hydrogen-5% nitrogen, heated from room temperature to 705 ° C., and maintained at 705 ° C. for 36 hours. The temperature distribution in the hot-rolled coil was made uniform, and then heated to 760 ° C. and held at 760 ° 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 prepare a sample for characteristic evaluation. In addition, the structure | tissue of the sample was measured by the method mentioned above.

Table 2 shows the results of measuring or evaluating the Vickers hardness of the manufactured sample, the ratio of the number of carbides on the ferrite grain boundary to the number of carbides in the ferrite grains, the pearlite area ratio, cold workability, and hardenability. .

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, P-1, R-1, T-1, W-1, X-1, Y-1, Z-1, AB-1, and AC-1 all have a ferrite grain boundary relative to the number of carbides in the ferrite grain. The ratio of the number of carbides exceeds 1, and the Vickers hardness is 170 HV or less, which is excellent in cold workability and hardenability.

In contrast, Comparative Steel G-1 had a high C content and cold workability decreased. The comparative steel O-1 has a high Mo content and Cr content and high carbide stability. Therefore, the carbide does not dissolve during quenching, the austenite generation amount is small, and the hardenability is inferior.

Comparative steels Q-1 and AD-1 have a high amount of Si and Al and a high A3 point. Therefore, the amount of austenite produced during quenching is small, and the hardenability is inferior. In Comparative Example U-1, the amount of S is high, coarse MnS is generated in the steel, and the cold workability is low. Comparative Example AA-1 has a low Mn content and inferior hardenability.

Comparative Example I-1 had a low hot-rolling finishing temperature, resulting in decreased productivity. In Comparative Example D-1, the hot rolling finishing temperature was high, and scale wrinkles were formed on the steel sheet surface. In Comparative Examples C-1 and S-1, the hot rolling coiling temperature is low, the low temperature transformation structure such as bainite and martensite is increased and embrittled, and cracks occur frequently when the hot rolled coil is discharged, resulting in increased productivity. Declined.

In Comparative Examples A-1 and V-1, the hot rolling milling temperature is high, and in the hot rolled structure, thick pearlite having a lamellar interval and acicular coarse carbide with high thermal stability are generated. Even after the step-step annealing, it remained in the steel sheet and the cold workability decreased.

(Example 2)
In order to investigate the influence of the annealing conditions, a steel slab having the composition shown in Table 1 was heated at 1240 ° C. for 1.8 hours, then subjected to hot rolling, and after finishing hot rolling at 820 ° C., 45 The steel sheet was cooled to 520 ° C. at a cooling rate of ° C./second, wound at 510 ° C. to produce a hot-rolled coil with a plate thickness of 3.0 mm, and subjected to a two-step type box annealing under the annealing conditions shown in Table 3, A sample having a thickness of 3.0 mm was produced.

Table 3 shows the carbide diameter, ferrite particle diameter, Vickers hardness, ratio of the number of carbides on the ferrite grain boundary to the number of carbides in the ferrite grains, pearlite area ratio, cold workability, and hardenability. The result of having been measured or evaluated is shown.

As shown in Table 3, the inventive steels B-2, C-2, E-2, F-2, H-2, I-2, J-2, K-2, M-2, N-2, For R-2, S-2, V-2, Z-2, and AC-2, the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains exceeds 1, and the Vickers hardness Is 170 HV or less, and is excellent in cold workability and hardenability.

In contrast, Comparative Steel G-1 had a high C content and cold workability decreased. The comparative steel O-1 had a high Mo content and Cr content, and cold workability decreased. Moreover, since the carbide has high stability, the carbide does not dissolve during quenching, the austenite generation amount is small, and the hardenability is inferior.

Comparative Steel Q-1 had a high Si content and a high hardness of ferrite, so that the workability was lowered. Further, since the A3 point is high, the amount of austenite produced during quenching is small, and the hardenability is inferior. Since the comparative steel AD-1 has a high Al content and a high A3 point, the amount of austenite produced during quenching is small and the hardenability is inferior. Comparative steel U-1 had a high amount of S, and coarse MnS was produced in the steel, resulting in a decrease in cold workability. Comparative steel AA-1 has a low Mn content and inferior hardenability.

The comparative steel T-2 has a low holding temperature during the first stage annealing of the two-step type box annealing, the carbide coarsening treatment below the Ac1 temperature is insufficient, and the thermal stability of the carbide is low. By being insufficient, the carbides remaining at the second stage of annealing decreased, the pearlite transformation could not be suppressed in the structure after the slow cooling, and the cold workability was lowered.

The comparative steel A-2 has a high holding temperature during the first stage annealing of the two-step type box annealing, austenite is generated during the annealing, and the stability of the carbide cannot be increased. The carbide remaining at the time of annealing decreased, and the pearlite transformation could not be suppressed in the structure after the slow cooling, resulting in a decrease in cold workability.

The comparative steel L-2 has a short holding time during the first stage annealing of the two-step type box annealing, the carbide coarsening treatment below the Ac1 temperature is insufficient, and the thermal stability of the carbide is low. By being insufficient, the carbides remaining at the second stage of annealing decreased, the pearlite transformation could not be suppressed in the structure after the slow cooling, and the cold workability was lowered.

The comparative steel W-2 had a long holding time during the first stage annealing during the two-step annealing, and the productivity decreased. The comparative steel X-2 has a low holding temperature during the second stage annealing during the two-step annealing, and the amount of carbides at the grain boundaries cannot be increased because the austenite generation amount is small, resulting in a decrease in cold workability. did.

The comparative steel AB-2 has a high holding temperature during the second stage annealing of the two-step type box annealing, and the dissolution of carbides is accelerated, so that the remaining carbides are reduced, and the pearlite transformation occurs in the structure after the slow cooling. Could not be suppressed, and cold forgeability was reduced.

The comparative steel P-2 has a low holding temperature during the second stage annealing of the two-step type box annealing, a small amount of austenite is generated, and the number ratio of carbides at the ferrite grain boundaries cannot be increased. Inter-workability decreased. The comparative steel Y-2 has a long holding time during the second stage annealing of the two-step type box annealing, and the dissolution of carbides is accelerated, so that the remaining carbides are reduced, and the pearlite transformation occurs in the structure after the slow cooling. Could not be suppressed, and cold forgeability was reduced.

Comparative steel D-2 had a high cooling rate from the end of the second stage annealing of the two-step type box annealing to 650 ° C., and pearlite transformation occurred during cooling, resulting in a decrease in cold workability.

As described above, according to the present invention, it is possible to manufacture and provide a steel sheet having excellent formability and wear resistance. Since the steel sheet of the present invention is a steel sheet suitable as a material for automobile parts, blades, and other machine parts manufactured through processing steps such as punching, bending, and pressing, the present invention has industrial applicability. It is expensive.

Claims

% By mass
C: 0.10 to 0.40%,
Si: 0.01 to 0.30%,
Mn: 1.00 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,
Cr: 0.50% or less,
Mo: 0.10% or less,
Nb: 0.10% or less,
V: 0.10% or less,
Cu: 0.10% or less,
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: a steel plate containing 0.050% or less, the balance being Fe and inevitable impurities,
The ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains of the steel sheet is more than 1,
The ferrite grain size is 5 μm or more and 50 μm or less, and the area ratio of pearlite is 6% or less,
A steel sheet, wherein the steel sheet has a Vickers hardness of 100 HV or more and 170 HV or less.
Instead of a part of the Fe,
Ti: 0.10% or less, and B: 0.010% or less,
The steel plate according to claim 1, comprising one or two of the following.
It is a manufacturing method which manufactures the steel plate according to claim 1 or 2,
The steel slab having the component composition according to claim 1 or 2 is subjected to hot rolling to complete finish rolling in a temperature range of 750 ° C. or higher and 850 ° C. or lower to obtain a hot-rolled steel plate,
Winding the hot-rolled steel sheet at 400 ° C. or higher and 550 ° C. or lower,
Pickled hot rolled steel sheet,
The pickled hot-rolled steel sheet is annealed in the first stage for holding in the temperature range of 650 ° C. to 720 ° C. for 3 hours to 60 hours,
Second-stage annealing is performed to hold the hot-rolled steel sheet in a temperature range of 725 ° C. to 790 ° C. for 3 hours to 50 hours,
A method for producing a steel sheet, comprising: cooling a hot-rolled steel sheet after annealing to 650 ° C. at a cooling rate of 1 ° C./hour to 30 ° C./hour.