KR102029565B1 - Steel plate and its manufacturing method - Google Patents

Abstract

It is a low carbon steel sheet that has excellent impact resistance after cold forging and carburizing quenching tempering, has a predetermined composition, and has an average particle diameter of carbide of 0.4 µm or more and 2.0 µm or less, an area ratio of pearlite of 6% or less, and the number of carbides in the ferrite particles. The ratio of the number of carbides in the ferrite grain boundary to is greater than 1, and the Vickers hardness is 100HV or more and 180HV or less.

Description

Steel plate and its manufacturing method

The present invention relates to a steel sheet and a method of manufacturing the same.

The steel sheet containing 0.1 to 0.4% of carbon by mass is subjected to cold forging from the blank material by press forming, hole expanding forming, bending forming, drawing forming, thickness increasing and thickness reducing forming and combination thereof. It is used as a raw material for drive system components such as gears and clutches. Compared with the conventional hot forging and the like, in cold forging, there is a problem that the amount of deformation accumulated in the raw material is increased, causing cracking of the raw material and buckling at the time of molding, thereby causing deterioration of component characteristics.

In particular, after carburizing, quenching and tempering are performed on the molded material in order to obtain wear resistance, residual stresses are generated by heat treatment, and thus, cracks and buckles are generated from the cracks and buckles. In order to use it as a drive system component, the above-described steel sheet has excellent cold forging because it is required to acquire impact resistance properties that do not break brittlely against a load of momentarily large loads, such as by a gear shifting during startup. Performance and impact resistance after carburizing quenching and tempering are required.

Until now, many proposals have been made about the technique of improving the cold forging of a steel plate and the impact resistance property after carburizing (for example, refer patent documents 1-5).

For example, Patent Literature 1 discloses a mechanical structural steel in which toughness is improved by suppressing coarsening of crystal grains in carburizing heat treatment, in terms of mass%, C: 0.10 to 0.30%, Si: 0.05 to 2.0%, and Mn: 0.10 to 0.50%, P: 0.030% or less, S: 0.030% or less, Cr: 1.80 to 3.00%, Al: 0.005 to 0.050%, Nb: 0.02 to 0.10%, N: 0.0300% or less, and the balance Fe and There is disclosed a mechanical structural steel composed of unavoidable impurities, wherein the structure before cold working is a ferrite / pearlite structure, and the average value of the ferrite particle diameter is 15 µm or more.

Patent Literature 2 discloses a steel having excellent cold workability and carburizing resistance, and includes C: 0.15 to 0.40%, Si: 1.00% or less, Mn: 0.40% or less, sol.Al:0.02% or less, N: 0.006% or less, and B: Disclosed is a steel containing 0.005 to 0.050%, a balance consisting of Fe and inevitable impurities, and having a structure mainly composed of a ferrite phase and a graphite phase.

Patent Document 3 discloses a steel for carburizing bevel gears excellent in impact strength, a high toughness carburizing bevel gear, and a manufacturing method thereof.

Patent Document 4 discloses cold forging after spheroidizing annealing, and has excellent workability for parts produced in a carburizing-quenching tempering step, while suppressing coarsening of crystal grains in subsequent carburization, and excellent impact resistance and impact fatigue characteristics. Steel for carburizing parts having the same is disclosed.

Patent document 5 contains C: 0.40 to 0.80%, Si: 0.05 to 1.50%, Mn: 0.05 to 1.50% and V: 1.8 to 6.0% as the cold tool steel for plasma carburization, Ni: 0.10 to 2.50%, Cr A steel containing one or two or more of 0.1 to 2.0% and Mo: 3.0% or less, and the balance is made of Fe and unavoidable impurities is disclosed.

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

The structure of the mechanical structural steel of Patent Literature 1 is a structure of ferrite + pearlite, and since the structure has a large hardness as compared with the ferrite + cementite structure, the wear and tear of the metal mold in cold forging cannot be suppressed, and it is not necessarily cold stage. It cannot be called a mechanical structural steel with excellent composition.

In the steel of patent document 2, annealing at high temperature becomes essential for the graphite process of cementite, and it cannot suppress the fall of a yield and the increase of a manufacturing cost.

Since the manufacturing method of patent document 3 needs to perform further hot forging after cold forging and carburizing, and hot forging becomes essential, it is not a manufacturing method which leads to radical cost reduction.

It is unclear whether the steel for carburizing parts of patent document 4 can exhibit the same effect in cold forging to which a large deformation | transformation is given, and since the specific structure form and the control method of a structure are also unclear, the board | substrate which becomes widespread in recent years In the case of forming forging by forging a large deformation by cold such as forging, it is not a steel showing excellent workability.

Patent Document 5 does not disclose any knowledge or technology regarding the optimum component and structure for improving the formability of steel, especially cold forging.

SUMMARY OF THE INVENTION The present invention provides a steel sheet excellent in cold forging and impact resistance after carburizing and quenching tempering, and is particularly suitable for obtaining parts such as high cycle gears by sheet forming, and a method of manufacturing the same. It is a task.

In order to solve the above problems and obtain a steel sheet suitable for a material such as a drive system component, in the steel sheet containing C necessary to increase the quenchability, the grain size of the ferrite is increased and the carbide (mainly cementite) is spheroidized to an appropriate grain diameter. It is understood that the perlite structure can be reduced. This is based on the following reasons.

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

The carbide is indispensable to drive system components because the carbide can be imparted with excellent wear resistance and rolling fatigue characteristics while maintaining material formability by appropriately dispersing it in the metal structure. In addition, the carbide in the steel sheet is a hard particle that prevents slipping, and the presence of carbide in the ferrite grain boundary prevents the propagation of slip exceeding the grain boundary, suppresses the formation of the shear zone, and improves the cold forging. The formability of the steel sheet is also improved.

However, cementite is hard and embrittlement structure, and when it exists in the state of pearlite which is a layer structure with ferrite, since steel is hard and embrittlement, it is necessary to exist in spherical form. In consideration of cold forging and generation of cracks during forging, the particle size needs to be in an appropriate range.

However, the manufacturing method for realizing the above structure is not disclosed until now. Therefore, the present inventors earnestly researched the manufacturing method for realizing the said structure.

As a result, in order to make the metal structure of the steel plate after the winding after hot rolling into the bainite structure which the cementite disperse | distributed in the fine pearlite or the fine ferrite with a small lamella spacing, it winds up at comparatively low temperature (400 degreeC-550 degreeC). By winding at a relatively low temperature, cementite dispersed in ferrite also becomes easier to spheroidize. Subsequently, cementite is partially spheroidized by annealing at a temperature just below the Ac1 point as the first stage annealing. Subsequently, the austenite transformation is performed while leaving part of the ferrite particles by annealing at a temperature between the Ac1 and Ac3 points (the so-called two-phase region of ferrite and austenite) as the second stage annealing. Then, it was found that by slowly cooling and growing the ferrite particles left, the austenite was ferrite transformed into a nucleus therein to precipitate cementite at grain boundaries while obtaining a large ferrite phase, thereby realizing the above structure.

That is, it has been found that a method for producing a steel sheet that satisfies quenchability and formability at the same time is difficult to realize even if a single hot rolling condition or annealing condition is devised, and can be realized by achieving optimization in a so-called consistent process such as a hot rolling and annealing process. .

In addition, it has been found that the reduction of plastic anisotropy is necessary for the improvement of drawing formability during cold forging, and the adjustment of hot rolling conditions is important for this improvement.

This invention is made | formed based on these knowledge, The summary is as follows.

(1) The component composition is, in 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%, Cr: 0.50 to 2.00%, and 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: 0.10% or less, 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, Ce: 0.050%, and is a low-carbon steel sheet which is a residual Fe and impurities, and the metal structure of the low carbon steel sheet is a carbide grain size. The ratio of the number of carbides of ferrite grain boundary to the number of carbides in the ferrite grains is 0.4 to 2.0 µm, the pearlite area ratio is 6% or less, and more than 1, and the Vickers hardness of the low carbon steel sheet is 100 HV or more and 180 HV or less. Steel sheet according to claim.

(2) It is a manufacturing method of manufacturing the steel plate of said (1), The steel piece of the component composition of said (1) is made hot-rolled to complete finishing hot rolling in the temperature range of 650 degreeC or more and 950 degrees C or less, and it is set as a hot rolled steel sheet. The hot rolled steel sheet is wound at 400 ° C. or higher and 600 ° C. or lower, and the pickled hot rolled steel sheet is pickled, and the pickled hot rolled steel sheet is heated at a heating rate of 30 ° C./hour or more and 150 ° C./hour or less and 650 ° C. or higher and 720 ° C. or lower. The first stage annealing is performed at the following annealing temperature and maintained for 3 hours or more and 60 hours or less, and subsequently, the hot-rolled steel sheet is heated at a heating rate of 1 ° C./hour or more and 80 ° C./hour or less, from 725 ° C. to 790 ° C. It is heated to the following annealing temperature, and performs the 2nd stage annealing which hold | maintains for 3 hours or more and 50 hours or less, and cooling the hot-rolled steel sheet after annealing to 650 degreeC at the cooling rate of 1 degreeC / hour or more and 100 degrees C / hour or less. Characterized by cold end Composition and method of manufacturing the steel sheet.

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

BRIEF DESCRIPTION OF THE DRAWINGS It is a figure which shows the outline | summary of the cold forging test and the form of the crack introduce | transduced in cold forging. (a) shows the disk shaped test material cut out from the hot rolled steel plate, (b) shows the shape of the test material after cold forging, and (c) shows the cross-sectional shape of the test material after cold forging.
It is a figure which shows typically the outline of the drop test which evaluates the impact resistance characteristic of the sample which performed carburizing quenching tempering.
It is a figure which shows the relationship of the ratio of the number of grain boundary carbides with the number of carbides in a particle, the crack length of a cold forging test piece, and the impact resistance property after carburizing quenching tempering.
4 is a diagram showing another relationship between the ratio of the number of grain boundary carbides to the number of carbides in the particle, the crack length of the cold forging test piece, and the impact resistance after carburizing quenching tempering.

EMBODIMENT OF THE INVENTION Hereinafter, this invention is demonstrated in detail. First, the reason for limitation of the component composition of the steel plate of this invention is demonstrated. Here, "%" regarding a component composition means "mass%."

[C: 0.10 to 0.40%]

C is an element that forms carbide in steel and is effective for reinforcing steel and miniaturizing ferrite particles. In order to suppress creep generation in cold working and to secure the surface aesthetics of cold forging parts, suppression of the coarsening of the ferrite grain size is essential, but the volume fraction of carbide is insufficient at less than 0.10%, Since coarsening cannot be suppressed, C is made into 0.10% or more. Preferably it is 0.11% or more.

On the other hand, if the content exceeds 0.40%, the volume fraction of the carbide increases, and when a load is momentarily loaded, a large amount of cracks, which are the starting point of fracture, are generated, resulting in a decrease in the impact resistance, so that C is 0.40% or less. do. Preferably it is 0.38% or less.

[Si: 0.01 to 0.30%]

Si is an element that acts as a deoxidizer and also affects the form of carbides. In order to reduce the number of carbides in the ferrite particles to obtain the deoxidation effect and increase the number of carbides on the ferrite grain boundary, an austenite phase is generated during the annealing by annealing in two steps, and once the carbides are dissolved It is necessary to accelerate the slow cooling and to generate carbides in the ferrite grain boundary.

When Si exceeds 0.30%, the ductility of the ferrite is lowered, so that cracks tend to occur during cold forging, and the cold forging property and the impact resistance after carburizing and quenching tempering decrease, so that Si is 0.30% or less. Preferably it is 0.28% or less.

A smaller amount of Si is more preferable, but a reduction to less than 0.01% causes a significant increase in refining cost, so the Si is made 0.01% or more. Preferably it is 0.02% or more.

[Mn: 0.30 to 1.00%]

Mn is an element which controls the form of carbide in the two step step annealing. If it is less than 0.30%, it becomes difficult to produce carbide on the ferrite grain boundary in the slow cooling after the second stage annealing, so that Mn is 0.30% or more. Preferably it is 0.33% or more.

On the other hand, when it exceeds 1.00%, since toughness after carburizing-quenching tempering will fall, Mn shall be 1.00% or less. Preferably it is 0.96% or less.

[Al: 0.001 to 0.10%]

Al is an element that functions as a deoxidizer of steel to stabilize ferrite. If the addition effect is not sufficiently obtained at less than 0.001%, Al is made 0.001% or more. Preferably it is 0.004% or more.

On the other hand, when it exceeds 0.10%, since the number ratio of carbides in a grain boundary falls and the crack length at the time of cold forging is brought about, Al shall be 0.10% or less. Preferably it is 0.09% or less.

[Cr: 0.50 to 2.00%]

Cr and Mo are elements which improve toughness. Cr is an element effective for stabilizing carbides during heat treatment. If it is less than 0.50%, it becomes difficult to remain carbide at the time of carburizing, coarsening of the austenite grain size in the surface layer, and lowering of impact resistance characteristics, and Cr is made 0.50% or more. Preferably it is 0.52% or more.

On the other hand, when the content exceeds 2.00%, the amount of Cr concentration in the carbide increases, and many fine carbides remain in the austenite phase produced by the two-step annealing, so that carbides are present in the particles after slow cooling. Since the ratio of the increase of hardness and the number of grain boundary carbides falls and cold forging property falls, Cr is made into 2.00% or less. Preferably it is 1.94% or less.

[Mo: 0.001 to 1.00%]

Mo is an element effective for controlling the form of carbide. If the addition effect is not sufficiently obtained at less than 0.001%, Mo is made 0.001% or more. Preferably it is 0.017% or more.

On the other hand, if it exceeds 1.00%, Mo is concentrated in the carbide, and stable carbides increase in the austenite phase. Therefore, carbides are present in the particles after slow cooling, resulting in an increase in hardness and a decrease in the number ratio of grain boundary carbides. Since cold forging property falls, Mo is made into 1.00% or less. Preferably it is 0.94% or less.

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

[P: 0.020% or less]

P is an element which segregates at the ferrite grain boundary and suppresses the formation of grain boundary carbide. Less is preferable. Although content of P may be 0, in order to make it highly purified to less than 0.0001% in a refining process, refining requires a long time and causes a significant increase in manufacturing cost, and a practical minimum is 0.0001 to 0.0013%.

On the other hand, when it exceeds 0.020%, since the number ratio of grain boundary carbide falls and cold forging property falls, P shall be 0.020% or less. Preferably it is 0.018% or less.

[S: 0.010% or less]

S is an impurity element which forms nonmetallic inclusions, such as MnS. Since a nonmetallic inclusion becomes a starting point of a crack generation at the time of cold forging, less S is more preferable. Although content of S may be 0, when S is reduced to less than 0.0001%, refining cost will increase significantly, and a practical minimum is 0.0001 to 0.0012%.

On the other hand, when it exceeds 0.010%, since the crack length at the time of cold forging is increased, S shall be 0.010% or less. Preferably it is 0.009% or less.

[N: 0.020% or less]

N segregates to ferrite grain boundaries and is an element that suppresses the formation of carbides on grain boundaries. Less is preferable. Although content of N may be 0, when it reduces to less than 0.0001%, refining cost will increase significantly, and a practical minimum is 0.0001 to 0.0006%.

On the other hand, if it exceeds 0.020%, even if two-phase region annealing and slow cooling are performed, the ratio of the number of carbides on the ferrite grain boundary to the number of carbides in the ferrite particles is less than 1, and the cold forging property is lowered, so that N is 0.020%. It is set as follows. Preferably it is 0.017% or less.

[O: 0.0001 to 0.020%]

O is an element which forms an oxide in steel. Since the oxide present in the ferrite particles is a carbide production site, the smaller one is preferable. Although O content may be 0, when O is reduced to less than 0.0001%, refining cost will increase significantly, and a practical minimum 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 boundary to the number of carbides in the ferrite particles is less than 1, and the cold forging property is lowered, so that O is made 0.020% or less. Preferably it is 0.017% or less.

[Ti: 0.010% or less]

Ti is an important element for controlling the form of carbide, and is an element which promotes the production of carbide in ferrite particles by containing a large amount, and the smaller the more, the more preferable. Although content of Ti may be 0, when it reduces to less than 0.0001%, refining cost will increase significantly, and a practical minimum is 0.0001 to 0.0006%.

On the other hand, if it exceeds 0.010%, the ratio of the number of carbides on the ferrite grain boundary to the number of carbides in the ferrite particles is less than 1, and the cold forging property is lowered, so that Ti is 0.010% or less. Preferably it is 0.007% or less.

[B: 0.0005% or less]

B is an element effective for the control of slip of dislocation during cold forging. Since a large amount contains, the activity of a slip system is limited, so it is preferable that B is smaller. The content of B may be 0. Careful attention is required for the detection of less than 0.0001% of B, and depending on the analysis device, it reaches the detection limit or lower.

On the other hand, when it exceeds 0.0005%, since the cross slip of dislocation is suppressed in the shear zone formed by cold forging, deformation | concentration concentrates locally, and a crack generate | occur | produces, B shall be 0.0005% or less. Preferably it is 0.0005% or less.

[Sn: 0.050% or less]

Sn is an element mixed from steel raw materials (scrap), and it is so preferable that there is little. Although content of Sn may be 0, when it reduces to less than 0.001%, refining cost will increase significantly, and a practical minimum is 0.001 to 0.002%.

On the other hand, when it exceeds 0.050%, ferrite becomes brittle and cold forging property falls, so Sn is made into 0.050% or less. Preferably it is 0.048% or less.

[Sb: 0.050% or less]

Sb is an element mixed from steel raw material (scrap) similarly to Sn. Since Sb segregates at grain boundaries and lowers the number ratio of grain boundary carbides, the smaller the amount of Sb is, the more preferable. Although content of Sb may be 0, when it reduces to less than 0.001%, refining cost will increase significantly, and a practical minimum is 0.001 to 0.002%.

On the other hand, when it exceeds 0.050%, since cold forging property falls, Sb shall be 0.050% or less. Preferably it is 0.048% or less.

[As: 0.050% or less]

As is an element mixed from steel raw materials (scrap) similarly to Sn and Sb. As is segregated at the grain boundary and the number ratio of grain boundary carbides is lowered, so it is more preferable. Although content of As may be 0, when it reduces to less than 0.001%, refining cost will increase significantly, and a practical minimum is 0.001 to 0.002%.

On the other hand, when it exceeds 0.050%, since the number ratio of grain boundary carbides will fall and cold forging property will fall, As shall be 0.050% or less. Preferably it is 0.045% or less.

Although the steel plate of this invention makes the said element a basic element, in order to improve cold forging property and another characteristic, you may contain the following elements. Since the following elements are not essential in order to acquire the effect of this invention, content may be zero.

[Nb: 0.10% or less]

Nb is an element effective for controlling the form of carbide, and is an element which makes the structure fine and contributes to the improvement of toughness. If the addition effect is not sufficiently obtained at less than 0.001%, Nb is preferably at least 0.001%. More preferably, it is 0.002% or more.

On the other hand, when it exceeds 0.10%, many fine Nb carbides will precipitate, an intensity will increase excessively, the number ratio of grain boundary carbides will fall, and cold forging property will fall, and Nb shall be 0.10% or less. Preferably it is 0.09% or less.

[V: 0.10% or less]

V, like Nb, is also an element effective in controlling the shape of carbides, and is an element that makes the structure fine and contributes to the improvement of toughness. If the addition effect is not sufficiently obtained at less than 0.001%, V is preferably at least 0.001%. More preferably, it is 0.004% or more.

On the other hand, when it exceeds 0.10%, many fine V carbides will precipitate, intensity | strength will increase excessively, the ratio of the number of grain boundary carbides will fall, and cold forging property will fall, and V shall be 0.10% or less. Preferably it is 0.09% or less.

[Cu: 0.10% or less]

Cu is an element which forms fine precipitates and contributes to the improvement of strength. If it is less than 0.001%, since the strength improvement effect is not fully acquired, it is preferable to set Cu as 0.001% or more. More preferably, it is 0.008% or more.

On the other hand, when it exceeds 0.10%, red brittleness will be expressed during hot rolling, and since productivity will fall, Cu shall be 0.10% or less. Preferably it is 0.09% or less.

[W: 0.10% or less]

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

On the other hand, when it exceeds 0.10%, many fine W carbides will precipitate, intensity | strength will increase excessively, the number ratio of grain boundary carbides will fall, and cold forging property will fall, and W shall be 0.10% or less. Preferably it is 0.08% or less.

[Ta: 0.10% or less]

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

On the other hand, when it exceeds 0.10%, many fine W carbides will precipitate, intensity | strength will increase excessively, the number ratio of grain boundary carbide will fall, and cold forging property will fall, and Ta shall be 0.10% or less. Preferably it is 0.09% or less.

[Ni: 0.10% or less]

Ni is an element effective for improving the impact resistance property of a part. If the addition effect is not sufficiently obtained at less than 0.001%, Ni is preferably at least 0.001%. More preferably, it is 0.002% or more.

On the other hand, when it exceeds 0.10%, since the number ratio of grain boundary carbide falls and cold forging property falls, Ni shall be 0.10% or less. Preferably it is 0.09% or less.

[Mg: 0.050% or less]

Mg is an element that can control the form of sulfides by addition of trace amounts. If the addition effect is not sufficiently obtained at less than 0.0001%, the Mg is preferably 0.0001% or more. More preferably, it is 0.0008% or more.

On the other hand, when it exceeds 0.050%, ferrite becomes brittle and cold forging property falls, so Mg shall be 0.050% or less. Preferably it is 0.049% or less.

[Ca: 0.050% or less]

Ca, like Mg, is an element that can control the form of sulfides by addition of trace amounts. If the addition effect is not sufficiently obtained at less than 0.001%, Ca is preferably at least 0.001%. More preferably, it is 0.003% or more.

On the other hand, when it exceeds 0.050%, coarse Ca oxide will generate | occur | produce and since it will become a starting point of a crack generation at the time of cold forging, Ca shall be 0.050% or less. Preferably it is 0.04% 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 small amount. If the addition effect is not sufficiently obtained at less than 0.001%, Y is preferably at least 0.001%. More preferably, it is 0.003% or more.

On the other hand, when it exceeds 0.050%, coarse Y oxide will generate | occur | produce and it will be a starting point of a crack generation at the time of cold forging, so Y shall be 0.050% or less. Preferably it is 0.031% or less.

[Zr: 0.050% or less]

Zr, like Mg, Ca, and Y, is an element that can control the form of sulfide by addition of a small amount. If the addition effect is not sufficiently obtained at less than 0.001%, Zr is preferably at least 0.001%. More preferably, it is 0.004% or more.

On the other hand, if it exceeds 0.050%, coarse Zr oxide will be produced, and since it will be a starting point of crack generation at the time of cold forging, Zr shall be 0.050% or less. Preferably it is 0.045% or less.

[La: 0.050% or less]

La is an element which is effective for morphology control of sulfides by addition of a small amount, and segregates at grain boundaries and reduces the number ratio of grain boundary carbides. If the shape control effect is not sufficiently obtained at less than 0.001%, La is preferably at least 0.001%. More preferably, it is 0.003% or more.

On the other hand, when it exceeds 0.050%, since the number ratio of grain boundary carbide falls and cold forging property falls, La shall be 0.050% or less. Preferably it is 0.047% or less.

[Ce: 0.050% or less]

Like La, Ce is an element that can control the form of sulfide by addition of a small amount, and is an element that segregates at grain boundaries and reduces the ratio of the number of grain boundary carbides. If it is less than 0.001%, since a form control effect is not fully acquired, it is preferable to make Ce into 0.001% or more. More preferably, it is 0.003% or more.

On the other hand, when it exceeds 0.050%, since the number ratio of grain boundary carbides will fall and cold forging property will fall, Ce shall be 0.050% or less. Preferably it is 0.046% or less.

In addition, the balance of the component composition of the steel sheet of the present invention is Fe and unavoidable impurities.

Next, the steel plate structure of this invention is demonstrated.

The structure of the steel sheet of the present invention is a structure substantially composed of ferrite and carbide. Carbide is a compound obtained by substituting Fe atoms in cementite with Mn, Cr, etc., in addition to cementite (Fe ₃ C), which is a compound of iron and carbon, alloy carbides (M ₂₃ C ₆ , M ₆ C, MC, etc.), and M is Fe. And other metal elements).

When the steel sheet is molded into a predetermined part shape, a shear zone is formed in the macrostructure of the steel sheet, and slip deformation is concentrated in the vicinity of the shear zone. Sliding deformation involves dislocation propagation, and a region with a high dislocation density is formed near the shear stage. With the increase in the amount of deformation applied to the steel sheet, the slipping deformation is promoted, and the dislocation density increases.

In cold forging, steel processing in excess of one is carried out in considerable deformation. For this reason, in the conventional steel sheet, it is difficult to prevent the generation of voids and / or cracks accompanying the increase in dislocation density, and it is difficult to improve cold forging.

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

Carbide in the steel sheet is a hard particle that prevents slipping, and the presence of carbide at the ferrite grain boundary suppresses the formation of the shear zone and improves cold forging.

In order to obtain such an effect, the carbide needs to be dispersed in an appropriate size in the metal structure. Therefore, the average particle diameter of carbide is made into 0.4 micrometer or more and 2.0 micrometers or less. If the particle diameter of the carbide is less than 0.4 µm, the hardness of the steel sheet is significantly increased, and the cold forging property is lowered. More preferably, it is 0.6 micrometer or more.

On the other hand, when the average particle diameter of carbide exceeds 2.0 micrometers, carbide will become a starting point of a crack at the time of cold forming. More preferably, it is 1.95 micrometers or less.

In addition, cementite, which is a carbide of iron, is a hard and brittle structure. When steel is present in the state of pearlite, which is a layered structure with ferrite, steel is hard and brittle. Therefore, it is necessary to make pearlite as small as possible, and in the steel plate of this invention, you may be 6% or less by area ratio.

Since pearlite has a unique lamellar structure, it can be distinguished strictly by SEM and optical microscope observation. The area ratio of pearlite can be calculated | required by calculating the area | region of a lamellar structure in arbitrary cross sections.

Based on the theory and principle, cold forging is considered to be strongly influenced by the coverage of carbides of ferrite grain boundaries, and the measurement of the high precision is required, but the coverage of carbides to ferrite grain boundaries in three-dimensional space is required. In the scanning electron microscope, serial section SEM observation, or three-dimensional EBSP observation, which repeatedly performs sample cutting and observation by FIB in a scanning electron microscope, becomes essential, requires a large measurement time, and accumulates technical know-how. This is indispensable.

The inventors made this clear and searched for a simpler and more accurate evaluation index. As a result, when the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains was used as an index, the cold forging was evaluated. When the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the particles exceeds 1, the inventors have found that the cold forging property is remarkably improved.

In addition, since buckling, bending and folding of the steel sheet, which occur during cold forging, are all caused by localization of the deformation associated with the formation of the shear zone, similarly, carbides are present at the ferrite grain boundaries, thereby localizing the formation and deformation of the shear zone. When relaxed, the occurrence of buckling, bending and folding can be suppressed.

Carbide observation is performed with a scanning electron microscope. Prior to the 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 μm, and the observation surface was mirror-finished, and then the tissue was saturated with saturated picric acid-alcohol solution. Etch it.

The magnification of the observation was set to 3000 times, and eight images of 30 µm x 40 µm in a 1/4 layer thickness layer were taken at random. About the obtained tissue image, the area of each carbide contained in the area is measured in detail by image analysis software represented by Mitani Shoji Corporation (Win ROOF). From the area of each carbide, a circle equivalent diameter [= 2 × √ (area / 3.14)] is obtained, and the average value is a carbide particle diameter.

In addition, in order to suppress the influence of the measurement error by noise, carbide whose area is 0.01 micrometer <^2> or less is excluded from evaluation object.

The number of carbides present on the ferrite grain boundary is counted, the number of carbides on the grain boundary is subtracted from the total carbide number, and the number of carbides in the ferrite grains is obtained. Based on the measured number, the ratio of the number of grain boundaries carbide to the carbide in the ferrite particles is calculated.

As a structure after annealing, cold forging property can be improved by making a ferrite particle diameter into 3.0 micrometers or more and 50.0 micrometers or less. If the ferrite grain size is less than 3 µm, the hardness increases, and cracks and cracks tend to occur during cold forging, and therefore the ferrite grain size is preferably 3.0 µm or more. More preferably, it is 7.5 micrometers or more.

On the other hand, when the ferrite grain size exceeds 50.0 µm, the number of carbides on the grain boundary that suppresses the propagation of slip decreases and the cold forging property is lowered. Therefore, the ferrite grain size is preferably 50.0 µm or less. More preferably, it is 37.9 micrometers or less.

The ferrite particle size is the procedure described above, and after polishing the observation surface of the sample for tissue observation with a mirror surface, the structure of the observation surface etched with a 3% nitric acid-alcohol solution is observed with an optical microscope or a scanning electron microscope, and the photographing is performed. It measures by applying the line segment method to one image.

By setting the Vickers hardness of the steel sheet to 100 HV or more and 180 HV or less, the cold forging property and the impact resistance property after carburizing-quenching tempering can be improved. If the Vickers hardness is less than 100 HV, buckling tends to occur during cold forging, bending and folding of the buckling portion occur, and the impact resistance is lowered. Therefore, the Vickers hardness is set to 100 HV or more. Preferably it is 110 HV or more.

On the other hand, when the Vickers hardness exceeds 180 HV, the ductility decreases, and internal cracking tends to occur during cold forging, and the impact resistance deteriorates. Therefore, the Vickers hardness is set to 180 HV or less. Preferably it is 170 HV or less.

Next, the evaluation method of cold forging is demonstrated.

The outline of the cold forging test and the form of the crack introduced in cold forging are shown typically in FIG. Fig. 1 (a) shows a disk shaped test material cut out from a hot rolled steel sheet, Fig. 1 (b) shows the shape of the test material after cold forging, and Fig. 1 (c) shows the cross-sectional shape of the test material after cold forging. Shows.

As shown in FIG. 1, the disk shaped test material 1 of diameter 70mm was cut out from the hot rolled sheet steel of plate thickness 5.2mm (refer FIG.1 (a)), and the diameter of a bottom face is 30 mm by deep drawing. An in-cup type test piece is manufactured (not shown). Next, using the Moriteco one-shot forming press machine, the vertical wall portion of the cup-shaped test material was thickened and molded (cold forging) at a thickness increase ratio of 1.54 (= 8 mm / 5.2 mm), and the diameter was 30 mm, the height was 30 mm, and the vertical wall. A cup-shaped test material 2 having a thickness of 8 mm was produced (see FIG. 1B).

The cup-shaped test material 2 subjected to the thickness increasing molding is cut so that a cross section of the diameter portion appears in a wire cut electric discharge machine made by FANUC (see Fig. 1 (c)). Mirror-polishing the cut surface, confirming that the cracks 3 were present in the cut surface, and the ratio of the maximum length L of cracks present in the vertical wall portion to the thickness of the vertical wall portion after thickness increase (= maximum length L of cracks after thickness increase 8 mm of thickness of vertical wall part) is measured. Based on this measured value, cold forging property is evaluated.

In addition, even if the initial plate thickness is other than 5.2 mm, if the diameter of the disk-shaped test specimen to be cut out is adjusted so that the height of the vertical wall after the thickness increase is 30 mm, and molded at the same thickness increase ratio of 1.54, it is not dependent on the initial plate thickness. In addition, since an evaluation result can be reproduced, the hot-rolled steel sheet which this invention targets is not limited to the hot-rolled steel sheet of 5.2 mm of plate | board thickness. The present invention can improve the impact resistance after cold forging and carburizing quenching tempering even in a hot rolled steel sheet having a general sheet thickness (2 to 15 mm).

Next, the manufacturing method of this invention is demonstrated. The technical idea of the manufacturing method of this invention is to manage hot rolling conditions and annealing conditions consistently when manufacturing a steel plate from the steel piece of the above-mentioned component composition, and to improve the impact resistance after cold forging and carburizing-quenching tempering.

The characteristic of the manufacturing method of this invention is demonstrated.

[Features of Hot Rolling]

Continuously casting molten steel having the required component composition into a slab, and applying the slab to hot rolling as it is in the usual manner, or heating after cooling once, and providing to the hot rolling, a temperature of 650 ° C or more and 950 ° C or less Finish finishing hot rolling in the area. The hot rolled steel sheet after finishing rolling is cooled on ROT, and wound up at a coiling temperature of 400 ° C. or higher and 600 ° C. or lower.

[Features of annealing]

After the pickling, the hot rolled steel sheet is subjected to an annealing of a two-stage step type held in two temperature ranges. However, in the first stage annealing, the hot rolled steel sheet is 30 ° C / hour or more to 150 ° C / hour to the annealing temperature. It heats at a heating rate and performs annealing hold | maintains 3 hours or more and 60 hours or less in the temperature range of 650 degreeC or more and 720 degrees C or less.

In the second stage of annealing, the hot rolled steel sheet is heated to an annealing temperature at a heating rate of 1 ° C./hour or more and 80 ° C./hour or less, and held for 3 hours or more and 50 hours or less in a temperature range of 725 ° C. or more and 790 ° C. or less. Annealing is performed.

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

By cooperation of this hot rolling condition and annealing conditions, the low carbon steel plate excellent in cold forging property and impact resistance after carburizing-quenching tempering can be obtained.

Below, the process conditions of the manufacturing method of this invention are demonstrated concretely.

[Hot rolling]

Finishing hot rolled temperature: 650 degrees Celsius or more and 950 degrees Celsius or less

Winding temperature: 400 ° C or more and 600 ° C or less

Continuously casting molten steel having the required component composition into a slab, and heating it as it is or after cooling once, providing it to hot rolling, finishing finishing hot rolling in a temperature range of 650 ° C or more and 950 ° C or less, and hot-rolled steel sheet of 400 ° C or more. It winds up at 600 degrees C or less.

The slab heating temperature is preferably 1300 ° C. or less, and the heating time for which the temperature of the slab surface layer is maintained at 1000 ° C. or more is preferably 7 hours or less.

If the heating temperature exceeds 1300 ° C. or the heating time exceeds 7 hours, the decarburization of the slab surface layer becomes remarkable, and the austenite particles in the surface layer grow abnormally at the time of heating before quenching, and the impact resistance property is deteriorated. 1300 degrees C or less is preferable, and, as for heating temperature, 7 hours or less are preferable. More preferably, heating temperature is 1280 degrees C or less, and heating time is 6 hours or less.

Finishing hot rolling is finished at the temperature of 650 degreeC or more and 950 degrees C or less. When the finish hot rolling temperature is less than 650 ° C, the rolling load is significantly increased from the increase in the deformation resistance of the steel, the roll wear amount is increased, and the productivity is lowered. Therefore, the finish hot rolling temperature is at least 650 ° C. Preferably it is 680 degreeC or more.

On the other hand, if the finish hot rolling temperature exceeds 950 ° C, a thick scale is produced during the passage through the Run Out Table (ROT), scratches occur on the surface of the steel sheet due to the scale, and during cold forging and / or carburizing quenching and tempering Later, when the impact load is applied, cracks are generated from scratches and the impact resistance is lowered. Therefore, the finish hot rolling temperature is 950 ° C or lower. Preferably it is 920 degrees C or less.

As for the cooling rate at the time of cooling a hot rolled sheet steel on ROT, 10 degrees C / sec or more and 100 degrees C / sec or less are preferable. If the cooling rate is less than 10 ° C / sec, the generation of thick scales and generation of scratches due to it cannot be suppressed during cooling, and the impact resistance is lowered, so the cooling rate is preferably 10 ° C / sec or more. . More preferably, it is 20 degree-C / sec or more.

On the other hand, when the hot rolled steel sheet is cooled from the surface layer of the steel sheet at a cooling rate exceeding 100 ° C / sec, the outermost layer portion is excessively cooled, and low-temperature transformation structures such as bainite and martensite are generated in the outermost layer portion. do.

When winding up the hot rolled steel sheet at 100 ° C to room temperature after winding, microcracks occur in the low temperature transformation structure, and it is difficult to remove the cracks in the subsequent pickling process and cold rolling process, during cold forging and / or carburizing quenching tempering. When a shock load is applied afterwards, since a crack advances from a crack and a fall of impact resistance property is caused, a cooling rate of 100 degrees C / sec or less is preferable. More preferably, it is 80 degrees C / sec or less.

In addition, the cooling rate is from the cooling facility of each casting section from the time when the hot-rolled steel sheet after the finish hot rolling passes through the water-free section, from the time of receiving water cooling in the pouring section, to the target temperature of winding up on the ROT The cooling ability to be received is indicated, and the average cooling rate from the water injection start point to the temperature wound by the winding machine is not shown.

Winding temperature shall be 400 degreeC or more and 600 degrees C or less. This is a temperature lower than the general winding temperature. By winding up the hot rolled sheet steel manufactured on the conditions mentioned above in this temperature range, the structure of a steel plate can be made into the bainite structure in which carbide disperse | distributed in the fine ferrite.

If the coiling temperature is less than 400 ° C, the austenite that was untransformed before winding is transformed into hard martensite, and when discharging the wound hot rolled steel sheet, cracks occur in the surface layer, and the impact resistance is lowered, so the coiling temperature is 400 ° C. Do as above. Preferably it is 430 degreeC or more.

On the other hand, when the coiling temperature exceeds 600 ° C, pearlite having a large lamellar spacing is formed, thick needle-like carbides having high thermal stability are formed, and needle-like carbides remain even after the annealing of the two-stage step type. At the time of cold forging, since a crack generate | occur | produces and advances on this needle-like carbide as a starting point, a coiling temperature shall be 600 degrees C or less. Preferably it is 570 degreeC or less.

After pickling, the hot rolled steel sheet manufactured under the above conditions is subjected to an annealing of a two-stage step type held in two temperature ranges. By performing a two-step step annealing on the hot-rolled steel sheet, the stability of the carbide is controlled and the formation of carbide to the ferrite grain boundary is promoted.

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

By carrying out the first stage annealing in the temperature range below Ac1 point, the carbides are coarsened, the added metal element is concentrated, and the thermal stability of the carbides is increased. Thereafter, the temperature is raised to a temperature range of Ac1 or more, austenite is formed in the structure, carbides in fine ferrite particles are dissolved in austenite, and coarse carbide remains in austenite.

By subsequent slow cooling, austenite is transformed into ferrite and the carbon concentration in austenite is increased. By advancing slow cooling, carbon atoms adsorb | suck to the carbide which remain | survives in austenite, carbide and austenite will cover the grain boundary of a ferrite, and finally it becomes possible to form the structure by which a large amount of carbide exists in a ferrite grain boundary. Therefore, it is clear that the tissue defined by the present invention cannot be formed only by simple annealing.

Below, specific annealing conditions are demonstrated.

[1st stage annealing]

Heating rate up to annealing temperature: 30 ° C / hour or more and 150 ° C / hour

Annealing Temperature: Above 650 ℃, Below 720 ℃

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

The heating rate to the annealing temperature of the first stage is made 30 ° C / hour or more and 150 ° C / hour or less. If a heating rate is less than 30 degreeC / hour, time will require time for temperature rising and productivity will fall, and heating rate shall be 30 degreeC / hour or more. Preferably it is 40 degreeC / hour or more.

On the other hand, when the heating rate exceeds 150 ° C / hour, the temperature difference between the outer circumference of the coil and the inside increases, and friction marks and sintering occur due to the difference in thermal expansion, and irregularities are generated on the surface of the steel sheet. During cold forging, cracks occur from this unevenness as a starting point, and thus the cold forging property is lowered and the impact resistance after carburizing and quenching tempering is caused. Therefore, the heating rate is 150 ° C./hour or less. Preferably it is 120 degrees C / hour or less.

Annealing temperature (annealing temperature of 1st stage) in annealing of a 1st stage | paragraph shall be 650 degreeC or more and 720 degrees C or less. When the first stage annealing temperature is less than 650 ° C., the stability of the carbide is insufficient, and in the second stage annealing, it becomes difficult to leave the carbide in the austenite, so the first stage annealing temperature is set to 650 ° C. or higher. Preferably it is 670 degreeC or more.

On the other hand, when the annealing temperature exceeds 720 ° C, the austenite is produced before the stability of the carbide increases, and thus the structure change described above cannot be controlled, so the annealing temperature is set to 720 ° C or lower. Preferably it is 700 degrees C or less.

The holding time (stage 1 holding time) in annealing of a 1st stage | paragraph shall be 3 hours or more and 60 hours or less. If the first-stage holding time is less than 3 hours, the stabilization of the carbides is not sufficient, and it becomes difficult to leave the carbides in the second-stage annealing. Therefore, the first-stage holding time is set to 3 hours or more. Preferably it is 10 hours or more.

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

[2nd stage annealing]

Heating rate up to annealing temperature: 1 ° C / hour or more and 80 ° 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 the annealing temperature at a heating rate of 1 ° C./hour or more and 80 ° C./hour or less. In the case of cooling without performing the second stage annealing, the ferrite grain size does not increase, and an ideal structure cannot be obtained.

In the second stage of annealing, austenite is formed and grown from the ferrite grain boundary. By slowing down the heating rate, nucleation of austenite can be suppressed, and the grain boundary coverage of the carbide can be increased in the structure obtained after slow cooling. Therefore, the smaller the heating rate in the second stage annealing is preferable.

If a heating rate is less than 1 degree-C / hour, since time will require time for temperature rising and productivity will fall, heating rate shall be 1 degree-C / hour or more. Preferably it is 10 degreeC / hour or more.

On the other hand, when the heating rate exceeds 80 deg. C / hour, the temperature difference between the outer circumference of the coil and the inside increases, and friction marks and sintering occur due to the large difference in thermal expansion due to transformation, and irregularities are generated on the surface of the steel sheet. During cold forging, cracks are formed from this unevenness as a starting point, and the cold forging property is lowered and the impact resistance after carburizing and quenching tempering is caused, so that the heating rate is 80 ° C./hour or less.

The annealing temperature (annealing temperature of the 2nd stage) in the 2nd stage annealing shall be 725 degreeC or more and 790 degrees C or less. When the annealing temperature of the second stage is less than 725 ° C, the amount of austenite produced decreases, and after cooling after the second stage of annealing, the number ratio of carbides on the ferrite grain boundary decreases, and the ferrite grain size decreases. For this reason, annealing temperature of a 2nd stage | paragraph shall be 725 degreeC or more. Preferably it is 735 degreeC or more.

On the other hand, when the annealing temperature of the second stage exceeds 790 ° C, it becomes difficult to remain carbide in austenite, and it becomes difficult to control the above-described structure change. Therefore, the annealing temperature of the second stage is set to 790 ° C or lower. Preferably it is 780 degrees C or less.

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

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

[Cooling after annealing]

Cooling stop temperature: 650 ℃

Cooling rate: 1 ℃ / hour or more and 100 ℃ / hour or less

After the holding | maintenance in the 2nd stage annealing is complete | finished, the hot rolled sheet steel after annealing is cooled slowly to the cooling rate of 1 degree-C / hour or more and 100 degrees C / hour or less to 650 degreeC. When the stop temperature of the slow cooling exceeds 650 ° C, the unmodified austenite transforms to pearlite or bainite by the cooling rate exceeding 100 ° C / hour to the subsequent room temperature, the hardness increases, and the cold end Since a composition falls, cooling stop temperature shall be 650 degreeC.

In order to cool the austenite produced in the second stage annealing to transform it into ferrite and to adsorb carbon with the carbide remaining in the austenite, the cooling rate is preferably slower. If cooling rate is less than 1 degree-C / hour, since time required for cooling will increase and productivity will fall, cooling rate shall be 1 degree-C / hour or more. Preferably it is 10 degreeC / hour or more.

On the other hand, if the cooling rate exceeds 100 DEG C / hour, the austenite transforms into pearlite, the hardness of the steel sheet increases, leading to a decrease in cold forging properties and a drop in impact resistance after carburizing and quenching tempering. It may be 100 degrees C / hour or less. Preferably it is 90 degreeC / hour.

Here, cooling stop temperature is temperature which should be controlled by said cooling rate, and when cooling to 650 degreeC is performed at cooling rate 1 degreeC / hour or more and 100 degrees C / hour or less, it will restrict | limit especially about cooling to 650 degrees C or less. It doesn't work.

In addition, the atmosphere of annealing is not limited to a specific atmosphere. For example, any of an atmosphere of 95% or more of nitrogen, an atmosphere of 95% or more of hydrogen, and an atmospheric atmosphere may be used.

As explained above, according to the manufacturing method which manages the hot-rolling conditions and annealing conditions of this invention consistently, and controls the structure of a steel plate, it exhibits the outstanding cold forging property in the cold forging which combined drawing and thickness increasing molding, A low carbon steel sheet excellent in impact resistance after carburizing and quenching tempering can be produced.

Example

Next, although an Example is described, the level of an Example is an example of the execution conditions employ | adopted in order to confirm the feasibility and effect of this invention, and this invention is not limited to this one example of conditions. This invention can employ | adopt various conditions, as long as the objective of this invention is reached, without deviating from the summary of this invention.

The continuous casting cast piece (steel ingot) which has the component composition shown in Table 1 was heated at 1240 degreeC for 1.8 hours, and was used for hot rolling. The finish hot rolling was complete | finished at 890 degreeC, it cooled to 520 degreeC at the cooling rate of 45 degree-C / sec on ROT, and wound up at 510 degreeC, and produced the hot rolled coil of 5.2 mm of sheet thickness.

After pickling the hot rolled coil, charging the coil into a box-annealing furnace, controlling the atmosphere to 95% hydrogen-5% nitrogen, and then heating from room temperature to 705 ° C at a heating rate of 100 ° C / hour, and at 705 ° C. It maintained for 36 hours and uniformized the temperature distribution in a coil. Thereafter, heating was performed at 760 ° C. at a heating rate of 5 ° C./hour, held at 760 ° C. for 10 hours, and then cooled up to 650 ° C. at a cooling rate of 10 ° C./hour, after which the furnace was cooled to room temperature to provide characteristics. The sample for evaluation was produced.

The structure of the sample was observed by the method mentioned above, and the crack length which exists in the sample after cold forging was measured by the method mentioned above.

Carburization of the sample with increasing thickness was carried out by gas carburization. In order to diffuse carbon into the steel from the atmosphere gas in the furnace through the sample surface layer, a treatment was performed at 940 ° C. for 120 minutes in a furnace controlled at a carbon potential of 0.5% by mass C, and thereafter the furnace was cooled to room temperature. did.

Subsequently, after heating from room temperature to 840 degreeC, it hold | maintained for 20 minutes and quenched in 60 degreeC oil. The quenched sample was subjected to air cooling tempering treatment after holding at 170 ° C. for 60 minutes to produce a carburized quenching tempering sample.

The impact resistance of the carburized quenched tempering sample was evaluated by the drop test. The outline of the drop test which evaluates the impact resistance characteristic of the sample which performed carburizing quenching tempering is shown typically in FIG. The cup bottom of the cup-shaped sample 4 subjected to carburizing and quenching tempering was fixed with a jig, and a weight of 2 kg was dropped on the side of the cup (upper width: 50 mm, lower width: 10 mm, height: 80 mm, length: 110). Mm) is free-falling from the upper part spaced 4m, the impact of about 80J is given to the vertical wall part of the sample 4, the presence or absence of the crack of a sample is observed, and an impact resistance characteristic is evaluated.

As a result of the free fall, a sample showing no cracks or breakage was given a rating of "OK" excellent in impact resistance characteristics, and a sample showing cracks or breakage was rated "NG" having a low impact resistance.

Table 2 shows the maximum crack lengths of carbide diameter, pearlite area ratio, ferrite grain size, Vickers hardness, ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains, and the plate thickness of the vertical wall in the manufactured samples. The measurement result and evaluation result of ratio and impact resistance of are shown.

As shown in Table 2, invention steels A-1, B-1, C-1, D-1, E-1, F-1, G-1, H-1, I-1, J-1 and K The ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite particles is greater than -1, and all of -1 has a Vickers hardness of 100 HV or more and 180 HV or less, and is excellent in cold forging properties and impact resistance after carburizing quenching and tempering.

On the other hand, comparative steel L-1 has a low amount of C and has a low cold forging property because the hardness before cold forging is less than 100 HV. Comparative steels M-1, P-1, and Z-1 contain excessive amounts of P, Al, and N, and in the second stage of annealing, the amount of segregation to the γ / α interface is large, so that carbides are formed at grain boundaries. This is suppressed.

Comparative steel S-1 contains excessively Si and has low ductility of ferrite, and thus has low cold forging. Comparative steels N-1 and T-1 each contain excessive amounts of Mo and Cr, and carbides are finely dispersed in the ferrite particles, and the hardness exceeds 180 HV. Since comparative steel Q-1 contains excessively Mn, the impact resistance after carburizing-quenching tempering is remarkably low.

Comparative steel O-1 is low in Cr and has low impact resistance because the austenite particles in the surface layer were abnormally coarsened during carburization. Since comparative steel R-1 contains S excessively, coarse MnS is produced | generated in steel, and cold forging property is low. Since the comparative steel U-1 contains an excess of C, coarse carbides are formed inside the thickness increase of the steel, and coarse carbides remain even after carburizing and quenching, and thus the impact resistance is low.

Since comparative steel V-1 had little Mn amount and it was difficult to raise the stability of carbide, it has low cold forging property and the impact resistance after carburizing quenching tempering. Comparative steels W-1 and X-1 contain excessive amounts of O and Ti, so that oxides and TiCs present in the ferrite particles become carbide formation sites in slow cooling after two-phase region annealing, and carbides at grain boundaries. Is suppressed and cold forging is low. Comparative steel Y-1 contains B excessively and is low in cold forging.

Subsequently, in order to investigate the influence of manufacturing conditions, the slab which has the component composition of A, B, C, D, E, F, G, H, I, J, and K shown in Table 1 is shown in Table 3. And annealing conditions, a hot rolled sheet annealing sample having a plate thickness of 5.2 mm was produced.

Table 4 shows the ratios of carbide diameter, pearlite area ratio, ferrite grain size, Vickers hardness, ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains, and the maximum crack length with respect to the plate thickness of the vertical wall, for the prepared samples. The measurement result and evaluation result of ratio and impact resistance are shown.

Comparative steel E-3 has a low finish hot rolling temperature and an increased rolling load, resulting in low productivity. Comparative steel D-2 had a high finish hot rolling temperature and produced scale marks on the surface of the steel sheet. Therefore, when applied to the abrasion resistance test after quenching tempering, cracks and peeling occurred due to the scale marks as a starting point, and the wear resistance characteristics were improved. Degraded. Comparative steel F-2 has a low cooling rate in the ROT (Run Out Table), which causes a decrease in productivity and generation of scale marks.

In Comparative Steel C-4, the cooling rate in ROT was 100 ° C./sec, and microcracks were generated in the outermost layer portion because the outermost layer portion of the steel sheet was excessively cooled. Comparative steel C-2 has a low coiling temperature, generates a lot of low-temperature transformation structures such as bainite and martensite, is embrittled, cracks frequently occur during hot rolled coil discharging, and productivity is lowered. Moreover, the wear resistance in the sample collected from the cracked piece is low.

Comparative steel G-2 has a high coiling temperature, a pearlite having a large lamellar spacing in the hot rolled structure, high thermal stability of acicular coarse carbide, and even after two-stage step annealing. Since it remains in the middle, the machinability is low. Comparative steel H-4 has low productivity since the heating rate in the first stage of annealing of the two-step step type annealing is low.

In Comparative Steel E-3, the heating rate in the first stage of annealing is fast, so that the temperature difference between the inside of the coil and the inner and outer circumferences increases, and friction marks and sizing due to the thermal expansion difference occur, and wear resistance characteristics after quenching tempering. When it provided for the evaluation test of, the crack and peeling generate | occur | produced from the collection part and the wear-resistant characteristic fell.

Comparative steel G-4 has a low holding temperature (annealing temperature) in the first stage annealing, insufficient coarsening treatment of carbides at the Ac1 point or less, and insufficient thermal stability of the second stage. The amount of carbides remaining in the wafer decreases and the pearlite transformation cannot be suppressed in the structure after slow cooling, resulting in low machinability.

Comparative steel D-4 has a high holding temperature (annealing temperature) in the first stage of annealing, austenite is formed during the annealing, and stability of carbide cannot be improved. Thus, pearlite is produced after annealing, and Vickers hardness is 180 HV. In addition, the machinability is low. Comparative steel J-4 has a short holding time in the first stage annealing, cannot improve the stability of the carbide, and is low in machinability.

Comparative steel F-2 had a long holding time in the first stage annealing, had low productivity, had a scuffing mark, and had low abrasion resistance. Comparative steel B-4 has low productivity because the heating rate in the second stage annealing of the two-step step annealing is slow. In Comparative Steel A-3, the heating rate in the second stage of annealing is fast, so that the temperature difference between the inside and the outer circumference of the coil is large, and friction marks and sizing due to the large thermal expansion difference due to transformation occur, resulting in quenching and tempering. Low wear resistance

Comparative steel K-2 has a low holding temperature (annealing temperature) in the second stage annealing, a small amount of austenite production and a low number of carbides in the ferrite grain boundary. Comparative steel C-4 had a high holding temperature (annealing temperature) in the second stage of annealing and promoted dissolution of carbides during annealing, making it difficult to form grain boundary carbides after slow cooling, and pearlite was formed. Hardness exceeds 180 HV, and machinability is low.

Comparative steel J-3 has a long holding time in the second stage of annealing and has a low machinability since the dissolution of carbides is promoted. Comparative steel D-3 has a slow cooling rate from annealing at the second stage to 650 ° C, low productivity, coarse carbides are formed in the structure after slow cooling, and cracks are formed based on coarse carbides during cold forging. Generated, and cold forging was reduced. Comparative steel I-3 has a low cooling forging since the cooling rate from the second stage annealing to 650 ° C. is high, and a pearlite transformation occurs during cooling to increase the hardness.

Subsequently, in order to investigate the allowable content of other elements, the continuous casting cast piece (steel ingot) which has the component composition shown in Table 5 and Table 6 (following Table 5) was heated at 1240 degreeC for 1.8 hours, and was then used for hot rolling. . The finish hot rolling was complete | finished at 890 degreeC, it cooled to 520 degreeC at the cooling rate of 45 degree-C / sec on ROT, and wound up at 510 degreeC, and produced the hot rolled coil of 5.2 mm of sheet thickness.

The hot rolled coil was pickled, the hot rolled coil was charged into a box annealing furnace, the atmosphere was controlled with 95% hydrogen-5% nitrogen, and then heated from room temperature to 705 ° C. at a heating rate of 100 ° C./hour, and 705 ° C. The temperature distribution in the coil was uniformed for 36 hours at, followed by heating to 760 ° C. at a heating rate of 5 ° C./hour, and holding at 760 ° C. for 10 hours, and then cooling to 10 ° C./hour at 650 ° C. After cooling at a rate, the furnace was cooled to room temperature and a sample for characteristic evaluation was produced.

In addition, the structure of the sample was observed by the method mentioned above, and the crack length which exists in the sample after cold forging was measured by the method mentioned above.

Table 7 shows the maximum crack lengths of the carbide diameter, pearlite area ratio, ferrite grain size, Vickers hardness, ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains, and the plate thickness of the vertical wall in the manufactured samples. The measurement result and evaluation result of ratio and impact resistance of are shown.

As shown in Table 7, invention steels AA-1, AB-1, AC-1, AD-1, AE-1, AF-1, AG-1, AH-1, AI-1, AJ-1, AK -1, AL-1, AM-1, AN-1, AO-1, AP-1, and AQ-1 all have ratios of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite particles exceeding one; Vickers hardness is 100HV or more and 180HV or less, and is excellent in cold forging property and impact resistance property after carburizing-quenching tempering.

In comparison, the comparative steels AR-1, AS-1, AW-1, AZ-1, BB-1 and BF-1 each contain excessive amounts of La, As, Cu, Ni, Sb, and Ce. Since the amount of segregation to the γ / α interface increases at the time of annealing, generation of carbides at the grain boundary is suppressed. Comparative steel BG-1 contains excessively Si and is low in ductility of ferrite, and thus has low cold forging.

Comparative steels AT-1, AV-1, BA-1, BC-1, BH-1, and BJ-1 contain excessive amounts of Mo, Nb, Cr, Ta, W, and V, respectively, and therefore carbides in ferrite particles This fine dispersion and hardness exceed 180 HV. Since comparative steel BF-1 contains Mn excessively, the impact resistance after carburizing-quenching tempering is remarkably low.

Comparative steels AU-1, AX-1, AY-1 and BE-1 each contain excessive amounts of Zr, Ca, Mg and Y, coarse oxide or nonmetallic inclusions are formed in the steel, and are coarse at the time of cold forging. The crack generate | occur | produced starting from an oxide or a coarse nonmetallic interference | inclusion, and cold forging property fell. Comparative steel BD-1 contains Sn excessively, ferrite is embrittled, and cold forging is low. Since comparative steel BK-1 contained C excessively, coarse carbide was produced inside the thickness increase of steel, coarse carbide remained after carburizing quenching, and the impact resistance property fell.

Subsequently, in order to investigate the influence of the manufacturing conditions, the AA, AB, AC, AD, AE, AF, AG, AH, AI, AJ, AK, AL, AM, AN, AO, AP and AQ shown in Table 5 The hot rolled sheet annealing sample of the plate thickness of 5.2 mm was produced for the slab which has a component composition on the hot rolling conditions and annealing conditions shown in Table 8.

Table 9 shows the maximum crack lengths 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 boundary to the number of carbides in the ferrite grains, and the plate thickness of the vertical wall in the produced sample. The measurement result and evaluation result of ratio and impact resistance of are shown.

Comparative steel AC-2 has low finish hot rolling temperature and low productivity. Comparative steel AN-4 had a high finish hot rolling temperature, formed scale traces on the surface of the steel sheet, and when impact loads were applied after cold forging and carburizing quenching tempering, cracks were generated from scratches, and the impact resistance properties were deteriorated.

Inventive steel AB-3 has a slow cooling rate in ROT, resulting in a decrease in productivity and derivation of scale traces. Inventive steels AJ-3 and AD-4 had a cooling rate in ROT of 100 deg. C / sec, and the outermost layer portion of the steel sheet was excessively cooled to generate fine cracks in the outermost layer portion.

Comparative steel AN-3 had a low coiling temperature, produced a lot of low-temperature transformation structures such as bainite and martensite, was embrittled, cracked frequently during hot rolled coil delivery, and decreased productivity. In addition, the impact resistance characteristics after cold forging and carburizing quenching tempering in the samples collected from the cracked pieces were inferior.

Comparative steel AH-3 has a high winding temperature, a pearlite having a large lamellar spacing in the hot rolled structure, high thermal stability of acicular coarse carbide, and even after the two-stage step type annealing, the carbide is in the steel sheet. Since it remains, cold forging property is low.

Comparative steel AF-4 has low productivity due to a low heating rate in the first stage of annealing of the two-step step type annealing. In Comparative Steel AG-2, the heating rate in the first stage annealing is fast, so that the temperature difference between the inner and outer circumferences of the coil increases, friction marks and sizing due to the thermal expansion difference occur, and after cold forging and carburizing quenching tempering The impact resistance property fell.

Comparative steel AA-2 had a low holding temperature (annealing temperature) in the first stage annealing, insufficient coarsening treatment of carbides at the Ac1 point or less, insufficient thermal stability of the carbide, and annealing in the second stage. Remaining carbides decreased, and the pearlite transformation could not be suppressed in the structure after slow cooling, resulting in a decrease in cold forging.

Comparative steel AM-3 had a high holding temperature (annealing temperature) in the first stage, austenite was formed during annealing, the stability of carbides could not be improved, and cold forging properties and impact resistance after carburizing quenching and tempering were reduced. Comparative steel AF-2 has a short holding time in the first stage of annealing, cannot increase the stability of carbides, and is low in cold forging. Comparative steel AO-4 has a long holding time in the first stage annealing and has a low productivity.

Comparative steel AP-4 is low in productivity because the heating rate in the second stage of annealing of the two-stage step type annealing is slow. The comparative steel AI-3 has a high heating rate in the second stage annealing, so that the temperature difference between the coil inner and outer circumference becomes large, friction marks and sizing due to the large thermal expansion difference due to transformation occur, and after carburizing and quenching tempering. When the impact load was applied, cracks occurred from the scratches, and the impact resistance properties were lowered.

Comparative steel AL-3 had a low holding temperature (annealing temperature) in the second stage annealing, a small amount of austenite production, could not increase the number ratio of carbides in the ferrite grain boundary, and cold forging was deteriorated. . Comparative steel AD-2 had a high holding temperature (annealing temperature) in the second stage annealing and promoted dissolution of carbides during the annealing, making it difficult to produce grain boundary carbide after slow cooling, resulting in cold forging and carburizing. The impact resistance property after tempering fell.

Comparative steel AJ-4 has a low holding time in the second stage annealing and promotes dissolution of carbide, and thus has low cold forging. Comparative steel AQ-3 had a slow cooling rate from the second stage annealing to 650 ° C, low productivity, coarse carbides formed in the structure after slow cooling, and cracked with coarse carbides at the time of cold forging. This occurred and cold forging property fell. Comparative steel AP-2 had a fast cooling rate from the second stage annealing to 650 ° C, a pearlite transformation occurred at the time of cooling, and increased hardness, resulting in a decrease in cold forging.

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

From FIG. 3, when the number ratio (= number of grain boundary carbides / number of carbides in particles) exceeds 1, the ratio of crack length introduced in cold forging can be suppressed, and excellent impact resistance can be obtained after carburizing quenching and tempering. It can be seen that.

4 shows another relationship between the ratio of the number of grain boundary carbides to the number of carbides in the particle, the crack length of the cold forging test piece, and the impact resistance after carburizing quenching and tempering. It is a figure which shows that a crack length can be suppressed also in the steel plate which added the additional element.

From FIG. 4, even when the element of an appropriate range is added to the steel plate, when the number ratio (= number of grain boundary carbides / number of carbides in the particles) exceeds 1, the ratio of crack length introduced in cold forging is suppressed. It can be seen that excellent impact resistance can be obtained after carburizing and quenching tempering.

As described above, according to the present invention, it is possible to provide a low carbon steel sheet excellent in cold forging property and impact resistance after carburizing and quenching tempering and a manufacturing method thereof. The steel sheet of the present invention is suitable as a raw material for forming parts such as high cycle gears by forming by cold forging such as plate forming, for example, and thus the present invention has high industrial applicability.

1: disk-shaped test material
2: cup type test material
3: crack
4: sample
5: fall
L: maximum length of crack

Claims (3)

The component composition is in mass%
C: 0.10 to 0.40%,
Si: 0.01 to 0.30%,
Mn: 0.30 to 1.00%,
Al: 0.001-0.10%,
Cr: 0.50 to 2.00%,
Mo: 0.001-1.00%,
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: 0.10% or less,
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% or less
It is, and the remainder is a steel sheet which is Fe and impurities,
The metal structure of the steel sheet,
The pearlite area ratio is 6% or less, the balance consists of ferrite and carbides,
Carbide particle diameter is 0.4 to 2.0 mu m,
The particle size of the ferrite is 9.7 to 50.0 μm, and
The ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite particles is greater than one
To satisfy
Vickers hardness of the steel sheet is 100HV or more and 180HV or less
Steel sheet, characterized in that. The method of claim 1,
A steel sheet having a particle diameter of the ferrite of 10.2 to 50.0 μm. It is a manufacturing method of manufacturing the steel plate of Claim 1,
The steel piece of the component composition of Claim 1 is hot-rolled to complete finishing hot rolling in the temperature range of 650 degreeC or more and 950 degrees C or less, and it is set as a hot rolled sheet steel,
The hot rolled steel sheet is cooled at 10 ° C / sec or more and 100 ° C / sec or less on a ROT,
The hot rolled steel sheet is wound at 400 ° C or higher and 600 ° C or lower,
The pickled hot rolled steel sheet was pickled, and the pickled hot rolled steel sheet was heated at an annealing temperature of 650 ° C. or more and 720 ° C. or less at a heating rate of 30 ° C./hour or more and 150 ° C./hour or less, and maintained for 3 hours or more and 60 hours or less. We perform annealing of the first stage to say, and continue,
The hot rolled steel sheet is heated at an annealing temperature of 725 ° C or higher and 790 ° C or lower at a heating rate of 1 ° C / hour or more and 80 ° C / hour or less, and is subjected to the second stage of annealing to be maintained for 3 hours or more and 50 hours or less, followed by hot rolling. Cooling the steel sheet to 650 ° C at a cooling rate of 1 ° C / hour or more and 100 ° C / hour or less
The manufacturing method of the steel plate characterized by the above-mentioned.

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