JP6515332B2 - Low carbon steel sheet excellent in machinability and wear resistance after quenching and tempering, and method for producing the same - Google Patents

Description

  The present invention relates to a low carbon steel sheet which is excellent in machinability and wear resistance after quenching and tempering, and a method for producing the same.

  A low carbon steel sheet containing 0.1 to 0.4% of carbon by mass% is formed by press forming, hole expansion forming, bending forming, draw forming, thickness increase and thickness reduction formation, and a combination thereof in a blank material It is cold forged and used as a material to manufacture drive system parts such as gears and clutches of automobiles.

  Since it is difficult to form the final shape of the part with high precision only by cold forging, cutting is performed after cold forging. Further, in order to suppress the wear due to the mutual friction with the gear parts, which is generated when used as a drive system part of an automobile, the forming material is subjected to hardening and tempering to provide wear resistance.

  Until now, many proposals have been made on techniques for improving the machinability of low carbon steel sheets and the wear resistance characteristics after quenching and tempering (see, for example, Patent Documents 1 to 5).

  For example, in Patent Document 1, as a low-carbon sulfur-based free-cutting steel excellent in machinability, C: 0.05 to 0.15%, Mn: 0.5 to 2.0%, S: 0.1 to 0.4%, P: 0.05 to 0.10%, O: 0.005 to 0.040% as a basic component, and further, 0.1% or less of Si and 0.009 of Al There is disclosed a low carbon sulfur free-cutting steel which is limited to not more than 10%, contains N in the range of 20 to 150 ppm, further contains Ca in the range of 5 to 60 ppm, and the balance substantially consists of Fe. However, Patent Document 1 does not disclose any improvement in the wear resistance after heat treatment.

  Patent Document 2 discloses a general-purpose and high-strength alloy steel which is excellent in wear resistance and excellent in machinability. However, the steel of Patent Document 2 is a sintered steel using atomized iron powder, and has a problem that the price is higher than that of a general hot rolled steel sheet.

  Patent Document 3 solves the problem of not showing machinability at all by the cutting method and the problem of material deterioration, and further improves the chip treatability in deep drilling and significantly reduces the strength characteristics of the steel material. There is also disclosed a steel for machine construction for hot forging which has excellent machinability in a wide range of cutting methods and cutting conditions, and in particular, is excellent in carbide tool wear resistance and chip treatability.

  However, the steel of Patent Document 3 can not suppress an increase in environmental load and an increase in manufacturing cost due to the release of furnace heating gas in hot forging.

  In Patent Document 4, the fatigue strength of the non-hardened portion, the rolling resistance of the hardened portion, the pitting resistance, the resistance against the hardness of the non-hardened portion while minimizing the hardness increase after forging and securing the machinability and the cold workability. There is disclosed a steel material having improved wear resistance, fatigue strength and the like.

  However, in the steel material of Patent Document 4, the content of alloying elements is high, such as containing, by mass%, C: 0.40% or more and less than 0.50%, and further Si: 0.5 to 0.9%. Therefore, in the blast furnace integrated manufacturing of continuous casting-hot rolling-hot-rolled sheet annealing, the yield is lowered, and the increase in cost can not be suppressed.

  Patent Document 5 discloses a spheroidal graphite cast iron steel pipe forming roll having wear resistance to a forming surface and excellent in machinability when machining to form a forming surface. However, since the roll of Patent Document 5 is formed by a casting method, it has a problem of low productivity, and is not suitable for manufacturing widely used driving system parts of automobiles.

JP 07-252588 A Japanese Patent Application Publication No. 07-233402 JP 2000-87179 A JP 2004-149928 A JP, 2013-094823, A

  In view of the circumstances of the prior art, the present invention provides a low carbon steel sheet excellent in machinability and wear resistance characteristics after quenching and tempering, and a method for producing the same, in low cost integrated blast furnace manufacture of hot rolling and hot rolled sheet annealing. To be the subject.

  The present inventors diligently studied methods for solving the above problems. As a result, together with the S concentration in the steel, the dispersed state of carbides in the structure of the steel sheet before cutting is controlled by optimizing the manufacturing conditions in hot rolling to annealing, and carbides are mainly deposited on ferrite grain boundaries. And when the number of carbides on grain boundaries was increased more than the number of carbides in grains, it was found that excellent machinability could be ensured, and further the wear resistance after quenching and tempering could be improved.

  In addition, steel sheets based on the above findings are difficult to manufacture even if they are devised solely by hot rolling conditions, annealing conditions, etc., and process conditions are optimized in the continuous process of hot rolling-annealing process It was found through the accumulation of various researches that it could be manufactured.

  The present invention has been made based on the above findings, and the summary thereof 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%,
P: 0.0001 to 0.020%,
S: 0.010 to 0.10%,
Al: 0.001 to 0.10%,
N: 0.0001 to 0.020%,
O: 0.0001 to 0.020%,
Ti: 0.0001 to 0.010%,
B: 0 to 0.0005%
Containing the balance Fe and unavoidable impurities,
For carbides with an area of more than 0.01 μm ^2, the ratio of the number of carbide grain boundaries in the ferrite grain to the number of carbide grains in the ferrite grains exceeds 1,
Ferrite grain size is 5.0 μm or more,
A low carbon steel sheet excellent in machinability and wear resistance characteristics after quenching and tempering characterized in that Vickers hardness is 100 HV or more and 180 HV or less.

(2) The above-mentioned component composition is further in mass%,
Cr: 0.001 to 0.50%,
Mo: 0.001 to 0.50%,
Nb: 0.001 to 0.10%,
V: 0.001 to 0.10%,
Cu: 0.001 to 0.10%,
W: 0.001 to 0.10%,
Ta: 0.001 to 0.10%,
Ni: 0.001 to 0.10%,
Sn: 0.001 to 0.050%,
Sb: 0.001 to 0.050%,
As: 0.001 to 0.050%,
Mg: 0.0001 to 0.050%,
Ca: 0.001 to 0.050%,
Y: 0.001 to 0.050%,
Zr: 0.001 to 0.050%,
La: 0.001 to 0.050%,
The low carbon steel sheet excellent in the machinability and the wear resistance characteristics after quenching and tempering as described in the above (1), which comprises one or two or more species of Ce: 0.001 to 0.050%.

(3) A manufacturing method for manufacturing a low carbon steel sheet excellent in the machinability described in the above (1) or (2) and the wear resistance after quenching and tempering,
The steel piece of the component composition described in the above (1) or (2) is directly or temporarily cooled and then heated to be subjected to hot rolling, and finish hot rolling is completed in a temperature range of 650 ° C. to 950 ° C. When subjecting a hot-rolled steel sheet wound at 400 ° C. or more and 600 ° C. or less to two-step box-type annealing maintained in two temperature ranges after pickling,
(I-1) In the first stage annealing, the heat-rolled steel plate is heated to the annealing temperature at a heating rate of 30 ° C./hour or more and 150 ° C./hour or less, and in the temperature range of 650 ° C. or more and 720 ° C. or less Annealed to hold for less than 60 hours,
(I-2) In the second stage annealing, the heat-rolled steel plate is heated to the annealing temperature at a heating rate of 1 ° C./hour to 80 ° C./hour, and at a temperature range of 725 ° C. to 790 ° C. for 3 hours or more Apply annealing holding for less than 50 hours, then
(Ii) 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. The manufacturing method of the low carbon steel plate which is excellent in abrasion resistance characteristics.

  According to the present invention, it is possible to provide a low carbon steel sheet which is excellent in machinability and wear resistance after quenching and tempering, and a method for producing the same.

It is a figure which shows the relationship of the ratio of the number of carbides of a ferrite grain boundary to the number of carbides in a ferrite grain, and machinability. It is a figure which shows another relationship of the ratio of the number of the carbides of a ferrite grain boundary to the number of the carbides in a ferrite grain, and machinability. It is a figure which shows the relationship between S content in steel, and a wear-resistant characteristic. It is a figure which shows another relationship of S content in steel, and a wear-resistant characteristic.

The low carbon steel sheet (hereinafter sometimes referred to as “the steel sheet of the present invention”) excellent in the machinability and the wear resistance characteristics after quenching and tempering according to the present invention has a component composition in mass%,
C: 0.10 to 0.40%,
Si: 0.01 to 0.30%,
Mn: 0.30 to 1.00%,
P: 0.0001 to 0.020%,
S: 0.010 to 0.10%,
Al: 0.001 to 0.10%,
N: 0.0001 to 0.020%,
O: 0.0001 to 0.020%,
Ti: 0.0001 to 0.010%,
B: 0 to 0.0005%
Containing the balance Fe and unavoidable impurities,
For carbides with an area of more than 0.01 μm ^2, the ratio of the number of carbide grain boundaries in the ferrite grain to the number of carbide grains in the ferrite grains exceeds 1,
Ferrite grain size is 5.0 μm or more,
It is characterized in that Vickers hardness is 100 HV or more and 180 HV or less.

The method for producing a low carbon steel sheet excellent in the machinability and the wear resistance characteristics after quenching and tempering according to the present invention (hereinafter sometimes referred to as “the present invention producing method”) is a method for producing the steel sheet of the present invention
The billet of the component composition of the present invention steel sheet is heated directly or after cooling and heated, and subjected to hot rolling to complete finish hot rolling in a temperature range of 650 ° C. to 950 ° C., 400 ° C. to 600 ° C. When subjecting the hot-rolled steel sheet wound up below to two-step stepped box annealing maintained in two temperature ranges after pickling,
(I-1) In the first stage annealing, the heat-rolled steel plate is heated to the annealing temperature at a heating rate of 30 ° C./hour or more and 150 ° C./hour or less, and in the temperature range of 650 ° C. or more and 720 ° C. or less for 3 hours or more Annealed to hold for less than 60 hours,
(I-2) In the second stage annealing, the heat-rolled steel plate is heated to the annealing temperature at a heating rate of 1 ° C./hour to 80 ° C./hour, and at a temperature range of 725 ° C. to 790 ° C. for 3 hours or more Apply annealing holding for less than 50 hours, then
(Ii) The heat-rolled steel plate 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.

  Hereinafter, the steel sheet of the present invention and the manufacturing method of the present invention will be described.

  First, the reasons for limitation of the component composition of the steel sheet of the present invention will be described. Hereinafter, "%" which concerns on a component composition means "mass%."

C: 0.10 to 0.40%
C forms carbides in the steel and is an element effective for strengthening the steel and refining the ferrite grains. In order to suppress the generation of a matte surface in cold working and secure the surface appearance of a cold forged part, it is essential to suppress the coarsening of the ferrite grain size. If the amount is less than 0.10%, the volume fraction of carbides is insufficient, and coarsening of ferrite during box annealing can not be suppressed, so C is made 0.10% or more. Preferably, it is 0.11% or more.

  On the other hand, if it exceeds 0.40%, the volume fraction of carbide increases, and when a load is applied instantaneously, a large number of cracks serving as a starting point of fracture are generated, and the wear resistance is deteriorated. .40% or less. Preferably it is 0.38% or less.

Si: 0.01 to 0.30%
Si is an element which acts as a deoxidizing agent and influences the form of carbides. In order to reduce the number of carbides in the ferrite grains and to increase the number of carbides on the ferrite grain boundaries, an austenite phase was formed during annealing by two-step-step box annealing, and the carbides were once dissolved. After that, it is necessary to gradually cool and promote the formation of carbides at ferrite grain boundaries.

  If it exceeds 0.30%, the formation of a tribo film described later is suppressed, and the wear resistance after quenching and tempering is reduced, so Si is made 0.30% or less. Preferably it is 0.28% or less. The smaller the amount of Si, the better, but if it is reduced to less than 0.01%, the refining cost increases, so the amount of Si is made 0.01% or more. Preferably it is 0.04% or more.

Mn: 0.30 to 1.00%
Mn is an element that controls the form of carbides in the two-step box annealing. If it is less than 0.30%, it is difficult to form carbides on ferrite grain boundaries in slow cooling after the second stage annealing, so Mn is made 0.30% or more. Preferably it is 0.33% or more. On the other hand, if it exceeds 1.00%, the toughness after quenching and tempering will decrease, so the Mn is made 1.00% or less. Preferably it is 0.96% or less.

P: 0.0001 to 0.020%
P is an element which segregates in ferrite grain boundaries and suppresses the formation of grain boundary carbides. The smaller the amount, the better. However, if the purification is performed to less than 0.0001% in the refining process, the refining time will be long and the refining cost will be significantly increased, so P is made at least 0.0001%. Preferably, it is 0.0013% or more.

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

S: 0.010 to 0.10%
S is an element that greatly affects the machinability and wear resistance of the steel. Regarding the influence on the machinability, S is an element which is concentrated at the interface between ferrite and cementite to promote the exfoliation of the interface, and the machinability is improved by the effect of the interfacial exfoliation.

  As for the effect on wear resistance, S in the surface layer of parts in a wear resistant environment promotes the formation of a tribofilm at the lubricating oil / steel interface, and the hardness of the film is the concentration of S in the steel The higher the amount, the better the wear resistance.

  If it is less than 0.010%, the addition effect can not be sufficiently obtained, so S is made 0.010% or more. Preferably it is 0.015% or more. On the other hand, if it exceeds 0.10%, the formation of intergranular carbides is suppressed by the S significantly concentrated at the ferrite / austenite interface in the annealing step of the hot-rolled steel sheet described later, and the machinability is reduced. S is 0.10% or less. Preferably it is 0.08% 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 addition effect is not 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%, the number ratio of carbides on grain boundaries decreases and the machinability decreases, so Al is made 0.10% or less. Preferably it is 0.09% or less.

N: 0.0001 to 0.020%
N is an element which segregates in ferrite grain boundaries and suppresses the formation of carbides on the grain boundaries. The smaller the amount, the better. However, if it is reduced to less than 0.0001%, the refining cost is greatly increased, so N is made at least 0.0001%. Preferably it is 0.0006% or more.

  On the other hand, if it exceeds 0.020%, the ratio of the number of carbides on ferrite grain boundaries to the number of carbides in ferrite grains becomes 1 or less even if two-phase region annealing and gradual cooling is performed, and the machinability is lowered. Therefore, N is made 0.020% or less. Preferably it is 0.017% or less.

O: 0.0001 to 0.020%
O is an element that promotes the formation of oxides in the steel. The oxide present in the ferrite grains is the formation site of carbides, so less O is preferable, but if it is reduced to less than 0.0001%, the refining cost is significantly increased. Do. Preferably it is 0.0006% or more.

  On the other hand, if it exceeds 0.020%, the ratio of the number of carbides on ferrite grain boundaries to the number of carbides in ferrite grains becomes 1 or less, and the machinability decreases, so O is made 0.020% or less . Preferably it is 0.017% or less.

Ti: 0.0001 to 0.010%
Ti is an element important for controlling the form of carbides, and is an element that promotes the formation of carbides in ferrite grains with a large amount of addition. The smaller the amount of Ti, the better. However, if it is reduced to less than 0.0001%, the refining cost will increase significantly, so the amount of Ti should be at least 0.0001%. Preferably it is 0.0006% or more.

  On the other hand, if it exceeds 0.010%, the ratio of the number of carbides on ferrite grain boundaries to the number of carbides in ferrite grains becomes 1 or less, and the machinability decreases, so Ti is made 0.010% or less . Preferably it is 0.017% or less.

B: 0 to 0.0005%
B is an element which is concentrated at the interface between ferrite and carbide and suppresses segregation of S at the interface between ferrite and carbide. Although B is preferably as small as possible, it is necessary to pay close attention to analysis for identification of less than 0.0001%, and since the lower limit of detection is reached depending on the analysis apparatus, the lower limit is 0%. Preferably, it is at least 0.0001%.

  On the other hand, if the ratio exceeds 0.0005%, the ratio of the number of carbides on the ferrite grain boundary to the number of carbides in the ferrite grains is more than 1 even if the structure is optimized. Since the thickening is suppressed and the machinability does not improve, B is made 0.0005% or less. Preferably it is 0.0002% or less.

  The component composition of the steel plate of the present invention is Cr: 0.001 to 0.50%, Mo: 0.001 to 0.50% for the improvement of the machinability and other properties of the steel plate of the present invention in addition to the above elements. Nb: 0.001 to 0.10%, V: 0.001 to 0.10%, Cu: 0.001 to 0.10%, W: 0.001 to 0.10%, Ta: 0.001 ~ 0.10%, Ni: 0.001 to 0.10%, Sn: 0.001 to 0.050%, Sb: 0.001 to 0.050%, As: 0.001 to 0.050%, Mg: 0.0001 to 0.050%, Ca: 0.001 to 0.050%, Y: 0.001 to 0.050%, Zr: 0.001 to 0.050%, La: 0.001 to It may contain one or more of 0.050% and Ce: 0.001 to 0.050%.

Cr: 0.001 to 0.50%
Cr is an element effective for stabilizing carbides during two-phase region annealing. If it is less than 0.001%, the addition effect is not sufficiently obtained, so Cr is made 0.001% or more. Preferably, it is 0.010% or more. On the other hand, if it exceeds 0.50%, the formation of S-based tribofilms on the lubricating oil / steel surface is suppressed when the steel parts are cast in the abrading environment, and the wear resistance is lowered. Is less than 0.50%. Preferably it is 0.30% or less.

Mo: 0.001 to 0.50%
Mo, like Cr, is an element effective in stabilizing carbides during two-phase region annealing. If it is less than 0.001%, the addition effect is not sufficiently obtained, so Mo is made 0.001% or more. Preferably, it is 0.010% or more. On the other hand, if it exceeds 0.50%, the hardness of the carbide increases due to the concentration of Mo in the carbide and the machinability decreases, so the Mo content is made 0.50% or less. Preferably it is 0.30% or less.

Nb: 0.001 to 0.10%
Nb is an element effective for controlling the form of carbide, and by addition thereof, it is an element which refines the structure and contributes to the improvement of toughness. If it is less than 0.001%, the addition effect can not be sufficiently obtained, so Nb is made 0.001% or more. Preferably it is 0.002% or more.

  On the other hand, if it exceeds 0.10%, many fine and hard Nb carbides will precipitate, and the strength will increase too much, and the number ratio of intergranular carbides will decrease and the machinability will decrease. 10% or less. Preferably it is 0.09% or less.

V: 0.001 to 0.10%
V, like Nb, is also an element effective for shape control of carbides, and is an element contributing to the refinement of the structure and the improvement of toughness by the addition thereof. If it is less than 0.001%, the addition effect can not be sufficiently obtained, so V is made 0.001% or more. Preferably it is 0.004% or more.

  On the other hand, if it exceeds 0.10%, many fine V carbides are precipitated and the strength is increased too much, and the number ratio of intergranular carbides is decreased, so that the machinability is lowered, so V is 0.10% It is assumed that Preferably it is 0.09% or less.

Cu: 0.001 to 0.10%
Cu is an element which segregates in crystal grain boundaries of ferrite and is an element which forms fine precipitates and contributes to the improvement of strength. If it is less than 0.001%, the strength improvement effect can not be sufficiently obtained, so Cu is made 0.001% or more. Preferably it is 0.008% or more. On the other hand, if it exceeds 0.10%, it causes red shortness and productivity in hot rolling decreases, so Cu is made 0.10% or less. Preferably it is 0.09% or less.

W: 0.001 to 0.10%
W, like Nb and V, is also an element effective for controlling the form of carbide. If it is less than 0.001%, the addition effect is not sufficiently obtained, so W is made 0.001% or more. Preferably it is 0.003% or more. On the other hand, if it exceeds 0.10%, many fine W carbides are precipitated and the strength is increased too much, the number ratio of intergranular carbides is decreased, and the machinability is lowered, so W is 0.10%. It is assumed that Preferably it is 0.08% or less.

Ta: 0.001 to 0.10%
Ta, like Nb, V and W, is also an element effective for controlling the form of carbide. If it is less than 0.001%, the addition effect is not sufficiently obtained, so Ta is made 0.001% or more. Preferably it is 0.007% or more. On the other hand, if it exceeds 0.10%, many fine W carbides are precipitated and the strength is increased too much, the number ratio of intergranular carbides is decreased, and the machinability is lowered, so Ta is 0.10% It is assumed that Preferably it is 0.09% or less.

Ni: 0.001 to 0.10%
Ni is an element effective for improving the wear resistance of parts. If it is less than 0.001%, the addition effect is not sufficiently obtained, so Ni is made 0.001% or more. Preferably it is 0.002% or more. On the other hand, if it exceeds 0.10%, the number ratio of intergranular carbides decreases and the machinability decreases, so Ni is made 0.10% or less. Preferably it is 0.09% or less.

Sn: 0.001 to 0.050%
Sn is an element which mixes in steel from a steel raw material (scrap), segregates in grain boundaries, and reduces the number ratio of intergranular carbides. The smaller the amount, the better. However, if it is reduced to less than 0.001%, the refining cost will increase significantly, so Sn is made 0.001% or more. Preferably it is 0.002% or more. On the other hand, if it exceeds 0.050%, the ferrite becomes brittle and the wear resistance decreases, so Sn is made 0.050% or less. Preferably it is 0.048% or less.

Sb: 0.001 to 0.050%
Sb, like Sn, is an element which is mixed in from the steel material (scrap) into the steel and segregates at grain boundaries to reduce the number ratio of grain boundary carbides. The smaller the amount, the better, but if it is reduced to less than 0.001%, the refining cost will increase significantly, so Sb is made 0.001% or more. Preferably it is 0.002% or more. On the other hand, if it exceeds 0.050%, the ferrite becomes brittle and the wear resistance decreases, so Sb is made 0.050% or less. Preferably it is 0.048% or less.

As: 0.001 to 0.050%
As Sn and Sb, As is an element which is mixed into the steel from steel raw materials (scrap) and segregates in grain boundaries to reduce the number ratio of grain boundary carbides. The smaller amount is preferable, but the reduction to less than 0.001% significantly increases the refining cost, so As is made 0.001% or more. Preferably it is 0.002% or more. On the other hand, if it exceeds 0.050%, the number ratio of intergranular carbides decreases and the machinability decreases, so As is made 0.050% or less. Preferably it is 0.045% or less.

Mg: 0.0001 to 0.050%
Mg is an element that can control the form of sulfide by a small amount of addition. If it is less than 0.0001%, the addition effect can not be sufficiently obtained, so Mg is made 0.0001% or more. Preferably it is 0.0008% or more. On the other hand, if it exceeds 0.050%, coarse inclusions are formed and the wear resistance is reduced, so Mg is made 0.050% or less. Preferably it is 0.049% or less.

Ca: 0.001 to 0.050%
Ca, like Mg, is an element capable of controlling the form of sulfide with a small amount of addition. If it is less than 0.001%, the addition effect is not sufficiently obtained, so Ca is made 0.001% or more. Preferably it is 0.003% or more. On the other hand, if it exceeds 0.050%, coarse Ca oxide will be formed and it will be a starting point of cracking in the wear environment, so Ca should be 0.050% or less. Preferably it is 0.043% or less.

Y: 0.001 to 0.050%
Y, like Mg and Ca, is an element that can control the form of sulfide with a small amount of addition. If it is less than 0.001%, the addition effect is not sufficiently obtained, so Y is made 0.001% or more. Preferably it is 0.003% or more. On the other hand, when it exceeds 0.050%, coarse Y oxide is formed and the wear resistance is lowered, so Y is made 0.050% or less. Preferably it is 0.031% or less.

Zr: 0.001 to 0.050%
Zr, like Mg, Ca and Y, is an element that can control the form of sulfide with a small amount of addition. If it is less than 0.001%, the addition effect is not sufficiently obtained, so Zr is made 0.001% or more. Preferably it is 0.004% or more. On the other hand, if it exceeds 0.050%, coarse Zr oxide is formed and the wear resistance is lowered, so Zr is made 0.050% or less. Preferably it is 0.045% or less.

La: 0.001 to 0.050%
La is an element effective for controlling the morphology of sulfides with a small amount of addition, but it is also an element that segregates at grain boundaries and reduces the number ratio of grain boundary carbides. If it is less than 0.001%, the shape control effect can not be obtained sufficiently, so La is made 0.001% or more. Preferably it is 0.003% or more. On the other hand, if it exceeds 0.050%, the number ratio of intergranular carbides decreases and the machinability decreases, so La is made 0.050% or less. Preferably it is 0.047% or less.

Ce: 0.001 to 0.050%
Similar to La, Ce is an element that can control the form of sulfides with a small amount of addition, but is also an element that segregates at grain boundaries and reduces the number ratio of grain boundary carbides. If it is less than 0.001%, the shape control effect is not sufficiently obtained, so Ce is made 0.001% or more. Preferably it is 0.004% or more. On the other hand, if it exceeds 0.050%, the number ratio of intergranular carbides decreases and the machinability 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 unavoidable impurities.

  Hot rolling and annealing are performed under optimal conditions on steel sheets of the above-mentioned composition, and the ratio of the number of carbides in ferrite grain boundaries to the number of carbides in ferrite grains is more than 1, the ferrite grain size is 5 μm or more, and Vickers hardness By setting the length to 100 HV or more and 180 HV or less, the cold forgeability in cold forging combining drawing and thickness increase forming is improved, and further, the machinability and wear resistance characteristics of the member after quenching and tempering are improved. Is a novel finding found by the present inventors.

The steel sheet of the present invention is substantially composed of ferrite and carbide, and has a structure in which the ratio of the number of carbides at ferrite grain boundaries to the number of carbides in ferrite grains exceeds one. In addition to cementite (Fe ₃ C) which is a compound of iron and carbon, carbides are compounds obtained by replacing Fe atoms in cementite with Mn, Cr, etc., alloy carbides (M ₂₃ C ₆ , M ₆ C, MC etc. And M is Fe and other additive metal elements).

  Next, the reason for the above definition will be described.

  When cutting a material using a soft medium carbon steel plate as a material and using a tool that is harder than the material, if the cuttings cut out of the material are finely divided, it becomes difficult for the cuttings to be entangled in the tool, so Cutting performance is improved. When the cutting process is viewed from a micro viewpoint, the blade of the tool bites into the material, and can be scooped up and separated in the process of separating cutting waste from the material, and in order to obtain fine cutting waste, from the material It is necessary to start the process of separating cutting chips early.

  Conventionally, it is known that the formation of MnS or the like can promote the effect of separating cutting chips from the material, but the separating effect is limited only when the cutting edge of the tool reaches the place where MnS exists. There is a problem.

  Then, the present inventors advanced development of the technique of separating the above-mentioned cutting waste from a raw material early for the solution of the said subject. Specifically, paying attention to carbides that are more finely and uniformly dispersed than MnS, in processing where the blade of the tool is subjected to high strain that cuts into the material, the interface between carbide and ferrite peels off prematurely to cause voids. It has been found that, if generated, the machinability is improved.

  That is, a structure is formed in which the ratio of the number of carbide grains in ferrite grain boundaries to the number of carbide grains in ferrite grains exceeds 1 so that a large amount of carbides exist on ferrite grain boundaries, and By separating S and promoting the separation of carbide and ferrite, it was clarified that cutting chips are finely dispersed at the time of cutting, and further, the force required for cutting is reduced and the machinability is remarkably improved.

  Based on theory and principle, machinability is considered to be strongly affected by the coverage of carbides in ferrite grain boundaries, and high-accuracy measurement is required. However, carbides to ferrite grain boundaries in three-dimensional space In order to measure coverage, serial sectioning SEM observation or three-dimensional EBSP observation is required by repeatedly performing sample cutting and observation by FIB in a scanning electron microscope, which requires an enormous amount of measurement time, The inventors clarified that the accumulation of technical know-how is essential, and concluded that it is not worthy of a general analysis method.

  For this reason, a simpler and more accurate evaluation index was searched. As a result, if the ratio of the number of carbides in ferrite grain boundaries to the number of carbides in ferrite grains is used as an index, it becomes possible to evaluate the machinability, and the carbides of ferrite grain boundaries relative to the number of carbides in ferrite grains The present inventors have found that when the ratio of the number of elements exceeds 1, the machinability is significantly improved.

  The average particle diameter of the carbide is preferably 0.1 μm or more and 2.0 μm or less. If the grain size of the carbide is less than 0.1 μm, the hardness of the steel sheet significantly increases and brittle fracture tends to occur and the machinability decreases, so the average grain size of the carbide is 0.1 μm. The above is preferable. More preferably, it is 0.17 μm or more. On the other hand, if the average particle size of the carbide exceeds 2.0 μm, cracks start from the coarse carbide during cold forging and the machinability decreases, so the average particle size of the carbide is preferably 2.0 μm or less . More preferably, it is 1.95 μm or less.

  Subsequently, an observation method and a measurement method of the tissue will be described.

  The observation of carbides is conducted with a scanning electron microscope. Prior to observation, a sample for texture observation is polished by wet polishing with emery paper and diamond abrasive having an average particle size of 1 μm, and the observation surface is finished to a mirror surface. Thereafter, the tissue is etched with a saturated picric acid-alcohol solution.

  The magnification of observation is 3000 times, and eight 30 μm × 40 μm fields in a 1/4 thickness plate are randomly taken. In the obtained tissue image, the area of each carbide contained in the image area is measured in detail by image analysis software represented by Mitani Shoji Co., Ltd. (Win ROOF). The equivalent circle diameter (= 2 × √ (area / 3.14)) is determined from the area of each carbide, and the average value is taken as the carbide average particle size.

In order to suppress the influence of measurement errors due to noise, carbides with an area of 0.01 μm ² or less are excluded from the evaluation targets. The carbides are specified as carbides having an area of more than 0.01 μm ² .

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

  In the structure after annealing of the hot-rolled steel sheet, the ferrite particle size is 5.0 μm or more. If the ferrite grain size is less than 5.0 μm, the hardness increases and cracks and cracks are easily generated during cold forging, so the ferrite grain size is set to 5.0 μm or more. By setting the ferrite grain size to 5.0 μm or more, the machinability can be improved. Preferably it is 7.5 micrometers or more.

  The upper limit of the ferrite particle size is not particularly limited, but is preferably 50.0 μm or less. If it exceeds 50.0 μm, the number of carbides on the grain boundary which suppresses the propagation of slip decreases, and the machinability decreases, so the ferrite grain size is preferably 50.0 μm or less. More preferably, it is 38.0 μm or less.

  The ferrite grain size is determined by observing the structure etched with a 3% nitric acid-alcohol solution with an optical microscope or a scanning electron microscope after polishing the observation surface to a mirror surface in the above-mentioned procedure, and using the line segment method for the photographed image. Apply and measure.

  In the steel sheet of the present invention, by setting the Vickers hardness to 100 HV or more and 180 HV or less, the machinability and the wear resistance after quenching and tempering can be improved.

  If the Vickers hardness is less than 100 HV, buckling tends to occur during cold forging, folding and folding of the buckling portion occur, and the wear resistance property is lowered, so the Vickers hardness is 100 HV or more I assume. Preferably, it is 110 HV or more. On the other hand, if the Vickers hardness exceeds 180 HV, the ductility is reduced, internal cracking is likely to occur during cold forging, and the wear resistance is deteriorated. Therefore, the Vickers hardness is set to 180 HV or less. Preferably it is 170 HV or less.

  Subsequently, a method of evaluating the machinability will be described.

  Usually, when evaluating the machinability of steel materials, it is used to reduce the weight loss of drills due to wear and / or cut from steel materials using a drill oxide coated with a hard oxide layer or a cemented carbide such as Ti-N. Examine and evaluate the shape of chips and the like.

  However, with this evaluation method, it is not possible to prevent the material change of the steel material to be cut and the drill bit accompanying the heat generation in the wear process, and it is impossible to clarify the structure factor of the steel material that truly affects the machinability. For this reason, in the case of evaluation of the machinability of the steel plate of the present invention, with the purpose of suppressing the heat generation of the steel plate during cutting as much as possible, a method of evaluating machinability using a microcutter was developed.

  For evaluation, using a Microtech (RCA-134 type) made by Refinetech, the rotating grinding stone is made to approach and cut at a horizontal moving speed of 0.5 mm / sec against the steel plate for which the machinability is evaluated. The amount of current applied to the motor was read.

  It was determined that the larger the current (more than 15 amperes) required for cutting, the more the motor was loaded and the lower the machinability. At the time of the evaluation test, the rotary grindstone was rotated at a speed of 2670 rpm using a power supply of 200 V and 60 Hz as a power supply. The rotary grindstone used was WA120NA9C-403249 manufactured by Nippon Plastics Co., Ltd., and the amount of current required for idling was 5.5 amperes.

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

  The technical idea of the manufacturing method of the present invention is characterized by improving the machinability and the wear resistance after quenching and tempering by consistent management of hot conditions and annealing conditions using a steel plate of the above-mentioned component composition. Do.

  Specific features of the production method of the present invention are as follows.

  A continuously cast slab having a predetermined component composition is hot-rolled as it is or after cooling, as it is, as it is or after cooling, to finish hot-rolling in a temperature range of 650 ° C. or more and 950 ° C. or less. The steel strip after finish rolling is cooled on ROT (Run Out Table), and taken up in a temperature range of 400 ° C. to 600 ° C. to form a hot rolled coil.

  The hot rolled coil is pickled and then subjected to a two-step-step box annealing maintained in two temperature ranges. At this time, in the first stage annealing, the hot-rolled steel sheet is heated to an annealing temperature of 30 ° C / hour. Annealing is performed by heating at a heating rate of 150 ° C./hour or less, holding in a temperature range of 650 ° C. or more and 720 ° C. or less for 3 hours or more and 60 hours or less, and further, in the second stage annealing, a hot rolled steel plate Annealing is performed by heating at a heating rate of 1 ° C./hour to 80 ° C./hour to the annealing temperature, holding for 3 hours to 50 hours or less in the temperature range of 725 ° C. to 790 ° C., and then cooling rate to 650 ° C. By cooling at 1 ° C./hour or more and 100 ° C./hour or less and then cooling to room temperature, a low carbon steel plate excellent in the machinability and the wear resistance characteristics after quenching and tempering can be produced.

  Below, the manufacturing method of this invention is demonstrated concretely.

Hot rolling A continuous cast slab having a predetermined composition is used as it is or after being cooled and heated to be subjected to hot rolling, and finish hot rolling is completed in a temperature range of 650 ° C. or more and 950 ° C. or less, and thermal The rolled steel sheet is wound up in a temperature range of 400 ° C. or more and 600 ° C. or less.

  The heating temperature of the slab is preferably 1300 ° C. or less. The heating time during which the temperature of the surface layer of the slab is kept at 1000 ° C. or more is preferably 7 hours or less. When the heating temperature exceeds 1300 ° C. or the heating time exceeds 7 hours, decarburization from the surface layer of the slab becomes remarkable, and austenite grains in the surface layer abnormally grow at the time of heating before quenching, and the wear resistance is deteriorated. The heating temperature is preferably 1300 ° C. or less, and the heating time is preferably 7 hours or less. More preferably, the heating temperature is 1280 ° C. or less, and the heating time is 6 hours or less.

  Finishing hot rolling is completed in a temperature range of 650 ° C. or more and 950 ° C. or less. If the finish hot rolling temperature is less than 650 ° C., the deformation resistance of the steel material is increased, the rolling load is significantly increased, and the roll wear amount is increased to lower the productivity, so the finish hot rolling temperature is 650 ° C or more. Preferably it is 680 ° C or more.

  On the other hand, if the finish hot rolling temperature exceeds 950 ° C., thick scales are generated during passing through the ROT, and wrinkles due to the scales are generated on the steel sheet surface, and impact load during cold forging and / or after quenching and tempering At the time of addition, since a crack is generated starting from the weir and the wear resistance property is lowered, the finish hot rolling temperature is made 950 ° C. or less. Preferably it is 920 degrees C or less.

  The cooling rate on the ROT after finish hot rolling is preferably 10 ° C./second to 100 ° C./second. If the cooling rate is less than 10 ° C./sec, thick scales are generated during cooling, and it is not possible to prevent generation of wrinkles due to the scales on the surface of the steel sheet, and the wear resistance is degraded. The speed is preferably 10 ° C./sec or more. More preferably, it is 20 ° C./second or more.

  On the other hand, when the steel sheet is cooled at a cooling rate exceeding 100 ° C./sec from the surface layer to the inside of the steel sheet, the outermost layer portion is excessively cooled and a low temperature transformation structure such as bainite or martensite occurs. When taking out the hot-rolled coil cooled to 100 ° C. to room temperature after winding, microcracks occur in the low temperature transformation structure, and it is difficult to remove the microcracks in the subsequent pickling step and cold rolling step. When an impact load is applied during forging and / or after quenching and tempering, the crack develops from the crack as a starting point, and the wear resistance property decreases, so the cooling rate is preferably 100 ° C./sec or less. More preferably, it is 90 ° C./second or less.

  The above cooling rate is the time when the steel strip after finishing hot rolling receives water cooling in the water injection section after passing through the non-water injection section, and is cooled on the ROT to the target temperature of the skimming, It refers to the cooling capacity received from the cooling facility, 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 set to 400 ° C. or more and 600 ° C. or less. If the coiling temperature is less than 400 ° C., the austenite which has not been transformed before coiling transforms into hard martensite, and cracks are generated in the surface layer at the time of dispensation of the hot-rolled coil, so that the wear resistance is deteriorated The winding temperature is 400 ° C. or more. Preferably it is 430 degreeC or more.

  On the other hand, if the coiling temperature exceeds 600 ° C., pearlite with a large lamellar spacing is formed, thick needle-like carbides with high thermal stability are formed, and needle-like carbides even after two-step box annealing The coiling temperature is set to 600 ° C. or less, since a crack is generated during cold forging from this needle-like carbide as a starting point. Preferably it is 570 degrees C or less.

  After pickling, the hot rolled coil manufactured under the conditions described above is subjected to two-step-step box annealing maintained at two temperature ranges. Two-step box annealing controls the stability of the carbides and promotes the formation of carbides at ferrite grain boundaries.

  First, the first stage annealing is carried out in a temperature range of Ac 1 point or less to coarsen the carbides and concentrate the alloy elements to enhance the thermal stability of the carbides. Thereafter, the temperature is raised to a temperature range of Ac 1 point or more, austenite is generated in the structure, carbides in fine ferrite grains are dissolved in austenite, and coarse carbides are left in austenite. The subsequent slow cooling transforms the austenite into ferrite and increases the carbon concentration in the austenite.

  By gradually cooling, carbon atoms are adsorbed to the carbide remaining in austenite, and carbide and austenite cover ferrite grain boundaries, and finally, a structure in which a large amount of carbides is formed in ferrite grain boundaries. It is possible to form Therefore, it is obvious that the structural form defined by the steel sheet of the present invention can not be obtained by simple annealing alone.

  Specific process conditions will be described below.

First stage annealing Heating rate to annealing temperature: 30 ° C / hour or more and 150 ° C / hour or less Annealing temperature: 650 ° C or more and 720 ° C or less At the time of holding: 3 hours or more and 60 hours or less

  In the first stage annealing, the heating rate to the annealing temperature (the first stage heating rate) is set to 30 ° C./hour or more and 150 ° C./hour or less. If the heating speed at the first stage is less than 30 ° C./hour, it takes time to raise the temperature, and the productivity is reduced. Therefore, the heating speed at the first stage is 30 ° C./hour or more. Preferably 40 ° C./hour.

  On the other hand, if the heating speed at the first stage exceeds 150 ° C./hour, the temperature difference between the outer periphery of the coil and the inside increases, causing scorching or seizure due to the difference in thermal expansion, resulting in irregularities on the steel sheet surface Generates. During cold forging, cracks are generated from this unevenness and the machinability is reduced, and the wear resistance after quenching and tempering is reduced, so the heating speed at the first stage is 150 ° C./hour or less. . Preferably it is 120 degrees C / hour or less.

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

  On the other hand, if the annealing temperature of the first stage exceeds 720 ° C., austenite is formed before the stability of the carbides increases, and the above-mentioned structural change can not be controlled, so the annealing temperature of the first stage is 720 ° C. or less Do. Preferably it is 700 degrees C or less.

  The holding time (first holding time) in the first stage annealing is set to 3 hours or more and 60 hours or less. If the first stage holding time is less than 3 hours, stabilization of the carbides is not sufficient and it becomes difficult to leave carbides during the second stage annealing, so the first stage holding time is 3 hours or more I assume. Preferably, it is 10 hours or more.

  On the other hand, if the first stage holding time exceeds 60 hours, further improvement in the stability of the carbide can not be expected, and furthermore, the productivity decreases, so the first stage holding time is set to 60 hours or less. Preferably it is 50 hours or less.

Second stage annealing Heating rate up to annealing temperature: 1 ° C / hour to 80 ° C / hour Annealing temperature: 725 ° C to 790 ° C Holding time: 3 hours to 50 hours

  In the second stage annealing after the holding in the first stage annealing is completed, the heating rate to the annealing temperature (the 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 grown from ferrite grain boundaries. At this time, by controlling the heating rate at the second stage small, nucleation of austenite can be suppressed, and the grain boundary coverage of carbide can be increased in the structure obtained after slow cooling.

  For this reason, although it is preferable that the second-stage heating rate be slower, if it is less than 1 ° C./hour, it takes time to raise the temperature and productivity decreases, so the second-stage heating rate is 1 ° C. It will be more than time. Preferably it is 10 ° C./hour or more.

  On the other hand, if the heating speed at the second stage exceeds 80 ° C./hour, the temperature difference between the outer periphery of the coil and the inside increases, generating scratching and seizure due to a large difference in thermal expansion due to transformation, and steel plate Irregularities are generated on the surface. During cold forging, a crack is generated starting from this unevenness, the machinability is reduced, and the wear resistance after quenching and tempering is reduced, so the second stage heating rate is set to 80 ° C./hour or less. . Preferably it is 70 degrees C / hour or less.

  The annealing temperature (second annealing temperature) in the second annealing is set to 725 ° C. or more and 790 ° C. or less. If the annealing temperature at the second stage is less than 725 ° C., the amount of austenite formed is small, and the number ratio of carbides on ferrite grain boundaries decreases, so the annealing temperature at the second stage is 725 ° C. or more. Preferably it is 735 ° C or more.

  On the other hand, if the annealing temperature at the second stage exceeds 790 ° C., it becomes difficult to retain carbides in austenite, and it becomes difficult to control to the above-mentioned structure change, so the annealing temperature at the second stage is 790 ° C. or less I assume. Preferably it is 770 degrees C or less.

  In the second stage annealing, the holding time at the annealing temperature (second stage holding time) is set to 3 hours or more and 50 hours or less. If the second stage holding time is less than 1 hour, the amount of austenite formed is small, and the dissolution of carbides in ferrite grains is not sufficient, making it difficult to increase the number ratio of carbides on grain boundaries Therefore, the second stage holding time is 1 hour or more. Preferably, it is 5 hours or more.

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

Cooling stop temperature after annealing: 650 ° C
Cooling rate: 1 ° C./hour or more and 100 ° C./hour or less After holding in the second stage annealing is completed, cooling is performed at a cooling rate of 1 ° C./hour or more and 100 ° C./hour or less to 650 ° C.

  In order to transform austenite formed in the second stage annealing into ferrite by gradual cooling and to make carbon adsorbed to carbide remaining in austenite, it is preferable that the cooling rate be slower, but the cooling rate is less than 1 ° C / hour Since the time required for cooling will increase and productivity will fall, the cooling rate should be 1 ° C./hour or more. Preferably it is 10 ° C./hour or more.

  On the other hand, when the cooling rate exceeds 100 ° C./hour, austenite transforms to pearlite, the hardness of the steel plate increases, the machinability decreases, and the wear resistance after quenching and tempering decreases. The cooling rate is 100 ° C./hour or less. Preferably it is 90 degrees C / hr or less.

  The atmosphere of box annealing is not limited to a specific atmosphere. For example, an atmosphere of 95% or more of nitrogen, an atmosphere of 95% or more of hydrogen, or an air atmosphere may be used.

  As described above, according to the manufacturing method of the present invention, the structure of ferrite and carbide is substantially, the ratio of the number of carbides in ferrite grain boundaries to the number of carbides in ferrite grains exceeds 1, and the ferrite grain size is 5.0 μm The above is excellent in cold forgeability and machinability in cold forging in which Vickers hardness is 100 HV or more and 180 HV or less, in combination with drawing and thickness increase forming, and also for the wear resistance characteristics of members after quenching and tempering. An excellent low carbon steel sheet can be manufactured.

  EXAMPLES Next, the effects of the present invention will be described by way of examples. The level of the examples is an example of the execution conditions adopted to confirm the practicability of the present invention and the effects, and the present invention is directed to one example of the conditions. It is not limited to The present invention does not deviate from the scope of the present invention, and various conditions can be adopted as long as the object of the present invention is achieved.

  A continuous cast slab (steel ingot) having the composition shown in Table 1 is subjected to hot rolling after being heated at 1250 ° C. for 1.5 hours, and after finishing hot rolling at 870 ° C., 40 ° C. on ROT It cooled to 520 degreeC with the cooling rate of 1 / second, wound up at 520 degreeC, and manufactured the 4.4-mm-thick hot-rolled coil of board thickness.

  After pickling, the hot-rolled coil is placed in a box annealing furnace, the atmosphere is controlled to 95% hydrogen-5% nitrogen, and then heated from room temperature to 700 ° C. at a heating rate of 80 ° C./hour Hold at 700 ° C. for 24 hours to homogenize the temperature distribution in the coil, then heat to 745 ° C. at a heating rate of 10 ° C./hour and hold at 745 ° C. for 6 hours, then 8 ° C. to 650 ° C. It was cooled at a cooling rate of 1 / hour and then furnace cooled to room temperature to make samples for characterization. The structure of the sample was observed by the above-mentioned method, and the machinability was evaluated by the above-mentioned method.

  Quenching and tempering of the sample were carried out under the following conditions. In a furnace controlled to an atmosphere of 95% nitrogen, a treatment of holding at 880 ° C. for 60 minutes was applied and quenched in oil at 60 ° C. The quenched sample was tempered for 45 minutes at 180 ° C. and air-cooled for tempering to prepare a quenched and tempered sample.

  The abrasion resistance of the sample after heat treatment was evaluated by a block on ring test. The SUJ 2 was processed into a ring test piece of 35 35 mm × width 8.8 mm, and further subjected to the above-mentioned heat treatment to produce a ring test piece. The samples for evaluation are cut into thickness x width 15.8 mm x length 10.2 mm, and when two sheets are stacked in the thickness direction, the top and bottom of the thickness should be 6.3 mm so that the total thickness is 6.3 mm. It cut and the two boards were screwed together and the block test piece was produced.

  Set the sample on the tester so that the 6.3 mm × 15.8 mm surface of the block test piece contacts the ring surface, and the sliding speed is 0.4 m / s, oil temperature 140 ° C, load 2000 N, sliding distance 10000 m The weight change before and after the test of the block test piece is measured under the conditions described above, and if the weight loss is less than 5 mg, it is evaluated as excellent in wear resistance, and if the amount of reduction exceeds 5 mg, the wear resistance It was evaluated as inferior.

  Table 2 shows the measurement results and evaluation results of carbide diameter, ferrite particle size, Vickers hardness, ratio of the number of carbides in ferrite grain boundaries to the number of carbides in ferrite grains, motor current at the time of cutting, and wear amount. Show.

  As shown in Table 2, A-1, B-1, C-1, D-1, E-1, F-1, G-1, H-1, I-1, J-1, of the invention steels In each of K-1, L-1, and M-1, the ratio of the number of carbides at ferrite grain boundaries to the number of carbides in ferrite grains is more than 1, and the ferrite grain size is 5.0 μm or more, It is understood that the Vickers hardness is 100 HV or more and 180 HV or less, and the machinability and the wear resistance characteristics after quenching and tempering are excellent.

  On the other hand, Comparative Steel N-1 has a low amount of C and a hardness after annealing of less than 100 HV, so the cold forgeability is lowered and the strength is not increased after hardening and tempering, and the wear resistance property is It has fallen. Comparative steels O-1, Q-1, S-1 and Y-1 contain excessive amounts of S, P, Al and N, and the amount of segregation to the γ / α interface increases at the second stage of annealing. Therefore, the formation of carbides at ferrite grain boundaries is suppressed.

  The comparative steel P-1 contains an excessive amount of Si and the wear resistance is degraded. Comparative Steel R-1 has a small amount of Mn, and it was difficult to suppress the decrease in hardness after quenching and tempering, so the wear resistance properties after quenching and tempering were degraded. Comparative Steel T-1 has a small amount of S, so the amount of S concentrated at the interface between ferrite and carbide is insufficient, and the machinability and the wear resistance are degraded.

  Comparative Steel U-1 contains Mn in excess, so the Vickers hardness exceeds 180 HV and the machinability is reduced. Since the comparative steel V-1 contains C in excess, the grain size of ferrite after annealing becomes finer and the machinability is lowered, and coarse carbides remain even after quenching and tempering, so that the wear resistance characteristics are obtained. Is declining.

  Since comparative steel W-1 contains B in excess, the concentration of S at the ferrite / carbide interface is suppressed, and the machinability is reduced. Since the comparative steels X-1 and Z-1 contain O and Ti in excess, oxides and TiC present in ferrite grains become formation sites of carbides during slow cooling after annealing in the two-phase region, and ferrite grain boundaries The formation of carbides is suppressed, and the machinability is reduced.

  Subsequently, in order to investigate the influence of the production conditions, a slab having a component composition of A, B, C, D, E, F, G, H, I, J, K, L, and M shown in Table 1 is used Under the hot-rolling conditions and annealing conditions shown in Table 3, a hot-rolled steel plate annealed sample of 5.2 mm in thickness is prepared, and the ferrite diameter, ferrite grain size, Vickers hardness, and ferrite relative to the number of carbides in the ferrite grain The ratio of the number of carbides in grain boundaries, the motor current at the time of cutting, and the amount of wear were measured and evaluated.

  Table 4 shows the measurement results and evaluation results of carbide diameter, ferrite particle size, Vickers hardness, ratio of the number of carbides in ferrite grain boundaries to the number of carbides in ferrite grains, motor current at the time of cutting, and wear amount. Show.

  In Comparative Steel E-3, the finish hot rolling temperature was low, the rolling load increased, and the productivity decreased. The comparative steel D-2 has a high finishing hot rolling temperature, and scale wrinkles are formed on the surface of the steel plate, and when subjected to a wear resistance test after quenching and tempering, cracking and peeling occur from the scale wrinkle as a starting point. The wear characteristics are reduced.

  In the invention steel F-2, the cooling rate in the ROT was low, the productivity decreased, and the scale wrinkles were generated. In the invention steel C-4, a minute crack was generated in the outermost layer portion due to excessive cooling of the outermost layer portion of the steel plate at a cooling rate in ROT of 100 ° C./sec. In Comparative Steel K-4, the coiling temperature was low, and a large number of low temperature transformation structures such as bainite and martensite were formed to cause embrittlement, so cracking frequently occurred at the time of hot-rolled coil delivery and productivity decreased. Furthermore, the wear resistance properties of the samples taken from the fragments are low.

  The comparative steel L-2 has a high coiling temperature, and in the hot-rolled structure, thick pearlite with a lamellar spacing is formed, and the thermal stability of a needle-like coarse carbide is high, and the carbide is a two-step step type Even after box annealing, the machinability is reduced because the steel remains in the steel sheet. In Comparative Steel B-2, the productivity decreased because the heating rate to the annealing temperature was slow in the first stage annealing of the two-step stepped box annealing.

  In Comparative Steel B-3, since the heating rate to the annealing temperature is fast in the first stage annealing, the temperature difference between the inside and the outer periphery of the coil becomes large, and the generation of chipping and seizing due to the difference in thermal expansion occurs When subjected to an evaluation test of wear resistance after quenching and tempering, cracking and peeling occurred from the ridges, and the wear resistance deteriorated.

  The comparative steel A-4 has a low holding temperature (annealing temperature) in the first stage annealing, is insufficient in the coarsening treatment of the carbide at the Ac1 point or less, and the thermal stability of the carbide is insufficient. As a result, the carbide remaining in the second stage annealing decreases, and in the structure after slow cooling, pearlite transformation can not be suppressed, so the machinability decreases.

  Comparative steel F-4 has a high holding temperature (annealing temperature) in the first stage annealing, austenite is formed during annealing, and carbide stability can not be improved, so pearlite is formed after annealing, and Vickers hardness The machinability was reduced by more than 180 HV.

  In Comparative Steel C-3, the holding time in the first stage annealing was short, and the carbide stability could not be improved, and the machinability decreased. The comparative steel L-4 had a long holding time in the first stage of annealing, and in addition to a decrease in productivity, a seizure flaw was generated, and the wear resistance was deteriorated. In Comparative Steel J-2, the productivity decreased because the heating rate to the holding temperature (annealing temperature) was slow in the second stage annealing of the two-step stepped box annealing.

  Comparative steel K-3 has a high heating rate at the second stage of annealing, so the temperature difference between the inner and outer circumferences of the coil becomes large, and cracking and seizure occur due to a large thermal expansion difference due to transformation. And the wear resistance after quenching and tempering was reduced. Comparative steel H-4 has a low holding temperature (annealing temperature) in the second stage annealing, has a small amount of austenite formed, can not increase the number ratio of carbides in ferrite grain boundaries, and reduces the machinability. did.

  Comparative steel G-3 has a high holding temperature (annealing temperature) in the second stage annealing, and since dissolution of carbides is promoted during annealing, it is difficult to form grain boundary carbides after slow cooling, and pearlite And the Vickers hardness exceeded 180 HV, and the machinability decreased.

  In Comparative Steel E-2, since slow cooling was started with a short holding time in the second stage annealing and leaving fine carbides, the formation of carbides at ferrite grain boundaries was suppressed, and the proportion of carbides in the grains was And the machinability decreased. The comparative steel M-2 had a long holding time in the second stage annealing, and promoted the dissolution of the carbide, so the machinability decreased.

  In Comparative Steel I-3, the cooling rate from the second stage annealing to 650 ° C. is slow, the productivity is lowered, and coarse carbides are formed in the structure after slow cooling, and the coarse remaining after quenching and tempering Cracking occurred from the carbide, and the wear resistance was degraded. In Comparative Steel G-2, the cooling rate from the second stage annealing to 650 ° C. was fast, pearlite transformation occurred during cooling, the hardness increased, and the machinability decreased.

  Next, in order to investigate the influence of other elements, continuously cast slabs (steel ingots) having the component compositions shown in Table 5 and Table 6 (continuation of Table 5) are heated at 1250 ° C. for 1.5 hours, After completion of hot rolling at 870 ° C, it was cooled to 520 ° C at a cooling rate of 40 ° C / s on ROT and wound up at 520 ° C to produce a hot rolled coil with a plate thickness of 4.4 mm. .

  After pickling, the hot-rolled coil is placed in a box annealing furnace, the atmosphere is controlled to 95% hydrogen-5% nitrogen, and then heated from room temperature to 700 ° C. at a heating rate of 80 ° C./hour Hold at 700 ° C. for 24 hours to homogenize the temperature distribution in the coil, then heat to 745 ° C. at a heating rate of 10 ° C./hour and hold at 745 ° C. for 6 hours, and then to 650 ° C. 8 The sample was cooled at a cooling rate of ° C./hour and then furnace cooled to room temperature to prepare a sample for characterization. The structure of the sample was observed by the above-described observation method, and the machinability was evaluated by the above-described measurement method.

  Table 7 shows the measurement results and evaluation results of carbide diameter, ferrite particle size, Vickers hardness, ratio of the number of carbides in ferrite grain boundaries to the number of carbides in ferrite grains, motor current at the time of cutting, and wear amount. Show.

  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, AQ-1, AR-1, AS-1, and AT-1 are all carbides in ferrite grains The ratio of the number of carbides in ferrite grain boundaries to the number of ferrite grains is more than 1, the ferrite grain size is 5.0 μm or more, the Vickers hardness is 100 HV to 180 HV, and for the machinability and wear resistance characteristics after quenching and tempering I understand that it is excellent.

  On the other hand, Comparative Steel AU-1 has a low C content and a hardness of less than 100 HV after annealing, so it has low cold forgeability and is not strengthened after quenching and tempering, and has low wear resistance. The comparative steel BF-1 had a small amount of Mn, and it was difficult to suppress the decrease in hardness at the time of quenching and tempering, so the machinability and wear resistance characteristics after quenching and tempering were degraded. The comparative steel BG-1 had a small amount of S, so the amount of S concentrated at the interface between the ferrite and the carbide was insufficient, and the machinability and the wear resistance were degraded.

  Comparative steels BL-1, BJ-1, BO-1, BH-1, AW-1, BS-1, BC-1, AY-1, and BK-1 are S, P, Al, La, respectively. , As, Cu, Ni, Sb, and Ce are excessively contained, and the amount of segregation to the γ / α interface becomes large in the second stage annealing, so that the formation of carbides at ferrite grain boundaries is suppressed.

  The comparative steel BP-1 contained excessive Si and reduced the wear resistance. Since the comparative steels AZ-1, BQ-1, BE-1, BN-1, AX-1, and BR-1 contain Mo, Nb, Cr, Ta, W, and V in excess, respectively, The machinability decreased. The comparative steel BM-1 contained Mn in excess, so the wear resistance after quenching and tempering was significantly reduced.

  The comparative steels BD-1, BB-1, AV-1 and BA-1 contain excessive amounts of Zr, Ca, Mg and Y, respectively, and coarse oxides or nonmetallic inclusions are formed in the steel. As a result, at the time of the evaluation test of the wear resistance, cracking and peeling occurred from the coarse oxide or the coarse nonmetallic inclusion as a starting point, and the wear resistance deteriorated.

  The comparative steel BI-1 contained an excessive amount of Sn and the steel became embrittled, so that the wear resistance was deteriorated. Since the comparative steel BT-1 contains C in excess, the grain size of ferrite after annealing becomes finer and the machinability is lowered, and coarse carbides remain even after quenching and tempering, resulting in wear resistance characteristics. Decreased.

  Subsequently, in order to examine the influence of the production conditions, AA, AB, AC, AD, AE, AF, AG, AH, AI, AJ, AK, AL, AM, AN, AO, AP, AQ, AR shown in Table 5 The hot-rolled steel plate annealing sample of 5.2 mm of board thickness was produced on the hot rolling conditions shown in Table 8, and annealing conditions using the slab which has the component composition of AS, and AT.

  Table 9 shows the measurement results and evaluation results of the carbide diameter, ferrite particle size, Vickers hardness, the ratio of the number of carbide grains in ferrite grain boundaries to the number of carbide grains in ferrite grains, the motor current at the time of cutting, and the amount of wear. Show.

  Comparative steel AC-3 had a low finish hot rolling temperature, increased rolling load and decreased productivity. The comparative steel AC-2 has a high finish hot rolling temperature, scale wrinkles are formed on the surface of the steel plate, and when subjected to a wear resistance test after quenching and tempering, cracking and peeling occur from the scale wrinkle as a starting point, and wear resistance The characteristics have fallen.

  In the invention steel AO-2, the cooling rate in ROT was slow, the productivity decreased, and scale wrinkles occurred. In the invention steel AJ-4, the cooling rate in the ROT was 100 ° C./sec, and the outermost layer portion of the steel plate was excessively cooled, and fine cracks were formed in the outermost layer portion.

  In the comparative steel AH-3, the coiling temperature was low, a large amount of low temperature transformation structure such as bainite or martensite was formed to cause embrittlement, cracking frequently occurred at the time of hot-rolled coil delivery, and productivity decreased. Furthermore, the wear resistance properties of the samples taken from the fragments were low.

  Comparative steel AS-2 has a high coiling temperature, and in hot-rolled structure, thick pearlite with a lamellar spacing is formed, and the thermal stability of acicular coarse carbides is high, and after two-step box annealing Also in the above, the above-mentioned carbides remained in the steel plate, so the machinability decreased. In Comparative Steel AN-3, the heating rate to the holding temperature (annealing temperature) in the first stage annealing of the two-step-step box annealing was slow, so the productivity was reduced.

  Comparative steel AR-3 has a high heating rate at the first stage of annealing, so the temperature difference between the inner and outer circumferences of the coil becomes large, causing cracking and seizing due to the thermal expansion difference, and after quenching and tempering When subjected to an evaluation test of wear resistance, cracking and peeling occurred from the ridge and the wear resistance was reduced.

  The comparative steel AG-2 has a low holding temperature (annealing temperature) in the first stage annealing, is insufficient in the coarsening treatment of the carbide at the Ac1 point or less, and the thermal stability of the carbide is insufficient. As a result, the carbide remaining at the second stage of annealing decreases, and in the structure after slow cooling, pearlite transformation can not be suppressed, and the machinability is reduced.

  Comparative steel AQ-4 has a high holding temperature (annealing temperature) in the first stage annealing, austenite is generated during annealing, and carbide stability can not be improved, so pearlite is generated after annealing, and Vickers hardness The machinability was reduced beyond 180 HV. In the comparative steel AF-4, the holding time in the first stage annealing was short, the carbide stability could not be improved, and the machinability decreased. The comparative steel AM-2 had a long holding time at the first stage of annealing, and in addition to a decrease in productivity, a seizure flaw was generated and the wear resistance was deteriorated.

  The comparative steel AG-4 had a low heating rate to the holding temperature (annealing temperature) in the second stage annealing of the two-step-step-type annealing, and thus the productivity decreased. Comparative steel AF-2 has a high heating rate at the second stage of annealing, so the temperature difference between the inner and outer circumferences of the coil becomes large, and the generation of cracking and seizing due to a large difference in thermal expansion due to transformation And the wear resistance after quenching and tempering was reduced.

  In Comparative Steel AB-2, the holding temperature (annealing temperature) in the second stage annealing was low, the austenite formation amount was small, the number ratio of carbides in the grain boundaries could not be increased, and the machinability was reduced. Comparative steel AK-4 has a high holding temperature (annealing temperature) in the second stage annealing, and dissolution of carbides is promoted during annealing, so it becomes difficult to form grain boundary carbides after slow cooling, and pearlite It formed, Vickers hardness exceeded 180HV, and machinability fell.

  In Comparative Steel AT-2, since slow cooling was started with a short holding time in the second stage annealing and leaving fine carbides, the formation of carbides at ferrite grain boundaries was suppressed, and the proportion of carbides in the grains was And the machinability decreased. The comparative steel AP-2 had a long holding time at the second stage of annealing and promoted the dissolution of carbides, so the machinability decreased.

  The comparative steel AI-3 has a slow cooling rate from the second stage annealing to 650 ° C., and the productivity is lowered, and coarse carbides are formed in the structure after slow cooling, and the coarse remaining after quenching and tempering Cracking occurred from the carbide, and the wear resistance was degraded. The comparative steel AT-4 had a high cooling rate from the second stage annealing to 650 ° C., pearlite transformation occurred during cooling, the hardness increased, and the machinability decreased.

  Here, FIG. 1 shows the relationship between the ratio of the number of carbides in ferrite grain boundaries to the number of carbides in ferrite grains and the machinability. The machinability was evaluated by the amount of motor current (A) at the time of evaluation of machinability. It is understood that when the number ratio (= number of grain boundary carbides / number of intragranular carbides) exceeds 1, the amount of motor current falls below 15 amperes, the energy required for cutting the steel plate decreases, and the machinability improves. .

  FIG. 2 shows another relationship between the ratio of the number of carbides in ferrite grain boundaries to the number of carbides in ferrite grains and the machinability. The relationship shown in FIG. 2 is a relationship in steel plate samples having various additive elements. The machinability was evaluated by the amount of motor current (A) at the time of evaluation of machinability. From FIG. 2, as in FIG. 1, when the number ratio (= number of grain boundary carbides / number of intragranular carbides) exceeds 1, the amount of motor current falls below 15 amperes, and the energy required to cut the steel plate decreases. It is understood that the machinability is improved.

  The relationship between the S content in the steel and the wear resistance is shown in FIG. The amount of wear is the weight loss of the block specimen after the block on ring test. It is understood from FIG. 3 that when the S content exceeds 0.01%, a hard tribofilm is formed on the surface of the lubricating oil / steel sheet, and the abrasion resistance is significantly improved.

  FIG. 4 shows another relationship between the S content in the steel and the wear resistance. The relationship shown in FIG. 4 is a relationship in steel plate samples having various additive elements. The amount of wear is the weight loss of the block specimen after the block on ring test. It is understood from FIG. 4 that, as in the case of FIG. 3, when the S content exceeds 0.01%, a hard tribofilm is formed on the surface of the lubricating oil / steel material, and the abrasion resistance is remarkably improved.

  As described above, according to the present invention, it is possible to provide a low carbon steel sheet which is excellent in machinability and wear resistance after quenching and tempering, and a method for producing the same. The low carbon steel plate of the present invention is suitable as a material for producing, for example, a drive system component such as an automobile gear or a clutch by cold forging and cutting, so the present invention has high industrial applicability. It is a thing.

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%,
P: 0.0001 to 0.020%,
S: 0.010 to 0.10%,
Al: 0.001 to 0.10%,
N: 0.0001 to 0.020%,
O: 0.0001 to 0.020%,
Ti: 0.0001 to 0.010%,
B: 0 to 0.0005%
Containing the balance Fe and unavoidable impurities,
For carbides with an area of more than 0.01 μm ^2, the ratio of the number of carbide grain boundaries in the ferrite grain to the number of carbide grains in the ferrite grains exceeds 1,
Ferrite grain size is 5.0 μm or more,
A low carbon steel sheet excellent in machinability and wear resistance characteristics after quenching and tempering characterized in that Vickers hardness is 100 HV or more and 180 HV or less. The above-mentioned component composition is further in mass%,
Cr: 0.001 to 0.50%,
Mo: 0.001 to 0.50%,
Nb: 0.001 to 0.10%,
V: 0.001 to 0.10%,
Cu: 0.001 to 0.10%,
W: 0.001 to 0.10%,
Ta: 0.001 to 0.10%,
Ni: 0.001 to 0.10%,
Sn: 0.001 to 0.050%,
Sb: 0.001 to 0.050%,
As: 0.001 to 0.050%,
Mg: 0.0001 to 0.050%,
Ca: 0.001 to 0.050%,
Y: 0.001 to 0.050%,
Zr: 0.001 to 0.050%,
La: 0.001 to 0.050%,
The low carbon steel sheet having excellent machinability and wear resistance properties after quenching and tempering according to claim 1, characterized in that it contains one or more of Ce: 0.001 to 0.050%. It is a manufacturing method which manufactures the low carbon steel plate which is excellent in the machinability according to claim 1 or 2 and the wear resistance characteristic after hardening and tempering,
A steel piece having the component composition according to claim 1 or 2 is directly or temporarily cooled and then heated to be subjected to hot rolling, and finish hot rolling is completed in a temperature range of 650 ° C. or more and 950 ° C. or less When subjecting a hot-rolled steel plate wound at 400 ° C. or more and 600 ° C. or less to two-step box annealing in two temperature ranges after pickling,
(I-1) In the first stage annealing, the heat-rolled steel plate is heated to the annealing temperature at a heating rate of 30 ° C./hour or more and 150 ° C./hour or less, and in the temperature range of 650 ° C. or more and 720 ° C. or less for 3 hours or more Annealed to hold for less than 60 hours,
(I-2) In the second stage annealing, the heat-rolled steel plate is heated to the annealing temperature at a heating rate of 1 ° C./hour to 80 ° C./hour, and at a temperature range of 725 ° C. to 790 ° C. for 3 hours or more Apply annealing holding for less than 50 hours, then
(Ii) 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. The manufacturing method of the low carbon steel plate which is excellent in abrasion resistance characteristics.

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