JP2016216809A - Low carbon steel sheet excellent in cold moldability and toughness after heat treatment and manufacturing method therefor - Google Patents

Abstract

PROBLEM TO BE SOLVED: To enhance cold moldability and toughness after a heat treatment in a low carbon steel sheet.SOLUTION: There is provided a low carbon steel sheet excellent in cold moldability and toughness after a heat treatment containing, by mass%, C:0.10 to 0.40%, Si:0.01 to 0.30%, Mn:0.30 to 1.00%, Al:over 0.10 to 1.00%, P:0.0001 to 0.02%, S:0.0001 to 0.01% and the balance Fe with impurities, and having (x) a ratio of the number of carbides of a ferrite grain boundary:B to the number of carbides in a ferrite particle:A, B/A, of over 1, (y) a ferrite particle diameter of 5 μm to 50 μm inclusive and (z) a Vickers hardness of 100 HV to 150 HV inclusive.SELECTED DRAWING: None

Description

  The present invention relates to a low carbon steel sheet excellent in cold formability and toughness after heat treatment and a method for producing the same.

  Automotive parts, blades, and other machine parts are manufactured through processing steps such as punching, bending, and pressing. In the machining process, it is necessary to improve the workability of the carbon steel plate as a raw material in order to improve and stabilize the product quality and reduce the manufacturing cost. In particular, in the case of drive system parts, the parts may be coupled at high speed, and drive system parts, for example, gear teeth may be missing. In order to prevent the loss of the drive system parts manufactured with the carbon steel sheet, the carbon steel sheet as the material for the drive system parts needs toughness after heat treatment.

  Generally, a carbon steel sheet is subjected to cold rolling and spheroidizing annealing, and a carbon steel sheet is used as a soft material having good workability made of ferrite and spheroidized carbide. And until now, several techniques for improving the workability of carbon steel sheets have been proposed.

  For example, in Patent Document 1, C: 0.15 to 0.90 mass%, Si: 0.40 mass% or less, Mn: 0.3 to 1.0 mass%, P: 0.03 mass% or less, Total Al: 0.10 mass% or less, Ti: 0.01-0.05 mass%, B: 0.0005-0.0050 mass%, N: 0.01 mass% or less, Cr: 1.2 mass% A high-carbon steel sheet for precision punching that has a structure in which carbides with an average carbide particle size of 0.4 to 1.0 μm and a carbide spheroidization rate of 80% or more are dispersed in a ferrite matrix and have a notch tensile elongation of 20% or more. And its manufacturing method.

In Patent Document 2, C: 0.3 to 1.3 mass%, Si: 1.0 mass% or less, Mn: 0.2 to 1.5 mass%, P: 0.02 mass% or less, S: 0.02% by mass or less, and carbide is dispersed so that the relationship of C _GB / C _IG ≦ 0.8 is established between the carbide C _GB on the ferrite grain boundary and the number of carbides C _IG in the ferrite crystal grain A medium and high carbon steel sheet excellent in workability and a method for producing the same are disclosed, characterized by having a texture that is made and having a cross-sectional hardness of 160 HV or less.

In Patent Document 3, C: 0.30 to 1.00 mass%, Si: 1.0 mass% or less, Mn: 0.2 to 1.5 mass%, P: 0.02 mass% or less, S: The relationship of C _GB / C _IG ≦ 0.8 holds between the carbide C _GB on the ferrite grain boundary and the number of carbides C _IG in the ferrite crystal grain, including 0.02% by mass or less. A medium and high carbon steel sheet excellent in workability is disclosed, characterized by having a structure in which a carbide occupying 90% or more of the spheroidized carbide having a major axis / minor axis of 2 or less is dispersed in ferrite. Yes.

  In these conventional techniques, it is assumed that the workability improves as the proportion of carbide in the ferrite grains increases.

In Patent Document 4, C: 0.1 to 0.5 mass%, Si: 0.5 mass% or less, Mn: 0.2 to 1.5 mass%, P: 0.03 mass% or less, S: It has a composition consisting of 0.02% by mass or less and a structure mainly composed of ferrite and carbide, and S _gb = {S _on / (S _on + S _in )} × 100 (where S _on : per unit area) of the carbides present, the total area occupied by the carbides present on the grain boundary, S _in: out of the carbides present per unit area, the ferrite grain boundary carbides, which is defined by the total occupied area) of carbide present in the grain A steel sheet excellent in FB workability, mold life, and forming workability after FB processing, characterized in that the amount S _gb is 40% or more is disclosed.

Japanese Patent No. 4465057 Japanese Patent No. 4974285 Japanese Patent No. 5197076 Japanese Patent No. 5194454

In the technique described in Patent Document 1, aiming at coarsening of ferrite grain size and carbide, annealing is performed at a temperature of A _C1 point or higher for softening, but annealing was performed at a temperature of A _C1 point or higher. In this case, rod-like and plate-like carbides precipitate during annealing. Since this carbide is said to reduce workability, even if it can reduce hardness, it adversely affects workability.

  The techniques described in Patent Documents 2 and 3 both improve the spheroidization rate of grain boundary carbides as a cause of deterioration of workability due to the low spheroidization rate of carbides (grain boundary carbides) precipitated at the grain boundaries. Is not a problem. In the technique described in Patent Document 4, the tissue factor is only defined, and the relationship between workability and mechanical properties is not studied.

  In light of the prior art, the present invention aims to improve cold formability and post-heat-treatment toughness in a low-carbon steel sheet, and to provide a low-carbon steel sheet that solves the problem and a method for producing the same. .

  Here, cold formability means the deformability of a steel plate that can be easily plastically deformed into a required shape without defects when the steel plate is plastically deformed into a required shape by cold working or cold forging. The toughness after heat treatment means toughness after heat treatment of the steel sheet.

  The inventors of the present invention have intensively studied a method for solving the above-described problems. As a result, in the steel sheet structure before cold working of the low-carbon steel sheet with the optimized component composition, the dispersion state of the carbide is controlled by optimizing the process conditions from hot rolling to annealing, and the carbide It has been found that excellent cold formability and post-heat-treatment toughness can be secured in a low-carbon steel sheet by precipitating at the ferrite grain boundary, further by setting the ferrite grain size to 5 μm or more and Vickers hardness to 150 or less.

  In addition, the low carbon steel sheet having excellent cold formability and toughness after heat treatment is difficult to manufacture even if the hot rolling conditions and annealing conditions are simply adjusted individually. It has been found that a low-carbon steel sheet having excellent cold formability and post-heat treatment toughness can be produced by optimizing the production conditions in the integrated process leading to annealing.

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

(1) In mass%,
C: 0.10 to 0.40%,
Si: 0.01-0.30%,
Mn: 0.30 to 1.00%
Al: more than 0.10 to 1.00%,
P: 0.0001 to 0.02%,
S: 0.0001 to 0.01%
In the steel sheet containing the balance and Fe and impurities,
(X) Number of carbides in ferrite grains: Number of carbides at ferrite grain boundaries to A: Ratio of B: B / A exceeds 1,
(Y) The ferrite particle size is 5 μm or more and 50 μm or less,
(Z) A low carbon steel sheet excellent in cold formability and toughness after heat treatment, characterized in that the Vickers hardness is from 100 HV to 150 HV.

(2) The steel sheet is further in mass%,
N: 0.0001 to 0.01%
O: 0.0001 to 0.02%
The low-carbon steel sheet excellent in cold formability and post-heat-treatment toughness as described in (1) above.

(3) The steel sheet is further in mass%,
Ti: 0.001 to 0.10%,
Cr: 0.001 to 0.50%,
Mo: 0.001 to 0.50%,
B: 0.0004 to 0.01%
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 0.05%,
Sb: 0.001 to 0.05%,
As: 0.001 to 0.05%,
Mg: 0.0001 to 0.05%,
Ca: 0.001 to 0.05%,
Y: 0.001 to 0.05%,
Zr: 0.001 to 0.05%,
La: 0.001 to 0.05%,
Ce: 0.001 to 0.05%,
The low-carbon steel sheet excellent in cold formability and post-heat-treatment toughness according to the above (1) or (2), characterized by containing one or more of the above.

(4) A production method for producing a low carbon steel sheet having excellent cold formability and toughness after heat treatment according to any one of (1) to (3),
(I) The steel slab having the component composition according to any one of (1) to (3) is directly or once cooled, and then heated and subjected to hot rolling, and a temperature range of 800 ° C. or higher and 900 ° C. or lower. The hot-rolled steel sheet that has been finished and rolled in is scraped at 400 ° C. or higher and 550 ° C. or lower,
(Ii) After rolling out the hot-rolled steel sheet that has been scraped off and pickling, it is subjected to a first-stage box annealing that is held at a temperature range of 650 ° C. to 720 ° C. for 3 hours to 60 hours, and further 725 ° C. Apply a second-stage box annealing that is held for 3 hours or more and 50 hours or less in a temperature range of 790 ° C. or lower, and perform a two-step type box annealing,
(Iii) Cold formability, wherein the hot-rolled steel sheet after box annealing is cooled to 650 ° C. at a cooling rate controlled to 1 ° C./hour or more and 30 ° C./hour or less, and then cooled to room temperature. And manufacturing method of low carbon steel plate with excellent toughness after heat treatment.

  (5) The method for producing a low-carbon steel sheet excellent in cold formability and toughness after heat treatment as described in (4) above, wherein the temperature of the steel slab subjected to hot rolling is 1000 to 1250 ° C.

  According to the present invention, it is possible to provide a low carbon steel sheet excellent in cold formability and post-heat treatment toughness and a method for producing the same.

  The low-carbon steel sheet (hereinafter sometimes referred to as “the present invention steel sheet”) excellent in cold formability and post-heat treatment toughness according to the present invention is mass%, C: 0.10 to 0.40%, Si: 0. 0.01 to 0.30%, Mn: 0.30 to 1.00%, Al: more than 0.10 to 1.00%, P: 0.0001 to 0.02%, S: 0.0001 to 0.00. (X) Number of carbides in ferrite grains: Number of carbides in ferrite grain boundaries: A ratio of B to B: A exceeds B / A in a steel sheet containing 01% and the balance being Fe and impurities ( y) The ferrite grain size is 5 μm or more and 50 μm or less, and (z) the Vickers hardness is 100 HV or more and 150 HV or less.

  The method for producing a low carbon steel sheet excellent in cold formability and toughness after heat treatment according to the present invention (hereinafter sometimes referred to as “the present invention production method”) is a production method for producing a steel sheet of the present invention. ) A steel plate having the same composition as that of the steel plate of the present invention is directly or once cooled and then heated and subjected to hot rolling, and a hot rolled steel plate that has been finish-rolled in a temperature range of 800 ° C. or higher and 900 ° C. or lower is 400 Scattering at 550 ° C. or more and 550 ° C. or less, (ii) paying out the scooped hot-rolled steel sheet, pickling, and 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 Box annealing is performed, and further, a second stage box annealing is performed in which a second stage box annealing is performed in a temperature range of 725 ° C. to 790 ° C. for 3 hours to 50 hours, and (iii) the box annealing described above. Control the later hot-rolled steel sheet to 1 ℃ / hour or more and 30 ℃ / hour or less Was cooled to 650 ° C. at a cooling rate, then characterized by cooling to room temperature.

  First, the reasons for limiting the component composition of the steel sheet of the present invention will be described. Hereinafter,% means mass%.

C: 0.10 to 0.40%
C is an element that forms carbides and is effective in strengthening steel and refining ferrite grains. In order to suppress the occurrence of satin during cold forming and to ensure the surface appearance of the cold-formed product, it is necessary to suppress the coarsening of the ferrite grain size. If it is less than 0.10%, the volume fraction of carbide is insufficient, and coarsening of the carbide cannot be suppressed during box annealing, so C is made 0.10% or more. Preferably it is 0.14% or more. On the other hand, if C exceeds 0.40%, the volume fraction of carbides increases, and cold formability and post-heat treatment toughness decrease, so C is made 0.40% or less. Preferably it is 0.38% or less.

Si: 0.01-0.30%
In addition to functioning as a deoxidizer, Si is an element that affects the form of carbides and contributes to improvement of toughness after heat treatment. In order to reduce the number of carbides in ferrite grains and increase the number of carbides on ferrite grain boundaries, two-step type box annealing (hereinafter sometimes referred to as “two-stage box annealing”) is used. It is necessary to generate an austenite phase during annealing, dissolve the carbide once, and then gradually cool to promote precipitation of the carbide on the ferrite grain boundary.

  Si is preferably as small as possible, but if it is reduced to less than 0.01%, the refining cost increases significantly and the deoxidation effect does not appear, so Si is made 0.01% or more. Preferably it is 0.07% or more. On the other hand, if Si exceeds 0.30%, the hardness increases due to the solid solution strengthening of ferrite, the ductility decreases, the cold formability decreases, and cracking easily occurs, so Si is 0.30. % Or less. Preferably it is 0.28% or less.

Mn: 0.30 to 1.00%
Mn is an element that controls the form of carbide in the two-stage box annealing. If it is less than 0.30%, it becomes difficult to generate carbides at the ferrite grain boundaries in the slow cooling after the two-stage annealing, so Mn is set to 0.30% or more. Preferably it is 0.33% or more. On the other hand, if it exceeds 1.00%, the hardness of the ferrite increases and the cold formability decreases, so Mn is made 1.00% or less. Preferably it is 0.96% or less.

Al: 0.10 to 1.00%
Al is an element that acts as a deoxidizer and stabilizes ferrite. Al is an element that can improve the toughness after heat treatment without impairing the cold formability because of its low solid solution strengthening ability. If it is less than 0.10%, the effect of addition cannot be sufficiently obtained, so Al is made 0.10% or more. Preferably it is 0.30% or more. On the other hand, when it exceeds 1.00%, and A ₃ points increase, the conventional quenching temperature since quenching is difficult, Al is not more than 1.00%. Preferably it is 0.5% or less.

P: 0.0001 to 0.02%
P is an element that segregates at the ferrite grain boundaries and suppresses the formation of carbides at the ferrite grain boundaries. Therefore, the smaller the amount of P, the better. However, if the P content is reduced to less than 0.0001%, the refining cost is greatly increased. Preferably it is 0.0013% or more. On the other hand, if P exceeds 0.02%, the formation of carbides at the ferrite grain boundaries is suppressed, the number of carbides decreases, and the cold formability deteriorates, so P is made 0.02% or less. Preferably it is 0.01% or less.

S: 0.0001 to 0.01%
S is an impurity element that forms non-metallic inclusions such as MnS. Since non-metallic inclusions are the starting point of cracking during cold forming, S is preferably as small as possible, but if it is reduced to less than 0.0001%, the refining cost will increase significantly, so it is made 0.0001% or more. . Preferably it is 0.0012% or more. On the other hand, if it exceeds 0.01%, non-metallic inclusions are generated and the cold formability deteriorates, so S is made 0.01% or less. Preferably it is 0.009% or less.

  The steel sheet of the present invention may contain the following elements in addition to the above elements.

N: 0.0001 to 0.01%
N is an element that embrittles ferrite when present in a large amount. Therefore, N is preferably as small as possible, but if it is reduced to less than 0.0001%, the refining cost will increase significantly, so it is made 0.0001% or more. Preferably it is 0.0006% or more. On the other hand, if it exceeds 0.01%, the ferrite becomes brittle and the cold formability deteriorates, so N is made 0.01% or less. Preferably, it is 0.007% or less.

O: 0.0001 to 0.02%
O, when present in a large amount, is an element formed by a coarse oxide. Therefore, the smaller the O, the better. However, if the content is reduced to less than 0.0001%, the refining cost increases significantly, so the content is made 0.0001% or more. Preferably it is 0.0011% or more. On the other hand, if it exceeds 0.02%, a coarse oxide is generated in the steel and becomes a starting point of cracking during cold forming, so O is made 0.02% or less. Preferably, it is 0.01% or less.

  In the steel sheet of the present invention, in addition to the above elements, one or more of the following elements may be further contained.

Ti: 0.001 to 0.10%
Ti is an element that forms a nitride and contributes to refinement of crystal grains. If less than 0.001%, the effect of addition cannot be sufficiently obtained, so Ti is made 0.001% or more. Preferably it is 0.005% or more. On the other hand, if it exceeds 0.10%, coarse Ti nitrides are produced and cold formability deteriorates, so Ti is made 0.10% or less. Preferably it is 0.07% or less.

Cr: 0.001 to 0.50%
Cr is an element that contributes to improvement of hardenability, and stabilizes the carbide by being concentrated in the carbide, and forms a stable carbide even in the austenite phase. If it is less than 0.001%, the effect of improving hardenability cannot be obtained, so Cr is made 0.001% or more. Preferably it is 0.007% or more. On the other hand, if it exceeds 0.05%, stable carbides are generated in the austenite phase, the dissolution of carbides is delayed during quenching, and the required quenching strength cannot be obtained, so Cr is 0.50% or less. Preferably it is 0.48% or less.

Mo: 0.001 to 0.50%
Mo, like Mn, is an element effective for controlling the morphology of carbides. If it is less than 0.001%, the effect of addition cannot be obtained, so Mo is made 0.001% or more. Preferably it is 0.017% or more. On the other hand, if it exceeds 0.50%, the in-plane anisotropy of the r value is lowered and the cold formability is lowered, so Mo is made 0.50% or less. Preferably it is 0.45% or less.

B: 0.0004 to 0.01%
B is an element that contributes to improving hardenability. If it is less than 0.0004%, the effect of addition cannot be obtained, so B is made 0.0004% or more. Preferably it is 0.0010% or more. On the other hand, if it exceeds 0.01%, a coarse B compound is produced and the cold formability deteriorates, so B is made 0.01% or less. Preferably it is 0.008% or less.

Nb: 0.001 to 0.10%
Nb is an element that is effective for controlling the morphology of carbides, and is an element that contributes to improving toughness by refining the structure. If it is less than 0.001%, the effect of addition cannot be 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%, a large number of fine Nb carbides are formed, the strength is excessively increased, the number of carbides at the ferrite grain boundaries is decreased, and the cold formability is lowered. 10% or less. Preferably it is 0.09% or less.

V: 0.001 to 0.10%
V, like Nb, is an element that is effective in controlling the morphology of carbides, and is an element that contributes to improving toughness by refining the structure. If it is less than 0.001%, the effect of addition cannot be obtained, so V is set to 0.001% or more. Preferably it is 0.004% or more. On the other hand, if it exceeds 0.10%, a lot of fine V carbides are generated, the strength is excessively increased, the number of carbides at the ferrite grain boundaries is reduced, and the cold formability is lowered, so that V is 0. 10% or less. Preferably, it is 0.09% or less.

Cu: 0.001 to 0.10%
Cu is an element that segregates at the ferrite grain boundary, and is an element that contributes to improvement in strength by forming fine precipitates. If it is less than 0.001%, the effect of improving the strength cannot be obtained, so Cu is made 0.001% or more. Preferably it is 0.004% or more. On the other hand, if it exceeds 0.10%, segregation to the ferrite grain boundary causes red heat embrittlement, and the productivity in hot rolling decreases, so it 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 an element effective for controlling the form of carbide. If less than 0.001%, the effect of addition cannot be 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%, a large number of fine W carbides are formed and the strength is excessively increased, and the number of carbides at the ferrite grain boundaries is reduced and the cold formability is lowered. 10% or less. Preferably it is 0.08% or less.

Ta: 0.001 to 0.10%
Ta, as well as Nb, V, and W, is an element effective for controlling the morphology of carbides. If less than 0.001%, the effect of addition cannot be 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%, a large number of fine Ta carbides are produced and the strength is excessively increased, and the number of carbides at the ferrite grain boundaries is reduced and the cold formability is lowered. 10% or less. Preferably it is 0.09% or less.

Ni: 0.001 to 0.10%
Ni is an element effective for improving toughness. If it is less than 0.001%, the effect of addition cannot be 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 of carbides at the ferrite grain boundary decreases and the cold formability deteriorates, so Ni is made 0.10% or less. Preferably it is 0.09% or less.

Sn: 0.001 to 0.05%
Sn is an element inevitably mixed from the steel raw material. Therefore, Sn is preferably as small as possible, but 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 0.05%, the ferrite becomes brittle and the cold formability deteriorates, so Sn is made 0.05% or less. Preferably it is 0.04% or less.

Sb: 0.001 to 0.05%
Similar to Sn, Sb is an element that is inevitably mixed from the steel raw material, segregates at the ferrite grain boundary, and reduces the number of carbides at the ferrite grain boundary. Therefore, the smaller the Sb, the better. However, when the Sb is reduced to less than 0.001%, the refining cost is greatly increased, so the Sb is set to 0.001% or more. Preferably it is 0.002% or more. On the other hand, if it exceeds 0.050%, Sb segregates at the ferrite grain boundaries, the number of carbides at the ferrite grain boundaries decreases, and the cold formability deteriorates, so Sb is made 0.05% or less. Preferably it is 0.04% or less.

As: 0.001 to 0.05%
As is an element that is inevitably mixed in from the steel raw material and segregates at the ferrite grain boundaries, like Sn and Sb. Therefore, the smaller the As, the better. However, if the As is reduced to less than 0.001%, the refining cost increases significantly, so As is made 0.001% or more. Preferably it is 0.002% or more. On the other hand, if it exceeds 0.05%, As is segregated at the ferrite grain boundary, the number of carbides at the ferrite grain boundary is reduced, and the cold formability is lowered, so As is made 0.05% or less. Preferably it is 0.04% or less.

Mg: 0.0001 to 0.05%
Mg is an element that can control the form of sulfide by addition of a small amount. If less than 0.0001%, the effect of addition cannot be obtained, so Mg is made 0.0001% or more. Preferably it is 0.0008% or more. On the other hand, if it exceeds 0.05%, ferrite becomes brittle and cold formability deteriorates, so Mg is made 0.05% or less. Preferably it is 0.04% or less.

Ca: 0.001 to 0.05%
Ca, like Mg, is an element that can control the form of sulfide with a small amount of addition. If it is less than 0.001%, the effect of addition cannot be obtained, so Ca is made 0.001% or more. Preferably it is 0.003% or more. On the other hand, if it exceeds 0.005%, coarse Ca oxide is generated and becomes a starting point of cracking during cold forging, so Ca is made 0.05% or less. Preferably it is 0.04% or less.

Y: 0.001 to 0.05%
Y, like Mg and Ca, is an element that can control the form of sulfide by addition of a trace amount. If it is less than 0.001%, the effect of addition cannot be obtained, so Y is made 0.001% or more. Preferably it is 0.003% or more. On the other hand, if it exceeds 0.05%, coarse Y oxide is generated and becomes a starting point of cracking during cold forming, so Y is set to 0.05% or less. Preferably it is 0.03% or less.

Zr: 0.001 to 0.05%
Zr is an element that can control the form of the sulfide by addition of a small amount, similarly to Mg, Ca, and Y. If less than 0.001%, the addition effect cannot be obtained, so Zr is 0.001% or more. To do. Preferably it is 0.004% or more. On the other hand, if it exceeds 0.05%, coarse Zr oxide is generated and becomes a starting point of cracking during cold forming, so Zr is made 0.05% or less. Preferably it is 0.04% or less.

La: 0.001 to 0.05%
La is an element that can control the form of sulfide by adding a small amount, but is also an element that segregates at the ferrite grain boundary and reduces the number of carbides at the ferrite grain boundary. If it is less than 0.001%, the effect of controlling the form of sulfide cannot be obtained, so La is made 0.001% or more. Preferably it is 0.003% or more. On the other hand, if it exceeds 0.05%, La segregates at the ferrite grain boundary, the number of carbides at the ferrite grain boundary decreases, and cold formability deteriorates, so La is made 0.050%. Preferably it is 0.04% or less.

Ce: 0.001 to 0.05%
Ce, like La, is an element that can control the form of sulfide by addition of a small amount, but is also an element that segregates at the ferrite grain boundary and reduces the number of carbides at the ferrite grain boundary. If it is less than 0.001%, the effect of controlling the form of sulfide cannot be obtained, so Ce is made 0.001% or more. Preferably it is 0.003% or more. On the other hand, if it exceeds 0.05%, Ce is segregated at the ferrite grain boundary, the number of carbides at the ferrite grain boundary is reduced, and the cold formability is lowered, so Ce is made 0.05% or less. Preferably, it is 0.04% or less.

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

  In the steel sheet of the present invention, in addition to the above component composition, (x) number of carbides in ferrite grains: number of carbides at ferrite grain boundaries to A: ratio of B: B / A exceeds 1, (y) ferrite grains The diameter is 5 μm or more and 50 μm or less, and (y) the Vickers hardness is 100 HV or more and 150 HV or less.

  The steel sheet of the present invention can have excellent cold formability and post-heat-treatment toughness by including the above-described characteristic requirements (x), (y), and (z). This is a new finding found by the present inventors. This will be described below.

  The structure of the steel sheet of the present invention is substantially a structure composed of ferrite and carbide. Then, the number of carbides in the ferrite grains: the number of carbides in the ferrite grain boundary with respect to A: the ratio of B: B / A has a structure exceeding 1.

In addition to the cementite (Fe ₃ C) which is a compound of iron and carbon, the carbide is a compound obtained by substituting Fe atoms in the cementite with an alloy element such as Mn or Cr, or an alloy carbide (M ₂₃ C ₆ , M ₆ C, MC, etc. [M: Fe and other metal elements added as alloys]).

  Next, the feature requirements of the above (x), (y), and (z) will be described.

  When the steel plate is formed into a predetermined shape, a shear band is formed in the macro structure of the steel plate, and slip deformation is concentrated near the shear band. Slip deformation is accompanied by dislocation growth, and a region having a high dislocation density is formed in the vicinity of the shear band. As the amount of strain applied to the steel sheet increases, slip deformation is promoted and the dislocation density increases. In order to improve cold formability, it is effective to suppress the formation of shear bands.

  From the viewpoint of the microstructure, the formation of a shear band is understood as a phenomenon in which a slip generated by a single grain overcomes the grain boundary and continuously propagates to adjacent crystal grains. Therefore, in order to suppress the formation of shear bands, it is necessary to prevent the propagation of slip across the grain boundary. Carbide in the steel sheet is a strong particle that prevents slipping, and by allowing the carbide to exist at the ferrite grain boundary, it is possible to prevent the propagation of the slip across the crystal grain boundary and suppress the formation of shear bands. It becomes possible to improve cold formability.

  Based on the theory and principle, it is considered that cold formability is strongly influenced by the carbide coverage of ferrite grain boundaries, and high-precision measurement is required. However, serial sectioning SEM observation or three-dimensional EBSP observation, which repeats the cutting and observation of the sample by FIB in the scanning electron microscope, is indispensable for the measurement of the coverage of the ferrite grain boundary carbide in the three-dimensional space. Therefore, it takes a lot of measurement time and accumulation of technical know-how is indispensable.

  The present inventors did not adopt the above observation method as a general analysis method, but searched for a simpler and more accurate evaluation index. As a result, the number of carbides in the ferrite grains: the number of carbides at the ferrite grain boundary with respect to A: the ratio of B: B / A can be used as an index, and the cold formability can be quantitatively evaluated, and the ratio: B It has been found that when / A exceeds 1, the cold formability is remarkably improved.

  Any of buckling, folding, and folding that occurs during cold forming of a steel sheet is caused by the localization of strain associated with the formation of a shear band. Formation and strain localization are mitigated, and buckling, folding, and folding are suppressed.

  When the spheroidization rate of the carbides at the grain boundaries is 80% or less, strain is concentrated locally on the rod-like or plate-like carbides, and voids and cracks are likely to occur. Therefore, the spheroidization rate of the carbides at the grain boundaries is preferably 80% or more, and more preferably 90% or more.

  When the average particle size of the carbide is less than 0.1 μm, the hardness of the steel sheet is remarkably increased and the cold formability is lowered. Therefore, the average particle size of the carbide is preferably 0.1 μm or more. More preferably, it is 0.2 μm or more. On the other hand, if the average particle diameter of the carbide exceeds 2.0 μm, the carbide becomes the starting point of cracking during cold forming, and therefore the average particle diameter of the carbide is preferably 2.0 μm. More preferably, it is 1.95 μm or less.

  Subsequently, a method for observing and measuring the tissue will be described.

  Carbide is observed with a scanning electron microscope. Prior to observation, a tissue observation sample was wet-polished with emery paper and polished with diamond abrasive grains having an average particle size of 1 μm, and the observation surface was mirror-finished. Etch. The magnification for observation is selected from among magnifications of 3000 times so that the structure of ferrite and carbide can be distinguished. At the selected magnification, 8 images of a 30 μm × 40 μm field of view in a 1/4 layer thickness are taken at random.

About the obtained structure | tissue image, the area of the carbide | carbonized_material contained in an analysis area | region is measured in detail with image analysis software (Win ROOF by Mitani Corporation). The equivalent circle diameter (= 2 × √ (area / 3.14)) is determined from the carbide area, and the average value is defined as the carbide particle diameter. The spheroidization rate of carbide is approximated to an ellipse with an equal area and equal moment of inertia, and the ratio of the maximum length to the maximum length in the perpendicular direction is less than 3 to calculate the spheroidization rate. Rate. In order to suppress an increase in measurement error due to noise, carbides having an area of 0.01 μm ² or less are excluded from evaluation targets.

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

  In the steel sheet structure after annealing, the cold formability can be improved by setting the ferrite grain size to 5 μm or more. If the ferrite particle size is less than 5 μm, the hardness increases and cracks and cracks are likely to occur during cold forming, so the ferrite particle size is set to 5 μm or more. Preferably it is 7 micrometers or more. On the other hand, if the ferrite grain size exceeds 50 μm, the number of carbides on the grain boundaries that suppress the propagation of slip is reduced and the cold formability is lowered, so the ferrite grain size is set to 50 μm or less. Preferably it is 38 micrometers or less.

  The ferrite grain size is the image taken by observing the etched structure with a 3% nitric acid-alcohol solution and then observing the etched structure with an optical microscope or a scanning electron microscope, after polishing the sample observation surface to a mirror surface according to the procedure described above. Can be measured by applying the line segment method.

  The cold formability can be improved by setting the Vickers hardness of the steel plate to 100 HV or more and 150 HV or less. If the Vickers hardness is less than 100 HV, buckling easily occurs during cold forming, so the Vickers hardness is 100 HV or more. Preferably it is 120HV or more. On the other hand, if the Vickers hardness exceeds 150 HV, the ductility is lowered and internal cracks are likely to occur during cold forming, so the Vickers hardness is set to 150 HV or less. Preferably it is 140HV or less.

  Next, the manufacturing method of this invention steel plate is demonstrated.

  The manufacturing method of the present invention is characterized in that the hot rolling technique and the annealing technique are managed consistently to control the structure of the steel sheet.

  The molten steel having the required composition is continuously cast, and the steel slab is directly or once heated after cooling and subjected to hot rolling, and finish rolling is completed in a temperature range of 800 ° C. or higher and 900 ° C. or lower. The hot rolled steel sheet that has been finish-rolled is scraped off in a temperature range of 400 ° C. or higher and 550 ° C. or lower. The hot-rolled steel sheet that has been scraped off is taken out, pickled, then subjected to two-stage box annealing, and after annealing, it is cooled to 650 ° C. at a cooling rate controlled to 1 ° C./hour to 30 ° C./hour, Cool to room temperature.

  In the second-stage annealing, the hot-rolled steel sheet is held in a temperature range of 650 ° C. or more and 720 ° C. or less for 3 hours or more and 60 hours or less in the first-stage box annealing, and in the second-stage annealing, 725 ° C. or more and 790 ° C. Annealing is performed for 3 hours or more and 50 hours or less in the following temperature range.

  A structure composed of fine pearlite and bainite can be obtained by hot rolling the steel slab.

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

  If the slab temperature or slab heating temperature exceeds 1250 ° C, or if the slab heating time exceeds 3 hours, decarburization from the slab surface layer is significant, and the austenite grains on the steel sheet surface layer are abnormal during heating before quenching. And cold formability is reduced. For this reason, the slab temperature or the slab heating temperature is preferably 1250 ° C. or less, and the heating time is preferably 3 hours or less. More preferably, it is 1200 degrees C or less and 2.5 hours.

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

  Finish rolling is completed in a temperature range of 800 ° C. or higher and 900 ° C. or lower. When the finishing temperature is less than 800 ° C., the deformation resistance of the steel sheet increases, the rolling load increases remarkably, the roll wear amount increases, and the productivity decreases, so the finishing temperature is 800 ° C. or more. To do. Preferably it is 830 ° C or more.

  When the finishing temperature exceeds 900 ° C, a thick scale is generated in the run-out table (ROT) through the plate. Due to this scale, wrinkles are generated on the surface of the steel plate. Cracks occur. For this reason, finishing temperature shall be 900 degrees C or less. Preferably it is 870 degrees C or less.

  When the hot-rolled steel sheet after finish rolling is cooled by ROT, the cooling rate is preferably 10 ° C./second or more and 100 ° C./second or less. When the cooling rate is less than 10 ° C./second, a thick scale is generated during the cooling, and wrinkles caused by the scale cannot be suppressed. Therefore, the cooling rate is preferably 10 ° C./second or more. More preferably, it is 15 ° C./second or more.

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

  And at the time of cold forming, a crack generate | occur | produces from a microcrack as a starting point. In order to suppress the formation of a low temperature transformation structure such as bainite or martensite in the outermost layer, the cooling rate is preferably 100 ° C./second or less. More preferably, it is 90 ° C./second or less.

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

  The cutting temperature is 400 ° C. or higher and 550 ° C. or lower. If the milling temperature is less than 400 ° C., the austenite that has not been transformed before the milling is transformed into hard martensite, and when the hot-rolled steel sheet coil is discharged, cracks occur in the surface layer of the hot-rolled steel sheet, Formability is reduced. In order to suppress the transformation, the scraping temperature is set to 400 ° C. or higher. Preferably it is 430 degreeC or more.

  When the scraping temperature exceeds 550 ° C., pearlite with a large lamellar spacing is generated, and thick needle-like carbides with high thermal stability are generated. This acicular carbide remains even after two-stage box annealing. At the time of cold forming of the steel sheet, cracks are generated starting from the needle-like carbides, so the cutting temperature is 550 ° C. or less. Preferably it is 520 degrees C or less.

  After the hot-rolled steel sheet coil is dispensed and pickled, a two-step type box annealing (two-step box annealing) that is held in two temperature ranges is performed. By subjecting the hot-rolled steel sheet to two-stage box annealing, it is possible to control the stability of the carbide, promote the formation of carbide at the ferrite grain boundary, and increase the spheroidization rate of the carbide at the ferrite grain boundary.

The first stage box annealing is performed in a temperature range of A _C1 or lower. By this box annealing, the carbide is coarsened and the alloying elements are concentrated to increase the thermal stability of the carbide. Thereafter, the temperature is raised to a temperature range from A _C1 point to A ₃ point, and austenite is generated in the structure. Thereafter, it is gradually cooled to transform the austenite into ferrite, thereby increasing the carbon concentration of the austenite.

  By slow cooling, carbon atoms are adsorbed on the carbide remaining in the austenite, the carbide and austenite cover the ferrite grain boundary, and finally, a structure in which a large number of spheroidized carbides exist in the ferrite grain boundary can be obtained.

If the amount of residual carbides is small during holding in the temperature range from A _C1 point to A ₃ points, pearlite, rod-like carbide, and plate-like carbide are generated during cooling. When pearlite, rod-like carbide, and plate-like carbide are generated, the cold formability of the steel sheet is significantly lowered. Therefore, increasing the number of residual carbides by holding in the temperature range of A _C1 point or more and A ₃ point or less is important for improving the cold formability.

In the steel sheet structure to form hot-rolled condition described above, a temperature range of A less than point _C1, the thermal stabilization of carbides is promoted, maintained in the temperature range below A _C1 points or more A ₃ points above Thus, the number of residual carbides can be increased.

  The production method of the present invention will be specifically described below.

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

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

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

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

  An annealing temperature (second stage annealing temperature) in the second stage box annealing is set to 725 ° C. or higher and 790 ° C. or lower. If the second stage annealing temperature is less than 725 ° C., the amount of austenite produced is small, and the number of carbides at the ferrite grain boundaries decreases. For this reason, the second stage annealing temperature is set to 725 ° C. or higher. Preferably it is 735 ° C or more.

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

  The annealing time in the second stage box annealing (second stage annealing time) is 3 hours or more and 50 hours or less. If the first stage annealing time is less than 3 hours, the amount of austenite produced is small and the carbides in the ferrite grains are not sufficiently dissolved, making it difficult to increase the number of carbides at the ferrite grain boundaries. For this reason, the second stage annealing time is set to 3 hours or more. Preferably it is 6 hours or more.

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

  After the two-stage box annealing, the steel sheet is cooled to 650 ° C. at a cooling rate controlled to 1 ° C./hour or more and 30 ° C./hour or less. The austenite produced by the second-stage box annealing is gradually cooled to transform it into ferrite, and carbon is adsorbed on the carbide remaining in the austenite. Although it is preferable that the cooling rate is low, if it is less than 1 ° C./hour, the time required for cooling increases and the productivity decreases, so the cooling rate is 1 ° C./hour or more. Preferably, it is 5 ° C./hour.

  On the other hand, when the cooling rate exceeds 30 ° C./hour, austenite is transformed into pearlite, the hardness of the steel sheet is increased, and the cold formability is lowered. Therefore, the cooling rate is set to 30 ° C./hour or less. Preferably it is 26 degrees C / hour or less.

  After the steel plate after the two-stage box annealing is cooled to 650 ° C. at the cooling rate, it is cooled to room temperature. In cooling to room temperature, the cooling rate is not particularly limited.

  The atmosphere in the two-stage box annealing is not particularly limited to a specific atmosphere. For example, any atmosphere of 95% or more nitrogen atmosphere, 95% or more hydrogen atmosphere, or air atmosphere may be used.

  As described above, according to the manufacturing method of the present invention, the structure of ferrite and spheroidized carbide having a grain size of 5 μm or more and 50 μm or less is substantially included, and the number of carbides in ferrite grains: The number of carbides: B ratio: B / A exceeds 1, and furthermore, a low carbon steel sheet having a Vickers hardness of 100 HV to 150 HV and excellent in cold formability and post-heat treatment toughness can be obtained.

  Next, examples of the present invention will be described. The conditions in the examples are examples of conditions adopted for confirming the feasibility and effects of the present invention. It is not limited. The present invention can adopt various conditions as long as the object of the present invention is achieved without departing from the gist of the present invention.

Example 1
In order to investigate the influence of hot rolling conditions on continuous cast slabs (steel slabs) having the composition shown in Table 1 (component composition of invention steel plate) and Table 2 (component composition of comparative steel plate), various hot rolling conditions (table 4) and hot rolled steel sheet coils having a thickness of 3.0 mm were manufactured.

  The hot rolled steel sheet coil is discharged, pickled, charged into a box-type annealing furnace, the annealing atmosphere is controlled to 95% hydrogen-5% nitrogen, heated from room temperature to 705 ° C. and held for 36 hours. After uniformizing the temperature distribution in the steel sheet coil, it was heated to 760 ° C. and held for 10 hours, then cooled to 650 ° C. at a cooling rate of 10 ° C./hour, and then cooled to room temperature in a furnace for characteristic evaluation. A sample was prepared.

  The structure of the sample was observed by the method described above, and the ferrite particle size and the number of carbides were measured. Next, the sample was placed in an atmospheric annealing furnace, and held at 950 ° C. for 20 minutes. After holding, oil cooling at 50 ° C. was performed. Thereafter, tempering was performed so that the hardness was HV400. The toughness after heat treatment was evaluated using the Charpy test. The surface of the sample after the heat treatment was ground to prepare a V-notch Charpy test piece having a thickness of 2 mm, tested at room temperature, and the impact energy was obtained by dividing the obtained absorbed energy by the cross-sectional area.

Table 3 shows ferrite grain size (μm), Vickers hardness (HV), ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grain (number of grain boundary carbides / number of carbides in grain), and heat treatment. Post-toughness (impact value J / cm ² ) is shown. As shown in Table 3, in the invention steel plates (A-1 to Z-1), all have a Vickers hardness of 150 HV or less, and the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains. It can be seen that (the number of grain boundary carbides / the number of carbides in grains) exceeds 1, and the cold formability is excellent.

  On the other hand, the comparative steel plate AA-1 has a large amount of C, the comparative steel plate AB-1 has a large amount of Mn, the comparative steel plate AC-1 has a large amount of Si, and both have a Vickers hardness of 150 HV. Is over. In the comparative steel sheet AE-1, not only the amount of C was small and the Vickers hardness was less than 100 HV, but also not baked at 950 ° C. In other comparative steel sheets, the component composition is outside the range of the component composition of the steel sheet of the present invention, and thus the toughness (impact value) after heat treatment is lowered.

  In order to investigate the influence of annealing conditions, steel slabs having the composition shown in Tables 1 and 2 were heated at 1240 ° C. for 1.8 hours and then subjected to hot rolling, and finish rolling was performed under the hot rolling conditions shown in Table 4. After that, it was cooled on the ROT at a cooling rate of 45 ° C./second, and scraped off at the scraping temperature shown in Table 4 to produce a hot-rolled steel sheet coil having a thickness of 3.0 mm.

A sample for property evaluation with a thickness of 3.0 mm was taken from the hot-rolled steel sheet after annealing, which was subjected to two-step step box annealing under the annealing conditions shown in Table 4 after pickling. , Ferrite grain size (μm), Vickers hardness (HV), ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grain (number of grain boundary carbides / number of carbides in the grain), and toughness after heat treatment ( Impact value J / cm ² ) was measured. The results are also shown in Table 4.

  As shown in Table 4, in the invention steels, the Vickers hardness is 150 HV or less, and the ratio of the number of carbides at the ferrite grain boundary to the number of carbides in the ferrite grains exceeds 1, It can be seen that the moldability is excellent and the toughness after heat treatment is excellent.

  On the other hand, in the comparative steel sheet, since the manufacturing conditions are outside the range of the manufacturing conditions of the manufacturing method of the present invention, the Vickers hardness is increased. In some comparative steel sheets, the number of grain boundary carbides / number of intragranular carbides also decreases.

  As described above, according to the present invention, it is possible to provide a low carbon steel sheet excellent in cold formability and post-heat treatment toughness, and a method for producing the same. Therefore, this invention has a high applicability in steel plate manufacture and utilization industry.

Claims (5)

% By mass
C: 0.10 to 0.40%,
Si: 0.01-0.30%,
Mn: 0.30 to 1.00%
Al: more than 0.10 to 1.00%,
P: 0.0001 to 0.02%,
S: 0.0001 to 0.01%
In the steel sheet containing the balance and Fe and impurities,
(X) Number of carbides in ferrite grains: Number of carbides at ferrite grain boundaries to A: Ratio of B: B / A exceeds 1,
(Y) The ferrite particle size is 5 μm or more and 50 μm or less,
(Z) A low carbon steel sheet excellent in cold formability and toughness after heat treatment, characterized in that the Vickers hardness is from 100 HV to 150 HV. The steel sheet is further in mass%,
N: 0.0001 to 0.01%
O: 0.0001 to 0.02%
The low-carbon steel sheet having excellent cold formability and post-heat-treatment toughness according to claim 1, wherein The steel sheet is further in mass%,
Ti: 0.001 to 0.10%,
Cr: 0.001 to 0.50%,
Mo: 0.001 to 0.50%,
B: 0.0004 to 0.01%
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 0.05%,
Sb: 0.001 to 0.05%,
As: 0.001 to 0.05%,
Mg: 0.0001 to 0.05%,
Ca: 0.001 to 0.05%,
Y: 0.001 to 0.05%,
Zr: 0.001 to 0.05%,
La: 0.001 to 0.05%,
Ce: 0.001 to 0.05%
The low-carbon steel sheet excellent in cold formability and post-heat-treatment toughness according to claim 1 or 2, characterized by containing at least one of the following. A manufacturing method for manufacturing a low carbon steel sheet excellent in cold formability and post-heat treatment toughness according to any one of claims 1 to 3,
(I) The steel slab having the component composition according to any one of claims 1 to 3 is directly or once heated after being cooled and subjected to hot rolling, and finished in a temperature range of 800 ° C or higher and 900 ° C or lower. The rolled hot rolled steel sheet is scraped off at 400 ° C. or higher and 550 ° C. or lower,
(Ii) After rolling out the hot-rolled steel sheet that has been scraped off and pickling, it is subjected to a first-stage box annealing that is held at a temperature range of 650 ° C. to 720 ° C. for 3 hours to 60 hours, and further 725 ° C. Apply a second-stage box annealing that is held for 3 hours or more and 50 hours or less in a temperature range of 790 ° C. or lower, and perform a two-step type box annealing,
(Iii) Cold formability, wherein the hot-rolled steel sheet after box annealing is cooled to 650 ° C. at a cooling rate controlled to 1 ° C./hour or more and 30 ° C./hour or less, and then cooled to room temperature. And manufacturing method of low carbon steel plate with excellent toughness after heat treatment.   The method for producing a low-carbon steel sheet excellent in cold formability and post-heat-treatment toughness according to claim 4, wherein the temperature of the steel slab subjected to hot rolling is 1000 to 1250 ° C.